Failure Analysis of Heat Treated Steel Components

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© 2008 ASM International. All Rights Reserved. Failure Analysis of Heat Treated Steel Component (#05113G)

www.asminternational.org

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 Tecnologıá e Ingenıéra 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 Saõ Paulo Jose´ Ricardo Tarpani, Universidade de Saõ 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 Caracterizaçaõ e Desenvolvimento de Materiais Celso Roberto Ribeiro, Centro de Caracterizaçaõ e Desenvolvimento de Materiais Clever Ricardo Chinaglia, Centro de Caracterizaçaõ e Desenvolvimento de Materiais Jose´ Eduardo Mya, Centro de Caracterizaçaõ 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 Saõ Paulo Xin Yao, Portland State University G.E. Totten, Associaçaõ 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 Saõ Paulo G.E. Totten, Associaçaõ 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 Saõ 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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Failure Analysis of Heat Treated Steel Components L.C.F. Canale, R.A. Mesquita, and G.E. Totten, editors, p 1-42 DOI: 10.1361/faht2008p001

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Component Design Mario Solari, CTI Consultores de Tecnologıá e Ingenierıá 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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Component Design / 3

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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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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pg 5


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

pg 8


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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

pg 11


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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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pg 23


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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pg 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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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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Overview of the Mechanisms of Failure in Heat Treated Steel Components / 45

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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Hazardous Materials Accident Brief, Adopted July 22, 2004 A.W. Loginow, Stress-Corrosion Cracking of Steel in Liquefied Ammonia Service, Bull., Natl. Board Boiler Pressure Vessel Insp., Vol 45 (No. 6), Oct 1988 “Report by the Danish Accident Investigation Board into the Incident to Boeing 757, TF-FIJ at Kastrup Airport, Denmark on June 28, 2001,” M-04001/AIG-09, Danish Accident Review Board J.E. Dorn, Ed., Mechanical Behavior of Materials at Elevated Temperature, McGraw-Hill, New York, 1961 R.W. Guard, Prod. Eng., Vol 27 (No. 10), 1956, p 160–174 I. Finnie and W.R. Heller, Creep of Engineering Materials, McGraw-Hill, New York, 1959 E.N. da C. Andrade and B. Chalmers, Proc. R. Soc. (London) A, Vol 138, 1932, p 348 J.B. Conway, Trans. Metall. Soc. AIME, Vol 223, 1965, p 2018 F. Garofalo, Properties of Crystalline Solids, STP 283, ASTM, 1965 W. Rosenhahn and D. Ewen, J. Inst. Met., Vol 10, 1913, p 119 W.J. Plumbridge and D.A. Ryder, Metall. Rev., Vol 14, 1969, p 136 P.J. Forsyth and C.A. Stubbington, J. Inst. Met., Vol 83, 1955, p 395 W.A. Wood, Some Basic Studies of Fatigue in Metals, Wiley, New York, 1959

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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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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Mechanisms and Causes of Failures in Heat Treated Steel Parts / 89

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

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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

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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

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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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General Aspects of Failure Analysis Waldek Wladimir Bose-Filho and Jose´ Ricardo Tarpani, Universidade de Saõ 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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General Aspects of Failure Analysis / 113

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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The critical flaw size (ac) is exceeded by some preexisting discontinuity or when subcritical cracking mechanisms (for example, fatigue, stress-corrosion cracking, or creep) reach the critical crack size

Failures due to net-section instability typically occur when a damage process such as corrosion or wear reduces the thickness of a structural section. This type of failure can be evaluated by traditional stress analysis or finite element analysis (FEA), which are effective methods in evaluating the effects of loading and geometric conditions on the distribution of stress and strain in a body or structural system. However, stress analyses by traditional methods or FEA do not easily account for crack propagation from preexisting cracks or sharp discontinuities in the material. When a preexisting crack or discontinuity is present, the concentration of stresses at the crack tip becomes asymptotic (infinite) when using the conventional theory of elasticity. In this regard, fracture mechanics is a useful tool, because it is a method that quantifies stresses at a crack tip in terms of a stress-intensity parameter (K). The fracture mechanics of cracking from a discontinuity or crack in a statically loaded component has two possible situations:

The crack reaches a critical length with rapid (brittle) separation. The crack blunts, redistributing the stress state, with continued loading creating a tear zone (and sharpened crack-tip radius) in front of the crack. In steels, this tear zone can then cause the critical crack length to be exceeded, such that unstable cleavage fracture occurs or unstable microscale ductile fracture is induced.

Which event occurs depends on the temperature and the loading rate, but in either event, crack propagation is unstable (i.e., does not require an increasing load after creation of the tear zone). Fracture mechanics is a tool to help evaluate the implications of preexisting discontinuities or cracks. Testing under Simulated Service Conditions. During the concluding stages of an investigation, it may be necessary to conduct tests that simulate the conditions under which failure is believed to have occurred. Often, simulated-service testing is not practical because elaborate equipment is required, and even where practical it is possible that not all of the

service conditions are fully known or understood. Corrosion failures, for example, are difficult to reproduce in a laboratory, and some attempts to reproduce them have given misleading results. Serious errors can arise when attempts are made to reduce the time required for a test by artificially increasing the severity of one of the factors—such as the corrosive medium or the operating temperature. Similar problems are encountered in wear testing. On the other hand, when its limitations are clearly understood, the simulated testing and statistical experimental design analysis of the effects of certain selected variables encountered in service may be helpful in planning corrective action or, at least, may extend service life. Most of the metallurgical phenomena involved in failures can be satisfactorily reproduced on a laboratory scale, and the information derived from such experiments can be helpful to the investigator, provided the limitations of the tests are fully recognized. Analysis of All the Evidence, Formulation of Conclusions, and Writing the Report. Before starting this final step, some questions must already be answered: Fracture surface: a. What is the fracture mode? b. Is the origin of the fracture visible? c. What is the relation between the fracture direction and the normal or expected fracture directions? d. How many fracture origins are there? e. Is there evidence of corrosion, paint, or some other foreign material on the fracture surface? f. Was the stress unidirectional or was it reversed in direction? The surface of a part: a. What is the contact pattern on the surface of the part? b. Has the surface of the part been deformed by loading during service or by damage after fracture? c. Is there evidence of damage on the surface of the part by manufacturing, assembling, repairing, or service? Geometry and design: a. Are there any stress concentrations related to the fracture? b. Is the part intended to be relatively rigid, or is it intended to be flexible, like a spring? c. Does the part have a basically flawless design?

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d. How does the part—and its assembly— work? e. Is the part dimensionally correct? Manufacturing and processing: a. Are there internal discontinuities or stress concentrations that could cause a problem? b. If it is a wrought metal, does it contain serious seams, inclusions, or forging problems, such as end grains, laps, or other discontinuities, that could have an effect on performance? c. If it is a casting, does it contain shrinkage cavities, cold shuts, gas porosity, or other discontinuities, particularly near the surface of the part? d. If a weldment was involved, was the fracture through the weld itself or through the heat-affected zone in the parent metal adjacent to the weld? If through the weld, were these problems something like gas porosity, undercutting, underbead cracking, or lack of penetration? If through the heat-affected zone adjacent to the weld, how were the parent metal properties affected by the heat of welding? e. If the part was heat treated, was the treatment properly performed? Material properties: a. Are the mechanical properties of the metal within the specified range, if this can be ascertained? b. Are the properties of the metal suitable for the application? c. Residual and applied stress relationship. The residual-stress system that was within the part prior to fracture can have a powerful effect—good or bad—on the performance of a part. d. What was the influence of adjacent parts on the failed part? e. Were fasteners tight? Assembly: a. Is there evidence of misalignment of the assembly that could have had an effect on the fractured part? b. Is there evidence of inaccurate machining, forming, or accumulation of tolerances? c. Did the assembly deflect excessively under stress? Service conditions: It is important to determine if there were any unusual occurrences, such as strange noises, smells, fumes, or other happenings, that could help explain the

problem. The following questions should also be considered: a. Is there evidence that the mechanism was overspeeded or overloaded? b. Is there evidence that the mechanism was abused during service or used under conditions for which it was not intended? c. Did the mechanism or structure receive normal maintenance with the recommended materials? d. What is the general condition of the mechanism? Environmental reactions: The problems related to the environment can arise anywhere in the history of the part: manufacturing, shipping, storage, assembly, maintenance, and service. None of these stages should be overlooked in a thorough investigation that asks: a. What chemical reactions could have taken place with the part during its history? b. To what thermal conditions has the part been subjected during its existence? Report writing: Finally, the report analyzing the failure should be written in a clear, concise, logical manner. It should be clearly structured with sections covering the following (Ref 6): a. Description of the failed item b. Conditions at the time of failure c. Background history important to the failure d. Mechanical and metallurgical study of the failure e. Evaluation of the material quality f. Discussion of any anomalies g. Discussion of the mechanism or possible mechanisms that caused the failure h. Recommendations for the prevention of future failures or for action to be taken with similar pieces of equipment Irrelevant data should be omitted, and, depending on the nature of the problem and the data, not every report will need full treatments for every one of the sections listed previously. Many times, the readership may include purchasing, operating, or accounting personnel who are not technically trained. If this is the situation, the report should be written so that it is comprehensible to these persons. At least, those sections of the report that bear on their decision-making or information needs should be written in language that is accessible to them. Frequently, a cover letter summarizing the most important findings and the suggested action is a

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good vehicle for reaching top executives who are not as interested in the technical specifics but need key findings and recommendations as a basis for decision making. Followup on the recommendations is frequently a difficult task but should be undertaken for the more critical failures. Cooperation between the investigator, the designer, the manufacturer, and the user is critical in developing good, workable changes.

Fracture The process of fracture, in general terms, can be described in terms of the mechanisms of crack initiation and/or crack extension (growth). Different mechanisms may occur for crack initiation and the subsequent process of crack growth. For example, crack extension may occur by the brittle mechanism of cleavage, even though extensive elongation accompanied or preceded crack initiation. The fracture may be classified as either ductile or brittle, depending on whether the mechanism is describing crack initiation or crack growth, respectively. Likewise, the low-energy catastrophic fracture of a high-strength aluminum alloy by microvoid coalescence is also difficult to classify because, although the fracture energy is low and failure initiates by fracture or decohesion of brittle particles, the growth and coalescence of the microvoids occurs by plastic deformation. Another difficulty is that cleavage fracture may be initiated by dislocation interactions that, by definition, involve plasticity. This is why fractures are sometimes difficult to logically classify (Ref 5). Therefore, it is helpful to be clear whether fracture mechanisms are describing the process of crack initiation or extension. Crack extension also can be multimode over time (e.g., fatigue crack growth followed by overload). In terms of fracture appearances (or fracture modes, defined earlier in the section “Determination of Failure Mechanism” in this chapter), a general summary of the visual and microscopic aspects of fracture surfaces for metallic materials is provided in Table 1 (Ref 8). Several analytical procedures are available for distinguishing among the various types of fracture. For example, the presence or absence of plastic macrodeformation can be determined with the unaided eye or by use of a steel scale, a machinist’s micrometer, or a machinist’s or measuring microscope. Differences in some dimensional attribute of parts (such as width or

thickness) at and well away from the fracture can serve to define macrodeformation after assurance that both points of measurement had the same dimension before fracture. Fracture-surface matching is also used to determine the presence or absence of plastic deformation. It is very important, however, to resist the temptation to fit the matching fracture surfaces together, because this almost always destroys (smears) microscopic features. The fracture surfaces should never actually touch during fracture-surface matching. The origin of a fracture may be indicated by a discoloration or by the topography of the fracture surface. A discolored area on a fracture surface may be produced by a preexisting crack whose surfaces have been corroded or oxidized. For example, the surfaces of a quench crack can be oxidized during a subsequent tempering heat treatment; the oxide film gives a bluish-black color to the surfaces of the crack. Topographical features that often reveal the origin of a fracture are either chevron or river patterns or a set of diverging ledges. If the fracture surface is essentially featureless, the presence of a shear lip can be used to locate, within limits, the origin of a fracture. For example, a shear lip is not formed at the origin of a stress-corrosion crack, but when the crack begins to propagate rapidly, a shear lip is formed wherever the crack front exits from the interior to the free surface. Beach marks, which are associated with fatigueinitiated fractures, also provide a definite indication of the crack origin; however, it should be noted that fracture surfaces having an appearance similar to that of the beach-mark pattern can be produced by stress corrosion. Generally, cyclic loading produces only a single crack, which is usually located at a site of stress concentration or of a metallurgical defect, whereas additional cracks, formed independently of the main crack and at a distance from it, may be observed on the surface of a structural or machine component subjected to corrosion fatigue or stress corrosion. On the microscopic level, striations on the fracture surface are unique to fatigue, and the crack path, although normally transgranular, can be intergranular. For example, intergranular fatigue cracking can occur in the case of a carburized steel or in a material that has a high density of second-phase particles at the grain boundaries. Corrosion-fatigue and stress-corrosion cracks may propagate transgranularly, intergranularly,

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Table 1 Fracture mode identification chart Instantaneous failure mode(a) Method

Visual, 1 to 50· (fracture surface)

Ductile overload

Necking or distortion in direction consistent with applied loads Dull, fibrous fracture Shear lips

Brittle overload

Progressive failure mode(b) Fatigue

Corrosion

Wear

Creep

Multiple brittle Gouging, Flat progressive General Little or no appearing fissures abrasion, wastage, roughzone with beach distortion External surface ening, pitting, or polishing, marks Flat fracture and internal or erosion trenching Bright or coarse Overload zone fissures contain consistent with Stress-corrosion Galling or storing texture, reaction scale in direction of and hydrogen applied loading crystalline, coatings motion damage may direction grainy Fracture after create multiple Roughened areas Ratchet marks Rays or limited with compacted cracks that where origins chevrons dimensional powdered debris appear brittle join point to origin change (fretting) Smooth gradual transitions in wastage

Path of penetra- Wear debris and/or Multiple Progressive Cleavage or Microvoids Scanning intergranular abrasive can be tion may be zone: worn intergranular (dimples) electron characterized as to fissures covered irregular, appearance, fracture elongated microscopy, with reaction scale morphology and intergranular, or flat, may show Origin area may in direction of 20 to Grain faces may a selective phase composition striations at contain an loading 10,000· show porosity Rolling-contact attacked magnification imperfection Single crack with (fracture fatigue appears EDS may above 500 · or stress no branching surface) like wear in early help identify Overload zone: concentrator Surface slip band stages corrodent(c) may be either emergence ductile or brittle Metallographic inspection, 50 to 1000· (cross section)

Contributing factors

Grain distortion and flow near fracture Irregular, transgranular fracture

Load exceeded the strength of the part Check for proper alloy and processing by hardness check or destructive testing, chemical analysis Loading direction may show failure was secondary Short-term, high-temperature, high-stress rupture has ductile appearance (see creep)

Microstructural May show General or Little distortion Progressive change typical of localized surface localized zone: usually evident overheating distortion at attack (pitting, transgranular Intergranular or surface consistent Multiple intercracking) with little transgranular granular cracks with direction of Selective phase apparent May relate to Voids formed on motion attack distortion notches at surface or brittle Overload zone: Thickness and Identify embedded grain boundaries or wedge-shaped particles morphology of phases internally may be either cracks at grain corrosion scales ductile or brittle triple points Reaction scales or internal precipitation Some cold flow in last stages of failure Load exceeded the dynamic strength of the part Check for proper alloy and processing as well as proper toughness, grain size Loading direction may show failure was secondary or impact induced Low temperatures

Attack morphol- For gouging or Cyclic stress abrasive wear: ogy and alloy exceeded the check source of type must be endurance limit abrasives evaluated of the material Evaluate effec Check for proper Severity of tiveness of lubristrength, surface exposure cants finish, assembly, conditions may Seals or filters may be excessive; and operation have failed Prior damage by check: pH, Fretting induced temperature, mechanical or by slight looseness corrosion modes flow rate, in clamped joints dissolved may have subject to oxidants, elecinitiated vibration trical current, cracking metal coupling, Bearing or materi Alignment, als engineering aggressive vibration, design may reduce agents balance or eliminate High cycle low Check bulk composition and problem stress: large Water contaminants fatigue zone; contamination low cycle high High velocities stress: small or uneven flow fatigue zone distribution, cavitation

Mild overheating and/or mild overstressing at elevated temperature Unstable microstructures and small grain size increase creep rates Ruptures occur after long exposure times Verify proper alloy

(a) Failure at the time of load application without prior weakening. (b) Failure after a period of time where the strength has degraded due to the formation of cracks, internal defects, or wastage. (c) EDS, energy-dispersive spectroscopy. Compiled by C.R. Morin, S.L. Meiley, and Z.B. Flanders, Packer Engineering Associates, Inc.


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or by a combination of both modes. A distinguishing feature of stress corrosion is the branching of the main crack. If corrosion pits or corrosion products are found only on the slow-growth region of a fracture surface, the environment was in all probability sufficiently corrosive to affect the fracture mechanism. However, if evidence of corrosion is found on both the slow-growth and fast-growth areas, some corrosion took place subsequent to fracture, and the environment may or may not have influenced fracture. Ductile Fracture Ductile fracture takes place when a material capable of undergoing plastic deformation is subjected to stresses that culminate in its rupture. Macroscopically, the ductile fracture process presents some peculiarities that allow it to be identified immediately. The first feature is the presence of plastic deformation that may be accompanied by neck formation. In tensile testpieces of ductile materials, besides necking, the fracture surface presents a fibrous aspect and a cup-cone geometry, as seen in Fig. 1.

Fig. 1

The fracture process begins in the center of the testpiece with microvoid nucleation along grain boundaries or from interfaces such as those found in base metal/inclusions boundaries. As the applied stress increases, microvoids grow and coalesce, forming a crack in the center of the part. This process, depicted in Fig. 2, ends up in rapid crack propagation by shearing of the remaining ligament of the neck region, at an angle of 45 in relation to the loading direction. It is important to emphasize that a cup-cone geometry will depend on the geometry and dimensions of the part and mechanical properties of the material. Thin sheets, for instance, present neck formation and a fracture surface oriented at an angle of 45 in relation to the applied load, as observed in Fig. 3. Ductile fracture takes place intergranularly, unless some sort of mechanism weakens the grain boundaries. The microscopic aspect of the fracture surface consists of several small elliptical cavities, or microvoids, as depicted in Fig. 4. Brittle Fracture Brittle fracture occurs with little or no plastic deformation. This type of fracture is often

Fig. 3

Thin sheet testpiece of a low-carbon steel after fracture

Fig. 4

Microvoids on the fracture surface of AA6061-T1 tensile testpiece

Ductile fracture showing the typical cup-cone geometry

Microvoids

Fig. 2

Schematic representation of the cup-cone geometry formation during the ductile fracture process

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associated with materials of high strength and low ductility or materials that were subjected to an embrittlement process. The crack, once nucleated, propagates very quickly in a direction perpendicular to the applied load. Figure 5 presents an example of a gray cast iron testpiece that presented brittle fracture. Besides the mechanical properties, several other factors may result in a brittle behavior, such as temperature, loading rate, presence of stress concentrators, and dimensions. Low temperatures tend to reduce the ductility of metals, especially those possessing a body-centered cubic structure, resulting in a typically brittle fracture. Figure 6 shows that as the temperature drops, the brittle aspect on the fracture surface of impact testpieces increases. The presence of stress raisers or larger dimensions introduces a more severe triaxial stress state within the material, and thus, there is larger probability that brittle fracture will occur. However, it is known

Fig. 5

that the superposition of high hydrostatic stresses on the material reduces the triaxiality levels, increasing ductility. High applied loading rates are likely to make plastic deformation more difficult because shearing processes are timedependent, resulting in brittle behavior. Crack propagation by brittle fracture can occur across the grains (transgranular) or along the grain boundaries (intergranular). In the transgranular mode, the fracture process takes place by cleavage along specific crystallographic planes. Figure 7 presents cleavage regions in a microalloyed low-carbon steel, which can be identified by flat regions on the fracture surface. Additionally, it is worth mentioning that most parts of steels will present alternate regions consisting of cleavage areas and microvoids, evidencing a mixed mode of crack propagation. In another situation, fracture can take place intergranularly, because the grain boundary is a

Tensile testpiece of gray cast iron presenting brittle fracture

Fig. 7

Fig. 6

Fracture surfaces of SAE 4140 impact testpieces. Tested at room temperature, right, and at 196 C, left

(a) Cleavage region observed in low-carbon steel. (b) Magnification of the region delimited by the rectangle in (a) showing an inclusion in the center of the cleavage region

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weaker path for crack propagation. Normally, this fracture mode will occur when some embrittlement process resulted in grain boundaries being more susceptible to crack propagation than the core of the grain, such as an unsuitable heat treating or by environmental factors. Figure 8 presents an example of intergranular brittle fracture in an austenitic stainless steel SAE 316L, where grain boundaries can clearly be observed on the fracture surface. Fatigue Fracture According to the definition given by ASTM E1823, fatigue is “the process of progressive localized permanent structural change occurring in a material subjected to conditions that produce fluctuating stresses and strains at some point or points and that may culminate in cracks or complete fracture after a sufficient number of fluctuations.” A material subjected to fatigue can fracture at applied stresses much lower than those necessary to fracture the same material under monotonic conditions. The fluctuating stresses can be originated from mechanical, thermal, or vibration loading conditions, and the phenomenon is responsible for more than 80% of mechanical failures of components. For more than 150 years, the study of metals fatigue has involved engineers, physicists, chemists, and mathematicians, and everyday this study becomes more and more complex and important. The theory about fatigue is extremely vast, and for each question answered, another one, more instigating, appears, requiring a broad knowledge of materials science. In the following topics, a brief overview is given about the main mechanisms and factors influencing the fatigue

life of a component during both the nucleation and crack propagation phases. Fatigue Crack Initiation. Generally, fatigue cracks are initiated at free surfaces, where there is no constraint to material deformation; however, in some cases, cracks may be initiated in the interior of the material where interfaces are present, such as the interface of a carburized surface layer and the base metal or the interface of an inclusion and the base metal, or from gas bubbles. In other cases, subsurface cracks were found to nucleate below the surface where high compressive residual stresses were introduced by shot peening or surface rolling. One of the classic models of fatigue crack nucleation considers that when a material is under loading (monotonic or cyclic), slips occur at the high-shear-stress planes, creating steps on the material surface. Under cyclic loading, the formation of intrusions and extrusions is observed, as schematically represented in Fig. 9. Slip band intrusions are excellent stress raisers that can be sites of crack nucleation. Besides the applied stress amplitude, DS/2, several other factors are likely to affect the nucleation of a fatigue crack, such as the mean stress, Sm, or load ratio, R; geometry and surface finishing of the part; mechanical properties; and environment. Here, the R ratio is defined as the ratio between the minimum and maximum loads during the fatigue cycle. A large proportion of fatigue data found in the literature refers to tests conducted at Sm = 0, that is, for a load ratio R = 1. However, in many engineering situations, the fluctuating stresses are superimposed to a static stress. Larger mean stresses reduce the nucleation time because they facilitate the plastic deformation mechanism associated with this phenomenon. In an S-N graph, this can be represented by curves shifted to the left and down, as represented in Fig. 10.

Intrusion

Extrusion

Metal Surface

Fig. 8

Fractograph of SAE 316L showing intergranular brittle fracture

Fig. 9

Schematic representation of an intrusion formation on the surface of a metallic material


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The mechanism proposed in Fig. 9 is adequate to explain the initiation of cracks on polished testpieces or components without the presence of geometric discontinuities. However, in engineering components, there are several stress concentrators, such as scratches, notches, machining marks, corrosion pits, and microconstituents such as grain boundaries, triple points, and inclusions, that individually or synergistically can reduce the initiation time. Since the initiation depends essentially on plastic deformation mechanisms, high-strength materials normally present a higher resistance to fatigue crack nucleation. In this sense, several surface-hardening treatments are employed to selectively reinforce the material, aiming to retard crack initiation and therefore to increase fatigue life. The chemical composition and/or the microstructure of the surface can be modified by thermochemical treatments, such as carburizing or nitriding, or by cold deformation processes, such as shot peening or surface rolling. Mechanical parts that necessarily present stress concentrators, such as crankshafts, gears, and bolts, can be subjected to these treatments to increase the fatigue limit of the material. Figure 11 shows a micrograph of the transverse section of a bolt, where the thread was cold formed by surface rolling. As a consequence, surface grains are flattened due to the mechanical deformation imposed. In this case, besides increasing hardness and mechanical strength, the process avoids the introduction of harmful machining marks. Surface treatments may also increase fatigue life by the introduction of compressive residual stresses on the surface of the material. As long as the material remains in linear elastic conditions, the principle of stress superposition can be

employed to describe the actual stress state in materials containing residual stresses. Therefore, the effective stress, S0 , is given by the sum of the applied stress, S, to the residual stress, Sres: S0 =S+Sres

(Eq 1)

Similarly, the effective minimum and maximum stresses are defined, respectively, as: S0max =Smax +Sres

(Eq 2)

S0min =Smin +Sres

(Eq 3)

Consequently, the effective stress amplitude, mean stress, and load ratio are given, respectively, by: DS0 S0max 7S0min (Smax +Sres )7(Smin +Sres ) = = 2 2 2 Smax 7Smin DS = = (Eq 4) 2 2 S0max +S0min (Smax +Sres )+(Smin +Sres ) = 2 2 Smax +Smin +Sres =Sm +Sres = (Eq 5) 2

S0m =

R0 =

S0min S +Sres = min S0max Smax +Sres

(Eq 6)

Therefore, the presence of a residual-stress field does not affect the stress amplitude but affects the mean stress and the load ratio. A compressive residual stress reduces the mean stress and the load ratio, increasing the number

Increasing Sm ∆S/2

Nf

Fig. 10

Mean stress effect on S-N fatigue curves

Fig. 11

Optical micrograph of the transverse section of a thread fillet machined by surface rolling. The material consists of duplex stainless steel


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of cycles for crack nucleation and vice versa. In some situations, where high surface compressive residual stresses are found, such as in materials subjected to surface-hardening treatments, a crack may initiate below the surface, where the compressive residual-stress level is lower. An example of subsurface crack nucleation is observed in Fig. 12 for a surface-rolled ductile cast iron subjected to bending-rotating fatigue. Fatigue Crack Propagation. Basically, fatigue crack propagation can be divided into three stages: stage I (short cracks), stage II (long cracks), and stage III (final fracture). A fatigue crack, once initiated, propagates along high shear-stress planes (45 ), as schematically represented in Fig. 13. This is known as stage I or the short crack growth propagation stage. The crack propagates until it is decelerated by a microstructural barrier, such as a grain boundary, inclusions, or pearlitic zones, that cannot accommodate the initial crack growth direction. Therefore, grain refinement is capable of increasing fatigue strength of the material due to the insertion of a large quantity of microstructural barriers, that is, grain boundaries, that must be overcome in stage I of propagation. Surface mechanical treatments, such as shot

Fig. 12

Probable subsurface crack nucleation site in a surface-rolled ductile cast iron testpiece tested under bending-rotating conditions

Stage II

Surface

Stage I

Fig. 13

Stages I and II of fatigue crack propagation

peening and surface rolling, contribute to the increase in the number of microstructural barriers per unit of length due to the flattening of the grains. When the stress-intensity factor, K, increases as a consequence of crack growth or higher applied loads, slips start to occur in different planes close to the crack tip, initiating stage II of propagation. While stage I of propagation is orientated 45 in relation to the applied load, propagation in stage II is perpendicular to load direction, as depicted in Fig. 13. An important characteristic of stage II propagation is the presence of ripples on the fracture surface, known as striations, which are only visible with the aid of a scanning electron microscope. Not all engineering materials exhibit striations. They are clearly seen in pure metals and many ductile alloys, such as aluminum alloys. In steels, they are frequently observed in cold-worked alloys. Figure 14 shows examples of fatigue striations in an interstitial-free steel and in aluminum alloys. The most accepted mechanism for the formation of striations on the fatigue fracture surface of ductile metals (Ref 9) consists of successive blunting and resharpening of the crack tip, as represented in Fig. 15. Finally, stage III is related to the unstable crack growth as Kmax approaches KIc. At this stage, crack growth is controlled by static modes of failure and is very sensitive to the microstructure, load ratio, and stress state (planestress or plane-strain loading). Macroscopically, the fatigue fracture surface can be divided into two distinct regions, as shown by Fig. 16. The first region corresponds to the stable fatigue crack growth and presents a smooth aspect due to the friction between the crack-wake faces. Sometimes, concentric marks, known as beach marks, can be seen on the fatigue fracture surface as a result of successive arrests or decrease in the fatigue crack growth rate due to a temporary load drop or to an overload that introduces a compressive residualstress field ahead of the crack tip. The other region corresponds to the final fracture and presents a fibrous and irregular aspect. In this region, the fracture can be either brittle or ductile, depending on the mechanical properties of the material, dimensions of the part, and loading conditions. The exact fraction of area of each region will depend on the applied load level. High applied loads will result in a small stable fatigue crack propagation area, as depicted in Fig. 16(a). On the other hand,

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if lower loads are applied, the fatigue crack will have to grow longer before the applied stress-intensity factor, K, reaches the fracture toughness value of the material, resulting in a smaller area of fast fracture (Fig. 16b). Ratcheting marks are another macroscopic feature that can be observed in fatigue fracture surfaces. These marks originate when multiple

(a)

(d) (b) (e)

(c)

Fig. 15

Proposed mechanisms of striation formation in stage II of propagation. (a) No load. (b) Tensile load. (c) Maximum tensile load. (d) Load reversion. (e) Compressive load. Source: Ref 9

Fig. 14

Fatigue striations in (a) interstitial-free steel and (b) aluminum alloy AA2024-T42. (c) Fatigue fracture surface of a cast aluminum alloy where a fatigue crack was nucleated from a casting defect, presenting solidification dendrites on the surface. Arrow at top right indicates fatigue striations.

Fig. 16

Fatigue fracture surface. (a) High applied load. (b) Low applied load


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cracks, nucleated at different points, join together, creating steps on the fracture surface. Therefore, counting the number of ratchet marks is a good indicator of the number of nucleation sites. Figure 17 presents in detail some ratchet marks found on the fracture surface of a large SAE 1045 rotating shaft, fractured by fatigue. Similar to the initiation phase, many factors can affect long fatigue crack propagation rates. Among them, special attention should be given to the effects of load ratio and the presence of residual stresses. Increasing the load ratio has a tendency to increase the long crack growth rates in all regions of the fatigue crack growth rate versus applied stress-intensity factor range curve, or simply, da/dN versus applied DK curve. Generally, the effect of increasing load ratio is less significant in the Paris regime than in nearthreshold and near-failure regions (Fig. 18).

Fig. 17

Near the threshold stress-intensity factor, DKth, the effects of R ratio are mainly attributed to crack closure effects, where crack faces come in contact at an applied Kcl that is higher than the minimum applied stress-intensity factor, Kmin. Several different mechanisms may contribute to premature crack closure. One of them consists of plasticity-induced closure, represented in Fig. 19(a). As the crack grows, the material that has been previously permanently deformed within the plastic zone now forms an envelope of plastic zones in the wake of the crack front. This leads to displacements normal to the crack surfaces as the restraint is relieved. This is no problem while the crack is open; however, as the load decreases, the crack surfaces touch before the minimum load is reached, shielding the crack. This type of premature contact can also occur due to the crack-wake roughness and irregularities (Fig. 19b) or by the presence of corrosion subproducts, such as oxides (Fig. 19c). As observed in Fig. 20, the effect of closure produces a reduction in the effective DK range because of the increase in the effective Kmin, reducing the driving force for fatigue crack growth. The effect is more significant near the threshold region because the crack tip opening displacements are smaller and the crack faces are closer to each other. Additionally, for the same applied DK, higher R ratios increase the applied values of Kmax and Kmin, increasing DKeff. For most materials, the Paris regime is considered closure-free and Kmax-independent, and

Ratcheting marks, indicated by the arrows, in an SAE 1045 shaft fractured by fatigue

Plastic deformation envelope Final failure

Plastic zone Crack tip

(a)

da/dN

Paris´ regime

Increasing R (b)

Premature contact points

Near threshold

Oxides ∆K (c)

Fig. 18

Schematic representation of the R ratio effect on fatigue crack growth curves. The near-threshold, Paris regime, and final failure regions are also indicated on the curves.

Fig. 19

Crack closure mechanisms induced by (a) plasticity, (b) roughness, and (c) oxide


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0 j0, then: If Kmin

the crack growth rates are generally very similar for tests conducted under different R ratios. Near the final failure, the effects of R ratio are related to the higher monotonic fracture component as Kmax approaches KIc. Therefore, for the same applied DK, Kmax values are higher for tests conducted under higher applied R ratios, and consequently, da/dN values are higher. The effects of residual stress on fatigue crack growth are related to alterations in the R ratio and in the applied DK. In other terms, the residual stresses affect the two parameters that control the crack driving force, that is, Kmax and DKeff. When a crack is introduced in a plate subjected to a residual-stress field, a residual stress-intensity factor, Kr, arises that can either decrease or increase the crack driving force parameters. The superposition principle can also be applied in terms of the stress-intensity factor, provided that the material remains linearly elastic. In this sense, Kr can be added to Kmax and Kmin: 0 Kmax =Kmax +Kr

(Eq 7)

0 =Kmin +Kr Kmin

(Eq 8)

0 Kmin K +Kr = min 0 Kmax Kmax +Kr

(Eq 11)

0 =Kmax +Kr DK 0 =Kmax

(Eq 12)

It is important to note that these equations assume that the part of the fatigue cycle during which the crack is closed at its tip (i.e., K050) makes no contribution to crack growth.

Distortion Distortion is the least serious mode of failure, but it can lead a part to failure or a structure to collapse. It is easy to recognize but very difficult to prevent. This is due to the fact that distortion does not involve the part itself but its use and design. There are four reasons for distortion: yielding, buckling, creep, and residual stresses. Yielding. When a load is put on a part, and it causes the part to be permanently distorted, it is unable to perform the intended function and therefore must be considered failed. In a welldesigned part, the stresses never exceed the yield point, and the part deforms only elastically; that is, when the load is released, the part returns to its original dimensions. In a good design, the part operates in the elastic range, that is, below yielding point; beyond this, the part will be permanently deformed, and greater loads will cause the part to actually break. This point is considered to be a very basic point to design and applies when the load on a part is applied in a quasi-static way, such as the load on a building structure or the stress in the legs of a desk. A ductile failure is

As a result, R0 and DK0 are defined as follows. If 0 40, then: Kmin R0 =

R0 =0

(Eq 9)

0 0 DK 0 =Kmax 7Kmin =ðKmax +Kr Þ7ðKmin +Kr Þ =Kmax 7Kmin =DK

(Eq 10)

Kmax Kmax ∆Kap=∆Keff K

K

∆Keff

∆Kap

Kmin Kcl

Kcl

Kmin Time

Time (a)

Fig. 20

(b)

Load ratio effect on DKeff in a fatigue cycle. (a) Kmin5Kcl. (b) Kmin4Kcl


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one where there is a great deal of distortion of the failed part. Commonly, a ductile part fails when it distorts and can no longer carry the needed load. However, some ductile parts break into two pieces and can be identified because there is a great deal of distortion around the fracture face, similar to what would happen if too much is placed load on a low-carbon steel bolt. Buckling. The failure of an engineering component is not always caused by materials fracture. In many occasions, the component distortion may be sufficient to put it out of function. The distortion can be elastic or plastic. The elastic distortions are temporary; however, they may be sufficient to cause interference on the mobile parts. The plastic distortion is permanent and can be a result of an overload or creep deformation. The overload causes permanent plastic deformation when the material yield limit is overcome. This may happen in the presence of stress concentrators, high temperature, inadequate heat treatment, or incorrect materials selection for the component application. Compressive overloads may lead the material to overcome the buckling strength limit, such as the one shown in Fig. 21 for an aluminum part. The buckling strength is essentially a design problem (not metallurgical), and the load depends on the dimensions of the part and the Young’s modulus of the material (the only materials factors involved). Creep is a time-dependent phenomenon that causes a part failure if it is under both quasistatic load and temperatures higher than 0.3 Tm (absolute melting temperature). Creep strain may produce sufficiently large deformation or

distortion that a part can no longer perform its intended function. The two general types of creep processes are grain-boundary sliding and voids at grain boundaries (cavitation creep). The creep processes are easily identified by the local ductility and large numbers of intergranular cracks that will depend on the temperature and strain rate imposed. In general, a high strain rate combined with high temperature results in ductile fracture, followed by a large elongation and neck formation. Additionally, the grains near the fracture surface tend to be elongated. On the other side, the combination of low strain rate and high temperature results in intergranular brittle fracture, with low elongation or necking. Intergranular fracture in such conditions normally initiates by grain-boundary sliding from triple points or at grain-boundary intersections with second-phase particles, causing cavities on the material microstructure, as presented in Fig. 22. Once the crack nucleates, it propagates by grain boundaries, and given that some significant plastic deformation may take place, the fracture surface tends to exhibit grains of equiaxial shape. Therefore, to increase creep strength, the material is normally heat treated to increase the grain size, reducing the ratio between the grain surface area and volume. In turbines that work at very high temperatures, the creep mechanism must be considered. In this case, the component may be produced from monocrystals that significantly increase the creep resistance. Most creep curves show three distinct stages (Fig. 23). After the elastic strain, there is a region of increasing plastic flow at decreasing rate (first stage), followed by a region of approximately constant strain rate (secondary stage), and finally a region of intense increase in the strain rate, which rapidly extends to fracture (third stage).

Cavities

(a)

Fig. 22 Fig. 21

Aluminum part that suffered buckling

(b)

Intergranular crack formation at high temperature by grain-boundary sliding at (a) triple points and (b) inclusions

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Residual stresses can play a significant role in explaining or preventing failure of a component. One example of residual stresses preventing failure is the use of shot peening processes that increase the fatigue life of a component by inducing surface compressive stresses. Unfortunately, there are also processes or processing errors that can induce excessive tensile residual stresses in locations that may promote failure of a component. The internal state of stress is caused by thermal and/or mechanical processing of the parts. Common examples of these are bending, rolling, or forging a part. Thermal residual stresses are primarily due to differential expansion when a metal is heated or cooled. Two control factors are thermal treatment (heating or cooling) and restraint. Both the thermal treatment and restraint of the component must be present to generate residual stresses. Residual stresses can result in visible distortion of a component. However, in the case of residual stresses, the distortion can also be useful in estimating the magnitude or direction of these stresses.

Wear-Assisted Failure Wear may be defined as damage to a solid surface caused by the removal or displacement of material by the mechanical action of a contacting solid, liquid, or gas. It may cause significant surface damage, and the damage is usually thought of as gradual deterioration. While the terminology of wear is unresolved, the following categories are commonly used: adhesive wear, abrasive wear, erosive wear, fretting, cavitation, rolling, contact fatigue, and corrosive wear.

Adhesive wear has been commonly identified by the terms galling or seizing. It is caused by the material transference from one surface to another during their relative movement due to a solid-state welding process. Figure 24 shows a schematic representation of this process. High contact pressure among the surface roughness results in local plastic deformation and points of microwelding. The movement between the surfaces causes the rupture of the junctions, resulting in a rough peak in one surface and a valley on the other. Eventually, the tip of a peak may break, and an abrasive particle is formed. Abrasive wear, or abrasion, is caused by the displacement of material from a solid surface due to hard particles or protuberances sliding along the surface. The particles may be found free between two surfaces or attached to one of them, and the wear level depends on the relative hardness between the particle and the surface (Fig. 25). The abrasion may also happen due to the protuberances or sharp asperities on one of the surfaces in contact. The process of abrasive erosion may be considered as abrasive wear. Erosion, or erosive wear, is the loss of material from a solid surface due to relative motion in contact with a fluid that contains solid particles. In this case, the particle is found to be dispersed in a fluid or gas means, and it reaches the surface under relatively high velocity (Fig. 25d). Figure 26 shows the microstructure of the transversal section of an H11 tool steel that has been subject to abrasive erosion. Fatigue wear can be characterized by the formation of cracks superficially and/or subsuperficially and the removal of posterior material due to cyclic loading of solid surfaces. The sliding contact and/or rolling between solids Adhesion

Fracture X ∆ε ∆t = creep rate ε

∆ε ∆t

ε0

Stage I

Stage II

Particle

Stage III

Time

Fig. 23

Schematic strain-time curve at constant load and temperature showing the three stages of creep

Fig. 24

Transference mechanism of a material from one surface to another and the formation of an abrasive particle in the process of adhesive wear

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or the repetitive impact of solids and/or liquids in a surface are responsible for the superficial fatigue. When two surfaces of this nature interact due to load application, the area effectively in contact may be very small, resulting in high compressive and shear stresses that may lead to crack nucleation. If only rolling is present, the maximum shear stress takes place just below the surface, giving rise to cracks that propagate parallel to the surface and emerge at the surface, causing part of the material to separate from the component, as shown in Fig. 27. However, pure rolling is not found in inservice conditions. Normally, there is some sliding between the two surfaces, which alters the stress field due to an increase in the shear component, displacing the resulting stress closer to the surface. The cracks start to nucleate on the component surface, propagating at a very shallow angle, as shown in Fig. 28.

Fretting fatigue is considered a phenomenon where the damage is introduced by a conjunction of events consisting of adhesion, oscillatory movement of very low amplitude, oxidation, and abrasion. The small oscillatory movements may cause points of adhesion on the surface that eventually break, forming oxidized particles that

Fig. 27

(a)

(b)

(c)

(d)

Schematic representation of contact fatigue under pure rolling between two surfaces

Fig. 25

Abrasive wear. (a) Free particle between two surfaces. (b) Particle attached to one of the surfaces. (c) Sharp asperity. (d) Erosion

Fig. 28

Fig. 26

Fractography showing an H11 tool steel that has suffered abrasive erosion

Damage by contact fatigue in rolling combined with sliding conditions in gears produced from a quenched and tempered AISI 8620 carburized steel. (a) Transversal section. (b) Frontal view from a formed cavity

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act as abrasives on the surface, since the smallamplitude movements avoid their dispersion apart from the source point. Figure 29 presents a micrograph from a plasma nitrided Cr-Mo-V steel, where a microcrack formed in the fretting region. More than one mechanism can be responsible for the wear observed on a particular part. The most critical function provided by lubricants is to minimize friction and wear to extend equipment service life. Gear failures can be traced to mechanical problems or lubricant failure. Lubricant-related failures are usually traced to contamination, oil film collapse, additive depletion, and use of improper lubricant for the application. The most common failures are due to particle contamination of the lubricant. Dust particles are highly abrasive and can penetrate through the oil film, causing plowing wear or ridging on metal surfaces. Water contamination can cause rust on working surfaces of gears and eventually destroy metal integrity. To prevent premature failure, gear selection requires careful consideration of the following: gear tooth geometry, tooth action, tooth pressures, construction materials and surface characteristics, lubricant characteristics, and operating environment.

Environmentally Assisted Failure Corrosion is chemically induced damage to a material that results in deterioration of the material and its properties. Corrosion can seldom be totally prevented, but it can be minimized or controlled by proper choice of material, design, coatings, and occasionally by changing the environment. Various types of

Fig. 29

Fretting fatigue at the surface of a Cr-Mo-V steel

metallic and nonmetallic coatings are regularly used to protect metal parts from corrosion. Corrosion may result in failure of the component. Several factors should be considered during a failure analysis to determine the effect corrosion played in a failure, such as type of corrosion, corrosion rate, the extent of the corrosion, and the interaction between corrosion and other failure mechanisms. Uniform, pitting crevice, galvanic, and stresscorrosion cracking are the most common types of corrosion. Uniform corrosion is characterized by corrosive attack proceeding evenly over the entire surface area or a large fraction of the total area. General thinning takes place until failure. On the basis of tonnage wasted, this is the most important form of corrosion. Stress-corrosion cracking necessitates a tensile stress, which may be caused by residual stresses, and a specific environment to cause progressive fracture of a metal. Aluminum and stainless steel are well known for stresscorrosion cracking problems. However, all metals are susceptible to stress-corrosion cracking in the right environment. Pitting corrosion is a localized form of corrosion by which cavities or holes are produced in the material. Pitting is considered to be more dangerous than uniform corrosion damage because it is more difficult to detect, predict, and design against. Corrosion products often cover the pits. A small, narrow pit with minimal overall metal loss can lead to the failure of an entire engineering system. Pitting corrosion, which, for example, is almost a common denominator of all types of localized corrosion attack, may assume different shapes. Crevice corrosion is a localized form of corrosion usually associated with a stagnant solution on the microenvironmental level. Such stagnant microenvironments tend to occur in crevices (shielded areas) such as those formed under gaskets, washers, insulation material, fastener heads, surface deposits, disbonded coatings, threads, lap joints, and clamps. Crevice corrosion is initiated by changes in local chemistry within the crevice. Galvanic corrosion (also called dissimilarmetal corrosion or, wrongly, electrolysis) refers to corrosion damage induced when two dissimilar materials are coupled in a corrosive electrolyte. It occurs when two (or more) dissimilar metals are brought into electrical contact under water. When a galvanic couple forms, one of the metals in the couple becomes the anode

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and corrodes faster than it would all by itself, while the other becomes the cathode and corrodes slower than it would alone. 6. REFERENCES

1. D. Dennies, How to Organize a Failure Investigation, ASM International, 2005 2. D.J. Wulpi, Chapter 1: Techniques of Failure Analysis, Understanding How Components Fail, 2nd ed., ASM International, 2000, p 1–11 3. C.R. Brooks and A. Choudhury, Chapter 1: Introduction, Metallurgical Failure Analysis, McGraw-Hill, 1993, p 1–72 4. R. Graham, Strategies for Failure Analysis, Adv. Mater. Process. Aug 2004, p 45–50 5. D.A. Ryder, T.J. Davies, I. Brough, and F.R. Hutchings, General Practice in Failure

7.

8.

9.

Analysis, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986, p 15–46 G.F. Vander Voort, Conducting the Failure Examination, Prac. Fail. Anal., Vol 1 (No 2), April 2001, p 14–46 and Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 2002 A. Tanzer, Determination and Classification of Damage, Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 2002 G. Powell, Identification of Types of Failure, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986, p 75–81 C. Laird, The Influence of Metallurgical Structure on the Mechanisms of Fatigue Crack Propagation, Fatigue Crack Propagation, STP 415, ASTM, p 131–168

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Failure in Steel Forging Md. Maniruzzaman and Richard D. Sisson, Jr., Worcester Polytechnic Institute Stephen R. Crosby, The Stanely Works Charlie Gure (deceased)

IN-PROCESS OR SERVICE FAILURES of forgings may occur for a variety of reasons. The starting material may be of insufficient quality to be adequately formed without cracking, or the forging process may introduce various types of discontinuities that cause failure during services. For example, well-known forging-related discontinuities include:

Laps Bursts Flakes Segregation Cavity shrinkage Centerline pipe Parting-line grain flow Inclusions

Forging discontinuities are discussed in more detail in the texts on forging (Ref 1–4). This article describes six case studies of failures with steel forgings (summarized in Table 1). The case studies illustrate difficulties encountered in either cold forging or hot forging in terms of preforge factors and/or discontinuities generated by the forging process. Tables 2 and 3 summarize these factors for cold and hot forging, respectively. Supporting

topics that are discussed in the case studies include:

Validity checks for buster and blocker design Lubrication and wear Mechanical surface phenomenon Forging process design Forging tolerances

As case studies were being selected, each of the aforementioned supporting topics was reviewed for any impact that particular study had on the case being examined. It is a wellknown fact that forging solutions have several possible avenues to follow. There is no unique theory in plasticity that leads to the solution. Most of the work reported here was performed using the minimum amount of energy to create the particular product. Factors unrelated to the deformation process, such as chemistry, microstructure, phase, grain size, segregation, and prior strain history, are not addressed here. Instead, factors directly related to the deformation process itself are presented in this abbreviated discussion. Wear, plastic deformation processes, and laws of friction are introduced as a group of

Table 1 Failure analysis of steel forgings and components Case study

Crankshaft underfill Tube bending Spade bit Trim tear Upset forging Flow-through laps and avoidance

Defect

Solution

Unable to fill crankshaft flanges with existing press capacity Unable to control exterior wall thinning and interior wall thickening Unable to achieve center web thickness at programmed force and sufficient flow to wings Forge material tore at trimline when forging was trimmed immediately following finish forging Cracking at circumferential bulge after upset

Introduce creep stages for last increment of displacements Introduce induction heating and cooling to limit the heated axial tube length prior to making the bend Adjust the die angle to create more shear stress, enabling full flow to the wings Introduce a delay time after forge and prior to trim, allowing the forge material to cool and gain strength Re-examine the strain and strain rate and process map for stable flow Replace the input piece with a newly designed preform piece, following the design procedures given in this work

Material foldover at tops of rib and flange intersections and cases of material flow under previously filled flanges

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subjects that have been considered in the case studies. Added factors that were evaluated in the case studies are:

Crankshaft underfill: induction coil inside diameter and stock diameter, equivalent current depth and subsequent time for conduction to reach a uniform stock temperature, total heating time for scale control, transfer time to press, and forging force applied Tube bending: precise heat input, control of temperature, and heated axial length of tube Spade bit: direction of forging relative to part shape and assessment of shear effect in extended wings Trim tear: trimmer tool tolerances, part temperature, and process time Upset forge: principal strains and equivalent plastic strain Flow through: strains going from round or flattened piece to finish, and assessing the need for a more generalized shape for input to finish die. Preform design—streamline

shape (not the same for aluminum-, magnesium-, steel-, titanium-, or nickel-base alloys) Lubrication: select one that provides the lowest coefficient of friction and other acceptable properties Forge process: total process for entire manufacturing train, including heat treatment and product testing Forge checking: fixture check for critical dimensions Forge tolerances: component to fit the customer’s assembly Simulation of process: verify that laws of plasticity are met

Forging Process Design Forging process design requires the application of integrated engineering principles that bring together factors such as:

Relationship between the important subsystem of a deformation system (Fig. 1)

Table 2 Factors in analysis of cold forging failures Preforge factors 1. Raw material—chemistry, microstructure, mechanical properties, size, surface finish, and cleanliness 2. Shape sequencing—general nature of shape to be created; strain, strain rate, and load requirements 3. Forging—equilibrium forging temperature, strain and strain rate, workpiece volume control, forge equipment, loading and transfer devices, lubrication, parts collection, inspection, and annealing 4. Trimming Causes of defects during cold forging 1. Cracking—Three factors combine to produce cracks: stress from thermal expansion and contraction, hydrogen, and a susceptible microstructure. 2. Product underfill—poor flow, sufficient volume, and proper distribution 3. Unbalanced forces—laps/lap fillin, nonhomogeneous strain, strain rate, nonuniform microstructure, and work hardening a. Seams—external and internal—on or within a metal surface, an unwelded fold or lap that appears as a crack usually resulting from a discontinuity b. Inclusions—raw material; internal and external substance that is foreign and insoluble to the matrix; particles of a foreign material in metallic matrix. Particles are usually compounds, such as oxides, sulfides, or silicates but may be of any substance that is foreign and insoluble to the matrix. c. Tears—occur when the equivalent plastic strain exceeds the capability of the material d. Entrapped scale—forged in contamination consisting primarily of oxides but can include other products left on metals 4. Strain hardening—increase in hardness and strength of metals caused by plastic deformation at temperatures below the recrystallization range; also known as work hardening 5. Flow through/push through—condition at which excessive material is provided in the preform in error, such that as elements of the shape are completely filled, such as flanges or rails, the central material continues to displace outward underneath the filled flanges 6. Porosity/voids—small openings, interstices, or channels within a consolidated solid mass or agglomerate usually larger than atomic or molecular dimensions 7. Segregation—In the casting process, the solidifying front moves away from the surface of the casting as a plane front, and lower-melting-point constituents in the solidifying alloy are driven toward the center. This is called normal segregation. 8. Internal shearing—This effect can occur when material displacements cause excessive sliding of adjacent volumes of material. 9. Surface impurities—any foreign substance deposited on the part unintentionally 10. Grain size structure—The number of grains per unit volume and the phase of the material dictate the forging response. 11. Flakes, blisters—These flaws typically result from the raw material or other processing steps but may show up when materials are forged. 12. Residual stresses/distortion—Most materials (especially steels) will have residual stresses after cold forging; distortion occurs when the stresses are not symmetrical. 13. Lubrication—dies and workpiece—viscosity and flow, hydrodynamics of lubrication, friction, heat generation and power losses, coefficients of friction

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Failure in Steel Forging / 135

Table 3 Factors in analysis of hot forging failures Preforge factors 1. 2. 3. 4.

Raw material—chemistry, microstructure, mechanical properties, size, surface finish, and cleanliness Shape sequencing—shape nature, temperature, strain and strain rate, upset tooling, fuller (roll), open-die tooling Hot forging—temperature, strain and strain rate, forge center cell, loading and transfer device, lubrication, parts collection, and inspection Trimming—trimmer unit and capacity, flash removal, temperature trace of product flashline

Causes of defects during hot forging 1. Cracking—occurs when the imposed equivalent plastic strain exceeds the material capability at the temperature of operation—surface (hot tears), cooling (centerline cracking) 2. Product underfill—underachieved thickness goal, inadequate material displacements, poor 3-D flow, inability of input shape to subsequent follow-on dies to satisfy local volume requirements, control of centroid path of newly created shapes 3. Unbalanced forces—laps/lap fillin, nonhomogeneous strain, strain rate, nonuniform and continuous microstructure a. Seams—external and internal—on or within a metal surface, an unwelded fold or lap that appears as a crack usually resulting from a discontinuity b. Inclusions—raw material; internal and external substance that is foreign and insoluble to the matrix; particles of a foreign material in metallic matrix. The particles are usually compounds, such as oxides, sulfides, or silicates but may be of any substance that is foreign and insoluble to the matrix. c. Hot tears—occur when the equivalent plastic strain exceeds the capability of the material at the temperature of operation d. Entrapped scale—forged-in contamination consisting primarily of oxides but can include other products left on metals 4. Flow through/push through—condition at which excessive material is provided in the preform in error, such that as elements of the shape are completely filled, such as flanges or rails, the central material continues to displace outward underneath the filled flanges 5. Porosity and voids—small openings, interstices, or channels within a consolidated solid mass or agglomerate usually larger than atomic or molecular dimensions 6. Segregation—In the casting process, the solidifying front moves away from the surface of the casting as a plane front, and lower-melting-point constituents in the solidifying alloy are driven toward the center. This is called normal segregation. 7. Internal shearing—This effect can occur when material displacements cause excessive sliding of adjacent volumes of material. 8. Surface impurities—any foreign substance deposited on the part 9. Grain size structure—The number of grains per unit volume and the phase of the material dictate the forging response. 10. Flakes/blisters—These flaws typically result from the raw material or other processing steps but may show up when materials are forged. 11. Residual stresses/distortion—Most steel forgings will have inherently residual stresses and distortion due to cold straightening or following quenching. 12. Lubrication—dies and workpiece—viscosity and flow, hydrodynamics of lubrication, friction, heat generation and power losses, coefficients of friction

Constitutive Equation

behavior to achieve stable deformation at a specified rate and proper evolution of microstructures and properties

Material system

Equipment

W

or

ol ntr Co stem Sy

ka bil ity

Forging Tolerances

Control system

Fig. 1

Relationship between important subsystems of a deformation system. Source: Ref 5

Interdependence of forging process parameters (Fig. 2) Forging process design task overview (Fig. 3) Relationship between process and machine variables (Fig. 4) Characteristics of forging machines (Table 4) Workability modeling (process maps showing zones of stable flow) of workpiece

The need for verification of the nominal dimensions and application of forging tolerances is important for quality assurance. Tolerances are required on forged products to allow for practical variations in die preparation, temperature effects during forging, equipment, and distortion during and after heat treatment. Forging tolerance review is a basic requirement to ensure that the part meets the multitude of design features and tolerances. A listing of the more important forging tolerances includes:

Dimensional—length, width, center-tocenter, and external-internal Die wear—generally approximately 0.102 mm (0.004 in.)/surface Die closure—thickness of approximately 0.813 to 6.35 mm (0.032 to 0.250 in.) as a function of plan form area at trimline

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Match—alignment of the top and bottom dies Radii—strong influence on material displacements

Flash extension Straightness—taken as a separate feature and then assessing its effect on the remainder of other tolerances

Data on Billet material

Ram velocity

Strain rate

Billet/Forging Geometry, Volume and thickness

Contact time under pressure

Flow stress/ forgeability

Temperature distribution in forging Die temperature, cooling Friction Conditions and coefficient

Interface lubrication

• Metal flow • Forging load • Forging energy

Fig. 2

Interdependence of forging process parameters

Fig. 3

Forging process design task overview

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Draft angle Datum plane location for three-plane (x, y, x) setup Alternative machined tooling points Finish allowance between forging and machined part

Tolerance review is conducted in various ways, and in the past, numerous forgings have been rejected and held for material review until some decision could be reached regarding their disposition for rejection or alternative repair. Even though this has typically been from review of the part drawings, another useful way to assess a completed forging is a fixture check. A uniquely designed fixture in conjunction with a dimensional inspector sets the forging into fixed tooling point locations and proceeds with the check “go” or “no-go,” determining whether the part will or will not serve its function in the assembly. In many cases, special fixtures and

Fig. 4

gages can confirm the accuracy of dimensions that are critical to the function of the component dimension. Large forgings are good candidates for fixture checking. Wear and Lubrication Surface interactions of two materials are influenced by small regions where contact is made at the atomic level. The real area of contact is determined by elastic and plastic deformation under consideration of loading. Lubrication reduces friction by introducing a viscous and low-shear-strength layer at junctions. Surface interactions can lead to wear, or the removal of material as a result of mechanical action. Wear types include:

Adhesion wear: particle transfers (pulled off) from one and adheres to the other Abrasive wear: a hard, rough surface plows grooves into the softer one

Forging equipment characteristics; relationship between process and machine variables

Table 4 Characteristics of forging presses Equipment type

Hydraulic press Mechanical press Screw press Hammer

Deformation rate

Slow Slow to medium Moderate to high High

Temperature loss

High Moderate Moderate Low

Consistency

Production rate

Very good Good Fair to good Fair to good

Low Moderate to high Moderate to high Moderate


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Corrosive wear: mechanical action removes a protective layer from a surface and exposes it to corrosive attack Surface fatigue wear: spalling occurs after the formation of surface or subsurface cracks Volume wear: proportional to the load and distance traveled and inversely proportional to the material hardness

During the early 1950s, the importance of proper lubrication was recognized on the shop floor. If an inexperienced oiler inadequately applied lubricant (in spray or paste form), then forging problems could occur even for an acceptable preliminary workpiece (preform). Alternatively, a questionable preform for the first closed impression die would be proven acceptable if an experienced oiler knew where the lubricant should be applied over the die impression and also when the die impression needed additional heavy lubricant in given die locations that appeared difficult to fill. These anecdotes vanished quickly as more science replaced art in forging. Presently, there are numerous ways that lubricants are used in the forging industry. Wrapping the workpiece during heating is an approach to prevent the formation of scale in the case of steel or thin metal sheets or cloths with impregnated graphite, in addition to the automatic spraying of lubricants. Lubricants play an important part in forging by minimizing the load required for maximizing material flow, protecting the die surface finish (critical for a specific lubricant), and assisting the entire forging process. Lubricant performance factors include:

Adequate lubricity Stability in gas-fired and electric furnaces Protect stock against atmospheric contaminants Provide good surface finish Act as a release agent No buildup in die cavity Ease of application and removal Conform to Environmental Protection Agency (EPA) and Occupational Safety and Health Administration (OSHA) requirements Acceptable cost Compatible with die materials

Graphite products for forging lubrication are:

GPC—for hot and warm forging Die lubricants—GP series

GP 100—low dilution ratios and spray application LS—oil and water Precoat workpiece—contains graphite as a lubricant pigment

Adhesion colloids are reliable for high pressure and temperature. Types include:

Colloidal—dispersions Delta forge lubricants—for hammer, press, and upsetters Deltaglaze—protective lubricants for billets applicable to steel

Case Studies Case Study 1: Crankshaft Underfill. There are several large steel forging components, such as ship crankshafts and airplane landing gear, being manufactured successfully in the United States and throughout the world today (2008). Crankshaft forgings in the weight range of 2268 to 4536 kg (5000 to 10,000 lb) are products made by a forging process creating a pair of flanges and a pinion shaft diameter at one time. The inboard and outboard flanges along with the pinion diameter become integral parts of the main shaft diameter. The forging operation creates one set of flanges by means of a working stroke in line with the major shaft diameter, while a 90 off-set load forges the pinion shaft between the flanges. These operations are generally performed following one local heating of the starting bar diameter for forging a set of crankshaft throws, including the two flanges and an offset pin diameter. The forging process is repeated until all of the flanges plus the offset pinion diameter are created along the major diameter of the crankshaft. The nature of the ready-for-assembly finish forging design for the incrementally forged crankshafts includes locations where material is provided for machining along with selected as-forged surfaces. During the forging of the flanges, there had been cases of small amounts of underfill at the flange extremities, as shown in Fig. 5. That extent of underfill has caused the entire component to be rejected. A test run was planned to measure material displacements while the flanges were being forged at the prior selected process variables of strain, strain rate, temperature of workpiece and dies, and forging force exerted. The conclusion

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Fig. 5

(a) Crankshaft flanges not filled. Main shaft diameter shown between flanges of adjacent “throws”. (b) Crankshaft flange with left side filled and right side not filled. Pinion shaft diameter located between flanges of single “throw”

reached was that since the workpiece material temperature at the end of the press working stroke was still in the hot working range, an extended length of time with the maximum force applied would be helpful to displace the relatively small amount of missing material into the remote regions of the flange dies. The thought was that allowing material to creep would aid in the final filling of the die cavities. Creep is an example of viscous flow and is defined as continuing flow at constant stress. At characteristic stresses, the creep strain reaches a steady state in which the rate of straining is constant. This is called the steady-state creep rate, e_ (or d e_ =dt). In hot working, the relationship between temperature (T), stress (s), and strain rate (e_ ) in the steady-state condition is best expressed as: 0

e_ =A(sinh as)n exp (7Q=RT)

where A, a, and n0 are temperature-independent constants, Q is the activation energy, R is the universal gas constant, and T is the temperature in Kelvin. At low stresses characteristic of creep (as5 0.8), this equation reduces to: 0

e_ =A0 sn exp (7Q=RT)

which describes the relationship among three variables under creep conditions. Evaluation of experimental data of the activation energy, Q, indicates that some metals in hot working soften the recovery process of repolygonization, and others soften by dynamic recrystallization. Thus, there is a distinct correspondence between hot working and viscous creep deformation.

Fig. 6

Both crankshaft flanges filled. Pinion shaft diameter located between flanges of single “throw”

In this case study, several time elements (all less than 60 s) were established in subsequent trials where all process variables were monitored (including a lower-than-press-capacity force), and flange fill results were measured. Finally, the proper combination of the important process variables, of which temperatures played an important part, enabled the consistent filling of the flange extremities, as shown in Fig. 6. Case Study 2: Tube Bending. Bending large-diameter tube (310 stainless steel) created a 90 bend in the diameter range of 635 to 762 mm (25 to 30 in.) and at a nominal wall thickness of 12.7 mm (0.50 in.) This offered a major challenge to manufacture. The challenge was to create a 90 angle bend without excess thinning at the outer wall and excess thickening at the inner wall. Earlier efforts had centered on using gas heaters around the circumference of the pipe

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and, in some cases, heating inside the tube. However, success was limited since the axial length of the heated zone was excessively long to permit control of the incremental bending strain and strain rate. As developments continued, a European company (Cojafax) had designed and manufactured a large bending machine to handle a 762 mm (30 in.) diameter stainless steel as the input and to impart a 90 bend with controlled thinning/thickening of the tube walls. An external axially thin induction coil supplied the heat and enabled control of the heated zone while a 90 bend was being made. The approach in the manufacturing process was to exert an axial force in line with the straight tube axis and then push the tube for a low strain rate through the induction coil until the heat input was sufficient to cause material displacements through the heated zone. In some cases, cooling rings were added to control the axial heated length, as shown in Fig. 7. This process was used for a number of trials in an attempt to achieve the goal of a 90 bend on a 762 mm (30 in.) diameter tube with 12.7 mm (0.50 in.) initial wall thickness.

Fig. 7

During these trials, process data were generated to show the variables that were used, such as machine force, bending moment, axial tube length inside the induction coil, tube axial speed, and amount of heat supplied, all of which were used to control the shift of the tube neutral axis and thus the thinning and thickening of the tube walls. Optimizing the process variables for a 90 bend on the 762 mm (30 in.) stainless steel tube resulted in minimal thinning of the outer wall, as shown in Fig. 8. The thinning of the outer wall of several tubes met the initial program goal of 18%. Case Study 3: Spade Bit. The cutting end of a proposed wood boring spade bit (AISI 1000 series) consists of a central web connected to angular extensions from each side of the web, as shown in Fig. 9. The forging process for the spade bit forge was designed so that the finish shape could be cold forged in a continuous line, starting out with wire, straighten, clean, lubricate, room-temperature forge, trip flash, and followed by heat treatment. The inherent tooling design presented an opportunity to run a series of numerical experiments using computer

Closeup of tube bending assembly. (a) Induction coil and water ring. (b) Induction coil and partially heated tube and water ring

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Bent product tube with 90 bend and minimal thinning at outer tube wall of ~18% of starting tube wall thickness

Fig. 8

Fig. 9

(a) Spade bit drawing. (b) Photograph of each half

simulations varying the forging direction, so that the direct compressive force applied by the dies to the central web of the workpiece would also then have components of shear acting to cause material displacements in the extensions. Several die rotations were attempted to create the longest wing extensions with a minimal central web and the lowest forging force. The fundamental idea behind this approach was that steels in shear are weaker than steels in direct compression. Strain-limiting criteria for cold forging low-alloy steels are shown in Fig. 10. Initially, analyses were made to determine the strains to be encountered when forging a round bar to a flat central section with two attached wings off at different angles. Shear stresses were

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determined by Mohr’s three-dimensional (3-D) analysis to quantify the stresses and strains throughout the central web and extended wings. Forming simulation software Antares (from UES-Software, now defunct) was used for the deformation analysis, followed by a shop trial. Strains in the wings as well as in the center section were calculated along with the compressive force and stress on the center section, before/after which simulations verified the analysis. Following a few shop trials supported by simulations, an optimized process was determined that led to the minimum amount of energy to be used to forge the spade bit center and wings. Case Study 4: Trim Tear. During the forging sequence of a typical mechanic tool product (AISI 4000 series), the process was running satisfactorily, except that tearing occurred at the flash trimline and then propagated into the forging proper. Trimming of the flash around the perimeter of a finish forging has traditionally been a very dependable operation by maintaining the proper clearance between the punch and the trim tool blades, so that the deformation zone between the punch, forging with flash, and the trimmer sheared the flash with no bending of the flash extension.

Fig. 10

Strain-limiting criteria for cold forging low-alloy steels

Excessive clearance encourages material bending displacements, which lead to shearing in the zone. Figure 11 depicts a schematic of the effect of punch-to-die clearance on characteristics of edges of holes produced by piercing a low-carbon steel. In any event, the flash shearing is not clean and adjacent to the draft wall of the forging. In some cases, the shear mechanism initiates a crack on the edge of the forging and flash. The case that is being reported here is one where cracks initiated and propagated into the forging proper. Micrographs of crack formation at the flash edge are shown in Fig. 12. During the search for a single or multiple solutions, several conventional avenues of attack were followed:

Trimmer tool setup for proper clearance at the trimline, seating of the forging inside the trip plates, and proper contact of the workpiece in the trim plate nest should be assured. Raw material condition did show some evidence of banding and inclusions. Temperature in the workpiece may be too high due to high speed of production and resulting lower strength at trim. Measured clearance between punch and die is reduced.

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The final analysis concluded that the temperature of the trimming operation (immediately following forging) was excessive and triggered the initiation of a small crack. During further trials in the shop, the time between forge and trim was increased. This made the workpiece have a higher strength because of the lower temperature, and no cracking occurred even though there was evidence of material banding. Other changes were made, such as decreasing the clearance between the punch and die from 0.127 to 0.076 mm (0.005 to 0.003 in.), which made the trimming operation cleaner and more robust. Case Study 5: Upset Forging. One of the most important operations in the forging process is the upsetting of a billet of material with ratios of axial length to diameter of less than 3 to 1 and,

Edge characteristic

Fracture angle Rollover(a) Burnish(a) Fracture Burr

preferably in some cases, 2.5 to 1 using flat dies. The material displacements are primarily radial, extending out from the billet center and forming the outside diameter. Material radial displacements come from the decreasing axial length of the starting billet. Material displacements during upsetting have been well analyzed in the past by numerous investigators who have identified the steel alloy; axial, radial, and tangential strain; strain rate; temperature of the workpiece and die; along with the frictional effect at the workpiece-die interface. During an extended working stroke, there is a period of time where the entire reduction operation is considered as one of nonsteady state. During this operation, there is one plasticity law that states that the sum of the principal strain at any time is equal to 0. This has proven to be a significant benchmark to

Type 1

Type 2

Type 3

Type 4

Type 5

14–16 10–20% t 10–20% t(b) 70–80% t Large, tensile plus part distortion

8–11 8–10% t 15–25% t 60–75% t Normal, tensile only

7–11 6–8% t 25–40% t 50–60% t Normal, tensile only

6–11 4–7% t 35–55% t(c) 35–50% t(e) Medium, tensile plus compressive(g)

... 2–5% t 50–70% t(d) 25–45% t(f) Large, tensile plus compressive(g)

(a) Rollover plus burnish approximately equals punch penetration before fracture. (b) Burnish on edge of slug or blank may be small and irregular or even absent. (c) With spotty secondary shear. (d) In two separate portions, alternating with fracture. (e) With rough surface. (f) In two separate portions, alternating with burnish. (g) Amount of compressive burr depends on die sharpness.

Fig. 11

Piercing of low-carbon steels. Source: Ref 6

Fig. 12

Micrographs of crack formation at flash edge


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evaluate the effects of the lubrication at the diematerial interface. Many forged circular components and also circular components with off-center port bosses begin as a billet of material of a given diameter and are then forged between flat dies to a larger diameter and a reduced axial dimension. Subsequent forging operations on the forged initial preform create cylindrical walls along with various configurations, depending on the final forging design. Thus, an initial flattening (or sometimes referred to as “pancaking”) between flat dies is a common initial forge operation for a number of steel components, for example, missile cases and bowl shapes. A typical steel upset forged between flat dies made on a screw press is shown in Fig. 13. However, during the upsetting

Fig. 13

Steel upset forged between flat dies made on a screw press

operation, cracking on the barrel of the upset piece often occurs because of higher strain levels than the material can sustain and, in particular, higher effective strain levels when each of the principal strains—axial, tangential, and radial— are considered. The fix for a problem of this nature is to calculate the strain field in advance when plans are made to create the upset piece. A numerical or computer simulation of the process can provide a “go” or “no-go” on the planned process. In the past, there have been cases reported of increased upset temperature within the workpiece, and the solution to that problem is a reduction in the strain field. Most of the cases fall into the situation of an unanticipated strain-rate effect, requiring the rate of forging to be reduced. Table 5 (with Fig. 14) shows some selected results obtained from an MSC superforge simulation of forging a flattened disc (pancake) between two flat dies (Fig. 15). The analyses are based on plasticity laws. The damage variable is defined as the ratio of the total cavity area over the total area found in a representative volume element. Therefore, the damage variable is a dimensionless quantity between 0 and 1, where D=0 describes the undamaged representative volume element, and D=1 is the failure due to rupture. A critical value of damage for multiaxial states of stresses may be defined as a quantity that describes the occurrences of measurable cracks in the material. When the critical value of damage reaches a certain magnitude, one can

Table 5 Analysis and simulation of upset forging a disc between flat dies Analytical analysis

Numerical simulation (superforge)

Starting billet size: 101.6 mm (4 in.) diameter and 101.6 mm (4 in.) length Ending flattened disc size: 198.2 mm (7.8 in.) diameter and 26.67 mm (1.06 in.) length (axial thickness) Strain: Axial?ln [1.05/4.00]=1.337 Radial+tangential?2 (ln[7.8/4.0])=2 (0.668)=+1.336

Equivalent plastic strain: 1.389 1.389

Analysis: axial strain (1.337)+radial strain (+0.668)+tangential strain (+0.668) ~0 Plasticity law: The algebraic sum of the principal strains equals 0. Z stress average (disc center-to-edge readings): determined by maximum force divided by disc (after working stroke) plan form area (PVA)

1.894 · 108 Pa (27,469 psi) 1.575 · 108 Pa (22,843 psi) 1.240 · 108 Pa (17,984 psi)

Z stress=Forging force /PVA=1.081 · 106/47.78=1.56 · 108 Pa (22,624 psi)

(Disc thickness center-to-barrel-edge readings)

Ratio of the average stress exerted during the flattening operation to the yield strength is 19,809 psi/16,215 psi=1.16, which is in the range of Z stresses reported.

1.2554 · 108 Pa (18,207 psi) 1.211 · 108 Pa (17,563 psi) 1.225 · 108 Pa (17,766 psi)

Average of all Z stresses taken is 20,305 psi. This ratio agrees closely with the graph of the average stress/yield strength when plotting the dimensionless parameter of coefficient of friction disc · diameter/part thickness when forging a flat disc (Fig. 14)

(Disc center-to-barrel-edge readings) 1.162 · 108 Pa (16,853 psi) 8.247 · 108 Pa (11,961 psi) 4.466 · 108 Pa (6,477 psi) Yield strength: 1.118 · 108 Pa (16,215 psi)

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Fig. 14

Pressure multiplication factors for forging of thin panels. Source: Ref 7

Fig. 15

MSC superforge simulation. Disc upset forged between flat dies, showing (a) start position and (b) end position after 74.93 mm (2.95 in.) stroke

deduce that the material is irreversibly damaged. A typical result plot is shown in Fig. 16. Case Study 6: Avoidance of Flow Through, Lap, and Crack. During the forging of an “H”shaped cylinder of a flat web with projecting ribs and flanges, material displacements are required to turn 90 in the direction from the web within

the die cavity to fill the external/internal ribs and flanges. There have been numerous cases in the past where insufficient thickness of the material front moving horizontally caused the front to contact the die wall and upset on itself, enabling the filling of the outer flange. Following the filling of an outside flange, as shown in Fig. 17, material displacements continue to move outward toward the flash opening and underneath the filled flange. This combination of material displacements causes a flow through at the base of the flange. Another common case of underfill coupled with lap formation in many forging designs occurs at the intersection of outer flanges and cross ribs, where material displacements are primarily 3-D. Additional material is required to fill the top of the flange and rib intersection because of the volume required. That additional volume is provided by a preform shape with increased fillet radii or taper at the base of the web. The term lap describes a defect that forms whenever material folds over itself during the forging of a new shape, using a previously designed preform as the input to follow on the set of dies. Laps occur when both vertical and

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Fig. 16

Typical state of strain in hot upset forging of steel showing fracture criteria in MSC superforge simulation

Fig. 17

Typical forging defect caused by excessive natural flow through the forging of a rib (flange)-and-web part. Flow through is the tendency of a metal to flow naturally past the rib (flange) opening.

horizontal sections intersect. When this occurs, it is an indication that a preliminary shape is required as input to the next die in the forging sequence to provide material to fill intersecting elements. Also, these types of problems are analyzed by a technique referred to as

streamlining. The series of shapes are tracked backward from the finish shape to blocker and preblocker shapes to the cogged, rolled, or upset piece. A series of “part-way” downs were forged using aluminum alloy 7075 on a hydraulic press at the temperature of 399 C (750 F) in a set of preform (blocker) dies that had been designed to show the material displacement field when attempting to fill a flange from a web surface. The samples were approximately 101.6 mm (4.0 in.) in diameter and 203.3 mm (8.0 in.) long to minimize the end effect. The strain rate during the trials was selected so that the displacement field could also be applied to other alloy systems, such as carbon steels, so that the same level of forging pressure would be used with similar results. The results are shown in Fig. 18. In the fifth and sixth “part-way” down, there is evidence that the material front is moving away from the entrance to the flange, and die contact is lost until an additional stroke is applied. Then, the outer die design adjacent to the flash gutter creates sufficient back pressure to displace material into the flange. Thus, a shape is made that will be suitable as an input to the finish dies,

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meeting the web thickness tolerances and filling the flange. The process is to reverse-integrate from the finish design to determine a more generalized shape with precise volume distribution along the

Fig. 18

three principal axes. Preforms in the series must satisfy subsequent shapes of the finish product design and local volume requirements, providing die materials for intersecting product features and other geometric attributes—round

Creation of a streamlined preform serving as an input to a finish die rib (flange) and web design, avoiding flow through

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(simple and compound), elliptical, and tapered shapes must all be accommodated. Admissible criteria and specific tasks for preform designs include:

Specific alloy characteristics Microstructure requirements of a finished product in terms of percent reduction at each forging operation Match areas and volumes along principal axes and location of centroids Perform reverse-integration, streamlining the finish shape to more generalized features—lower ribs and rails (flanges) coupled with increased web thickness and connecting radii Calculate principal strains when comparing finish cross sections to preform cross sections Generate an overall preform shape to obtain uniform deformation in the finish die Examine the nature of material displacements over the die contour for unsupported material fronts as the working stroke progresses Examine preform locations in the finish (or subsequent) die and initial die contacts Examine unsupported webs at die contact to prevent buckling Make short plots of the displacement field in terms of the material contacting fixed die boundaries and the change of shape of the material front being generated as the deformation progresses Follow the continuous trace of the displacement field in terms of the material contacting fixed die boundaries and the change of the shape of the materials front being generated as the deformation progresses Determine the amount of energy expended for each preform evaluated and then the entire shape sequence

REFERENCES

1. G.E. Dieter, H.A. Kuhn, and S.L. Semiatin, Ed., Handbook of Workability and Process Design, ASM International, 2003 2. Metalworking: Bulk Forming, Vol 14A, ASM Handbook, ASM International, 2005 3. T. Altan, G. Ngaile, and G. Shen, Cold and Hot Forging: Fundamentals and Applications, ASM International, 2005

4. J.E. Johnson, Ed., Forging Industry Handbook, Forging Industry Association, Cleveland, OH, 1966 5. H. Gegel, G. Huang, and S. Manna, “Precision Forging—Quality—Productivity— Equipment—A Technical Article,” UES Software Inc., Dayton, OH 6. Piercing of Low-Carbon Steel, Metalworking: Sheet Forming, Vol 14B, ASM Handbook, ASM International, 2006, p 159 7. M.D. Stone, The Design and Construction of Large Forging and Extrusion Presses for Light Metals, United Engineering and Foundry, Pittsburg, PA

SELECTED REFERENCES

J. Burke and V. Weiss, Advances in Deformation Processing, Sagamore Army Materials Research, Vol 21, Army Materials and Mechanics Research, Massachusetts and Syracuse University, NY H. Chandler, Metallurgy for the Non Metallurgist, ASM International, 1998 G.E. Dieter, Mechanical Metallurgy, 3rd ed., McGraw-Hill Book Co., 1986 D.D. Fuller, Theory and Practice of Lubrication for Engineers, John Wiley and Sons, Inc., Chapman and Hall, Ltd., 1956 C.G. Johnson, Forging Practice, American Technical Society Publisher, Chicago, IL, 1954 S. Kalpakjian, Manufacturing Processes of Engineering Materials, 3rd ed., AddisonWesley, 1977 A. Kannappan, “Wear in Forging Dies—A Technical Paper,” Swedish Institute of Product Engineering Research, Goteborg, Sweden, 1969 C. Lipson, “Wear—Consideration in Design, Residual Stresses and Contact Stresses—A Technical Paper,” University of Michigan, Ann Arbor, MI F.A. McClintock and A.S. Argon, Ed., An Introduction to the Mechanical Behavior of Metals, School of Engineering, Massachusetts Institute of Technology, Cambridge, MA, 1962 “Research Report: Work at IIT Research Institute,” Committee of Hot Rolled and Cold Finished Bar Products, American Iron and Steel Institute, New York A.M. Sabroff, F.W. Boulger, and H.J. Henning, Forging Materials and Practices,

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Battelle Memorial Institute, Columbus, OH, Reinhold Book Company, 1968 J.A. Schey, Introduction to Manufacturing Processes, McGraw-Hill Book Company, New York, 1977 J.A. Schey, Ed., Metal Deformation Process, Marcel Dekker Inc., 1970 J.A. Schey and P.W. Wallace, Research Report: Metal Flow in Closed Die Forging of Steel, Part 2: Speed and Lubrication Effects, American Iron and Steel Institute, New York, 1966 J.A. Schey, P.W. Wallace, and F.A. Shunk, Research Report: Metal Flow in Closed Die Forging of Steel, Part 1: Fundamental

Study, American Iron and Steel Institute, New York, 1966 T.M. Silva and T.A. Dear, “Wear in Drop Forging Dies—A Technical Paper,” Department of Mechanical Engineering, University of Birmingham, 1969 J.W. Spretnak, “Technical Notes on Forging,” Forging Industry Education and Research Foundation, Cleveland, OH, 1976 “Technical Notes: Mechanical and Physical Properties of Ferrous Forging,” Committee of Hot Rolled and Cold Finished Bar Products, American Iron and Steel Institute, New York

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Failures from the Casting Process Omar Maluf and Luciana Sgarbi Rossino, Instituto de Materiais Tecnolo´gicos do Brasil Ltda. Camilo Bento Carletti, Celso Roberto Ribeiro, Clever Ricardo Chinaglia, and Jose´ Eduardo May, Universidade Federal de Saõ Carlos

THE HEAT TREATMENT of a steel component is often the last step or near the end of a somewhat complex manufacturing process. Finished products require attention to each step of the long operation chain from raw material to finished product. Early in-service failures of components after heat treatment may result from improper planning, lack of required equipment, nonqualified personnel, not enough time to execute the expected operations, or even a combination of some or all of these deficiencies (Ref 1). However, most of the early failures that happen during the heat treatment process are the result of features generated in previous manufacturing stages. A component lifetime basically depends on the following factors:

Global component project Materials selection Material quality Processing methods, such as casting and machining operations prior to the heat treatment Heat treatment Final finishing operations Mechanical solicitation of the component and the service environment

However, it is not a simple task to identify which of these items is responsible for the early failure of a component during heat treatment or in service. The failure analyst uses these seven items as a guide in a failure analysis. This chapter deals specifically with improper casting projects and those features that originated in the casting process itself, including porosity (generated by the presence of gas as well as by shrinkage pores), decarburization, cold joint, and inclusions. These features may not be called defects because, according to ASTM

E1316-2005, “Standard Terminology for Nondestructive Testing,” components have defects only when they fail to meet their specification requirements. If a component has a large amount of porosity, for example, it is not a defect unless (1) an inspection porosity is specified, (2) its amount exceeds the required acceptance criterion, or (3) the component fails because of this porosity. This chapter describes cast steel features that may be identified or attributed to component failure during heat treatment or subsequent processing or service. As such, these casting features are referred to as defects in this chapter.

Failures due to Improper Cast Design The engineer’s designing job may face inevitable weak points due to some inherent component characteristic in use. The engineer should, however, try to overcome these inconveniences by looking for alternative solutions and finding the middle ground between the component functionality and the manufacturing difficulties. The project aspects that should be avoided at all costs in the cast component production are (Ref 2–7)

Sharp edges, sharp corners, and nonround edges Abrupt section changes Holes, especially when located near the external wall of the component Sections with cross connections Unfavorable length/width relationship

Rounding the corners, as shown in Fig. 1, should always be performed in order to avoid the stress concentration that can originate from

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cracks formed during casting solidification in the mold, during heat treatment, especially quenching and tempering, or even during heating for austenitization (Ref 2). The risks will be reduced if steel with increased temperability is chosen, which requires a less severe quenching medium, such as oil or air. Other strategies in the design of cast tooling are to avoid creating components with right angles or to machine the corners to make them round. Another option is to quench and temper the component and then remove the exceeding material to give the component sharp corners if they are required for its function. This last strategy requires a steel with good temperability. Otherwise, when the exceeding material is removed, that region in the component will present a surface with lower hardness and less resistance to abrasion than the previous one. The quenching of components with abrupt section variations in a liquid environment always represents a serious problem due to the associated stress concentrations, even if the transitions are made using the apparently correct concordance radius resource (Ref 4, 7). In this case, the solution is to create the component in different parts, treat them separately, and assemble later on. However, if the component must be made as one unit due to a functional imposition, the solution is to choose an airquenchable steel that presents a lower crack probability. The existence of holes raises a problem mainly in high-carbon steels and/or alloying elements. The abrupt section variation and other specific aspects of the holes (Fig. 2) must be considered. The accumulation of quenching liquid in the interior of blind holes leads to an improper heat loss of the internal walls, thus lowering the hardness. In open-ended holes, the heat removal may not be as effective as in the rest of the component, which is more exposed to the quenching liquid (Ref 5). Therefore, when the chosen steel is quenched in a liquid medium,

the components containing holes must be quenched in specially designed devices so that they receive strong gushes of liquid in the interior (of the holes), or they must be arranged so that all of the set is subject to a strong stirring. In cases where the holes do not need to be hardened, they can be made of low-alloy lowcarbon steel components already inserted during molding for the casting, for example, in tool steel components. Another possibility during heat treatment is to fill the holes from casting with any material that can totally inhibit contact with the quenching fluid that would result in hardening of this region. Regarding stress concentration, it is preferable to have the existence of a completely quenched hole than the presence of a mixed structure (hardened and soft). Both methods, particularly the first, make the achievement and finishing of holes and threads easier. Cotter holes, especially the rectangular section ones, are places with high stress concentration (Ref 2). Therefore, whenever possible, they should be eliminated or substituted by channel sections, whose locking efficiency is equivalent, but the stressconcentration factor is three to four times lower. Another geometry that should be avoided during project design of the cast tooling component is the cross type, such as the furnace grid and heating equipment shown in Fig. 3(a). It causes serious crack problems in the cross area during the solidification process while still inside the sand mold or during heat treatment (Ref 3). The solution is to use node dislocation, as illustrated in Fig. 3(b). Lengthy components with very thin sections or small diameters show serious bending problems during heat treatments, mainly quenching, even when the steel is favorable to less severe environments. The problems start from the moment the component is put into the furnace for austenitization. If it is secured only in two extremities, there is the risk of deflection

Fig. 1

Fig. 2

Sharp edge elimination. (a) Sharp corners create high strain concentration. (b) Exaggerated relief causes a shrinkage cavity. (c) Ideal relief

Types of holes. (a) Blind hole with a parallel bottom. (b) Blind hole with a steeple bottom. (c) Passing hole— the most economic

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Failures from the Casting Process / 153

Fig. 3

Grid crossings. (a) Crossed node region of crack formation. (b) With dislocated nodes, the occurrence of cracks is less likely if the distance between nodes, d, is larger than 2r+e, where e is the thickness, and r is the curvature radius.

due to its own weight. If it is supported on the furnace hearth, its heating will not be homogeneous, making the component subject to bending and/or to the appearance of soft spots (Ref 6). The correct heating method, in this case, consists of hanging the component by one of its extremities and using a furnace that allows the austenitization of the component hung in the vertical position. Also to be avoided is the manufacturing of too thin, lengthy components. They should be split into components whose length-height ratio is more favorable (Ref 3). For example, highalloy, high-carbon, steel sheets used in the guillotine have been replaced by shorter sheets that, after quenching and tempering are assembled in a chassis to present a continuous edge. An advantage of this solution is that only the damaged component of the edge can be replaced when there is an in-service failure of one of the sections, leading to an easier and cheaper operation. Very big components with a circular section larger than 25 cm (10 in.) in diameter, or rectangular with equivalent mass, also present problems during quenching. When carried out in a liquid medium, the surface reaches the starting martensitic transformation temperature long before the central region, with the generation of stress that may cause internal cracks as a consequence. It is recommended that big, bulky components be replaced by sets of smaller

Fig. 4

Crack resulting from the normalization heat treatment of an AISI 1045 steel cast hull caused by thinning of the wall due to deficiency of the tooling or the core alignment

pieces, despite the adjustment problems that result from this operation. Figures 4 to 6 illustrate failures that happened during the heat treatment operation due to poor design considerations. Figure 4 is a crack that occurred during the normalization heat treatment of an AISI 1045 steel cast hull. A prior thinning of the wall (due to deficiency of the tooling or core alignment) promoted the cracking during normalization. Figure 5 is a crack that happened in the normalization heat treatment, caused by stress buildup in the sharp edge region. In Figure 6, poor design of an edge led to cracking after quenching.

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Effects due to Porosity During the production of casting components, the interactions between the liquid metal and elements from the gaseous atmosphere, furnace refractory, foundry ladle, molds, and core materials are important factors from a technological and metallurgical point of view and are responsible for desirable or undesirable changes in the chemical, physical, and mechanical properties of the metallic materials. The dimensional precision grade of the feeding system should also be taken into consideration. Therefore, the quality of the casting product is related to the physical integrity of the

component, that is, the absence or minimization of the quantity of defects present. Among these defects, the most important ones are those generated by the interaction of gas and metal that promote the appearance of voids. In general, there are two kinds of voids: those generated by gas, and shrinkage pores. Porosity Caused by Gas One of the factors that must be considered in steel casting is the behavior of the gases in the process. Generally, there are three major sources that may contribute to porosity formation (voids caused by gases) in steel castings. These are:

Fig. 5

Crack in the bottom of a machine molded from AISI 1030 steel that happened in the normalization heat treatment, caused by stress buildup in the sharp edge region

Fig. 6 treatment

Plastic injection mold casting in AISI H13 steel with a crack from a sharp edge after a quenching heat

High initial gas content of the melt originating from the charge ingredients, melting practice, or atmospheric humidity Reaction of carbon and dissolved oxygen under certain melt conditions Mold-metal reactions between the evolved mold and core gases at the solidifying casting surface

In addition, any combination of these three sources may have an accumulative effect in promoting porosity formation. However, the gases normally held responsible for subsurface porosity defects are nitrogen and hydrogen. Types of Gas Porosity Defects. Pinholes and blowholes are the two main kinds of porosity caused by the presence of gas (Ref 8). Gas porosity (pinholes) refers to hydrogen, oxygen, and nitrogen gases within a casting. Molten metal has such an affinity for H2, O2, and N2 that it will disassociate it from other molecules, such as water or atmosphere gases, and form a solution with it. As with most solutions, as the temperature drops, these gases become less soluble and precipitate as gas. The greater the amount of gas in the molten metal and the slower it solidifies, the greater the gas voids. It should be remembered that the H2 comes from the mold humidity, when the H2 from the metal is eliminated in the cleaning process performed before pouring. These voids are generally smooth, round, or slightly elongated and may be somewhat localized in the areas of the casting that solidify last. This type of porosity is generally undetectable visually, since the surface of the casting solidifies the quickest, preventing the gases from forming holes large enough to be visible on the surface, except through

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fluorescent-penetrant inspection or crack detection during or after the heat treatment. Gas holes (blowholes) are generally larger and more localized voids than gas porosity, but they retain the smooth, round, or slightly elongated shape. They are usually caused by a reaction in the mold medium, producing gas that bubbles through the molten metal. The humidity contained in the mold walls and cores is the main source of the vapor that is necessary for defect (gas bubble) formation in the casting component. During mold filling, the gas generated by the metal-mold reaction is eliminated to the environment through permeability, a hole from the exit of gases, and/or a rising gate. The residual quantity of gas that could lead to bubbles is almost nonexistent or negligible. The exception would be the use of low-permeable molds, for example, the ones whose sand contains a high percentage of fines, making the passage of the gases to the environment difficult. For the cores, which may become completely surrounded by liquid metal during mold filling, the problem can be more serious. Gas elimination to the exterior, including the gases generated by the binder and collapsible materials, is extremely difficult. It may require the use of devices such as internal wax wicks in all the core extensions, so the gases are “sucked” toward the core prints and then eliminated. A gas bubble can also occur, even though it is not very common, as the result of an inadequate measurement of the descent channels, distribution, and attack, which, during pouring, can cause turbulence in the liquid metal flow or can cause air to be inhaled to the interior of the mold cavity, where it mixes with the liquid metal. Behavior of Gases. Dissolved hydrogen and nitrogen in the molten steel can cause a porosity defect such as a pinhole. The extent of gas porosity depends on the amount of these gases, the alloy, chemical kinetics, and the alloy surface tension (Ref 9–12). During solidification of most steel alloys, the component that is still liquid becomes more concentrated in alloy elements due to its solubility. This solubility difference is expressed by the component ratio KS/L, which is a relation between the quantity of solute present in the solid and in the liquid. For most alloys, this value is usually lower than 1, which indicates a liquid enrichment during solidification. The hydrogen solubility in the austenite is nearly 7 ppm, meaning that its solubility in the molten and solid conditions is approximately the

same, and there is low hydrogen segregation during the solidification, which reinforces the fact that the presence of gas bubbles caused by hydrogen has other causes, for example, the reaction with the moisture from the mold and/or the cores (Ref 9). For nitrogen, its component ratio for stable and metastable eutectic solidification is 1.9 and 2.2, respectively (Ref 13). This shows that nitrogen is less soluble in the liquid metal than in the solid metal. However, it is good to highlight that in a real situation, the solidification involves the liquid-solid diffusion and vice versa of a larger quantity of elements, nitrogen being just one of them. When nitrogen behavior in iron alloys is analyzed, the most important element whose mutual presence must be considered is carbon. For example, in a 3.8% alloy at 1500 C (2730 F), the nitrogen solubility at equilibrium in the liquid metal is 110 ppm. At the time of eutectic solidification, this value is reduced in the liquid to 97.5 ppm due to austenite enrichment. In this sequence, because there is an increase in the carbon percentage in the liquid, the nitrogen solubility is reduced to 90 ppm. There is nitrogen saturation in the liquid, and thus, this excess will cause the evolution of the gas and may originate pinholes, as seen in Fig. 7. When the possibility of pinhole formation is evaluated, the component pressure of all involved gases must be considered. When this sum is higher than 1 atm, pinholes form. So, in the previous example, the formation of pinholes could happen with lower nitrogen concentrations than those mentioned, needing just the presence of other gases, such as hydrogen or oxygen, for example. It is also important

Fig. 7

Typical morphology of a defect called a pinhole, caused by gases

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to consider the dynamic formation of these bubbles in pinhole formation. However, the thermodynamics of this event are very complex and are not covered in this chapter. How to Treat Pinhole and Blowhole Problems. In order to minimize or eliminate the pinhole formation problem, a relatively simple method can be used: bubbling argon in the desulfurization reactor or ladle. When the argon is blown into the bath, the gases in the atomic form combine themselves on the bubble surface, forming molecules from the respective gases (N2 and H2). However, it is known that bubbling is more effective in hydrogen elimination than in nitrogen. One of the most common alternatives to reduce or eliminate blowholes is to increase the pouring temperature. There will be a higher fluidity and time interval for the beginning of component solidification, giving time for the gases to escape to the atmosphere. However, care must be exercised in the decision to increase the pouring temperature. The concentration will also be bigger, and there is the risk of the riser becoming undermeasured, thus allowing the occurrence of a shrinkage cavity, in addition to the possibility of a molding and core system collapse, causing other defects in the casting. In summary, to minimize the effect of void appearance from gases, the main measures to be taken are:

Control of the furnace atmosphere using vacuum or gases with low solubility values Develop a project of feeding channels to avoid turbulence Use sand molds and cores with the lowest humidity and the maximum permeability possible Use low-solubility gases that, when injected in the liquid metal, carry the dissolved gases to be eliminated to the surface

system, directional solidification, and pouring temperature. During the liquid-to-solid transformation, there is a grouping of atoms that forms ordered structures. In the majority of cases, this transformation is followed by a density increase (Fig. 8) and thus a shrinkage, because the metal as a liquid occupies a larger volume than in the solid state. The defect known as shrinkage pores can be characterized as the appearance of nonsuperficial cavities in the casting component due to the lack of predetermined and precise compensation devices for the liquid metal shrinkage that occurs during solidification and/or the metallic inserts for directional freezing. If the concentration is a little higher than the capacity of the system to compensate for it, or if thicker pieces of the component work as risers for thinner components, small, irregular voids will be formed, as is seen in Fig. 9. However, if the concentration is much higher than the compensating mechanisms, there are large voids of irregular shapes on the surface of the component. These are called primary shrinkage cavities or simply shrinkage cavities, as seen in Fig. 10, and are not discussed further in this chapter. In order to better understand shrinkage, an example is given of the production of a cast iron component with two kinds of molds: a nonstiff mold with synthetic sand (green) and a stiff mold with phenolic no-bake sand. In the one produced with synthetic sand, there is a factor that must be considered: the mold walls deform, increasing the volume of the cavity when it receives the liquid metal, and thus, it requires more metal to

Porosity Caused by Shrinkage Pores This type of porosity has a rough, irregular shape. It is caused by a lack of adequate feed metal during solidification. This defect is an internal void known as shrinkage pores. It usually is detected only through ultrasonic or radiographic tests or, during heat treatment, when it causes disruption in the components. Shrinkage pores are not related to the high or low presence of any kind of gases but to the feeding

Fig. 8

Density variation with temperature in metals

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compensate the increase in the volume. With the rigid phenolic resin mold, there is no volume increase in the cavity, and the additional liquid metal is not necessary. On the contrary, when the

Fig. 9

Example of a shrinkage pore

(a)

equivalent carbon is larger than 3.9%, there is a graphitic expansion, which is larger than the solidification shrinkage and could cause a metal reflux for the mold exterior. Each metal or metal alloy presents a characteristic concentration rate during the solidification process. Therefore, it is possible to estimate quite precisely which feeding condition will be necessary to avoid the occurrence of porosity and shrinkage cavity problems. The theoretical calculations to predict metal volume shrinkage during the casting process are based on a model proposed by Campbell (Ref 14), using a sphere as an example. There are basically three different types of shrinkage that may occur during the solidification process, as shown in Fig. 11: liquid shrinkage, solidification shrinkage, and solid shrinkage. During the casting process, the first type of shrinkage observed is the liquid one, which happens with temperature decrease. However, this does not represent significant problems in the quality of a casting component when the volume reduction occurs linearly with temperature decrease, and the necessary volume of liquid material to compensate this volume reduction can be given by risers. On the other hand, the volume shrinkage that occurs during solidification of the liquid metal can bring more serious problems to the casting component, and thus, it requires more attention. The major concern is to make the feeding process, which replaces the liquid metal necessary to compensate the shrinkage in the system, very precise in a way that allows for the attainment of perfect components. This shrinkage compensation process determines the precision and perfection of the casting component and is inversely

(b)

Fig. 10

Primary shrinkage cavity forming large voids of irregular shapes on the component surface. (a) Schematic drawing. (b) Shrinkage cavity compensated for riser

Fig. 11

Schematic representation of the three regimens of shrinkage: in the liquid state, during solidification, and in the solid state. Source: Ref 14

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proportional to the quantity of shrinkage cavities or shrinkage pores present in the obtained component. During the shrinkage process that occurs in the solid state, the component size starts to be reduced. At this moment, the casting component faces the resistance of the mold and/or core. This kind of stress from the casting component, when trying to contract, generates residual stresses that may cause plastic deformation of the casting component, hot tearing, or cracks during heat treatment later on. Yet, this shrinkage depends more on the volume reduction intrinsic to the cast alloy and the project of the mold than on the casting parameters. Six Rules for Casting Component Feeding. In the absence of gases and if the feeding of liquid metal is appropriate, no porosity will be found in the casting. However, because there are many complex casting projects, there may be regions of the mold with feeding problems, allowing the internal hydrostatic tension in the liquid metal to generate the conditions for the formation of internal pores. In the design of a component to be cast, it is necessary to have an effective supply of material in order to compensate the shrinkages previously mentioned. For the additional liquid metal supplied to the system to compensate the volume shrinkage that occurs during cooling, a riser must be provided in the casting of the component. The use of these risers, also known as feeders, exothermic sleeves, or hot tops, can eliminate the problem of shrinkage pores. The quantity, form, and volume of these risers vary according to the form and complexity of the component to be cast. However, despite the fact that there is a vast amount of literature on the calculation and quantity of these risers, the correct location of them depends on the experience of the process controller. The following criteria, however, are considered fundamental for proper feeding of the component, and thus, the defects caused by shrinkage pores are reduced or eliminated:

Thermal transfer criterion: The riser must solidify at the same time or slower than the casting. Volume criterion: The riser must contain enough mass to fulfill the volume shrinkage needs of the component.

However, there are still rules that are eventually observed, and they define additional geometric,

thermal, and pressure criteria that are absolutely necessary for perfect solidification:

The junction between the casting and the riser must not create a hot spot. This place cannot have a larger solidification period than the riser or the casting component; otherwise, it can cause the formation of a shrinkage porosity. There must be a way in which the liquid metal of the riser can reach all the required regions. There must be a pressure variation in order to cause a liquid material flow in the right direction. There must be enough pressure in all the regions of the mold to avoid the formation and growth of cavities.

Internal Porosity Formed from the Surface. If there is not enough internal pressure inside the component being cast and if the liquid inside the mold is still connected with the liquid in the external surface, it can be sucked to the inside, causing the growth of porosities that are connected with the surface (Fig. 12), because the liquid naturally drags air with itself that stays in the interdendritic spacings of the casting component. This preforming mechanism is much more common than imagined. It occurs mainly in alloys with a very long cooling range, when the development of the dendritic lattice means that the aspiration of liquid in the neighboring surfaces becomes easier than feeding from a more distant point. The point at which the liquid may be pulled from the surface can be anywhere for an alloy with a long enough cooling period. Thus, in an alloy with an intermediary solidification period, the starting point is usually a hot spot, such as an internal corner or a recess angle of the component. The possibility of the connection of two opposed surfaces in the same component through the pores is one of the main reasons that alloys with long solidification times should not be employed in the manufacturing of components where high working pressures are applied, such as hydraulic valves or motor cylinder heads, because they would cause leakages. The prerequisite in such complex components is that the interior should have positive pressure in all points in order to avoid the connection of the surfaces through internal porosities, which is rarely achieved. Internal Porosity from Nucleation. Alloys with very short solidification intervals, such as

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the aluminum, brass, and eutectic aluminumsilicon alloys, do not present the connection problem between the surface and pores because they have a perfect solid layer in the first stages of solidification, while the liquid feeding occurs through the feeding channels. The internal pressure decrease due to an inefficient feeding at the end of solidification can create a pore through nucleation in the interior of the liquid. In this case, there is no connection with the external surface of the casting. So, in this kind of alloy, the porosity is usually nucleated and is concentrated near the center of the component. When it occurs in plates, for example, it is referred to as axial porosity. Unless subsequent machining operations pass through the pores, the casting in such alloys is tight. After nucleation, the subsequent solidification provides the necessary driving force for pore growth, which, if observed structurally, has many similarities with the one started from the surface. Growth of the Shrinkage Pores. The first stage of shrinkage pore growth is very fast. According to Davies (Ref 15), this period should be less than 60 ms. After this first nucleation stage, the growth of the pore happens more slowly, being controlled by the heat extraction rate of the mold. For the shrinkage pores that started from the surface—the primary shrinkage cavities, for example—there is not a fast first

Fig. 12

Internal porosity formed from the surface

stage of nucleation. In fact, such a pore or shrinkage cavity is simply formed as an answer to the shrinkage of the solidification. During the solidification of a casting component, the liquid flow from the reservoirs to the areas that are being solidified make the level of the reservoirs decrease. At the same time, there is the advance of the solidification front. This joint action of decreasing the liquid level and the advance of the solidification front creates a conic cavity, as shown in Fig. 13. This cavity is called a primary shrinkage cavity in order to differentiate it from a secondary shrinkage cavity, which is porosity islands observed from longitudinal cuts guided according to a line from the thinnest region of the primary shrinkage cavity. In fact, these islands are interconnected in the solid volume and are thus an extension of the primary shrinkage cavity. Example 1: Failure Analysis of a Mill Gear with Defect Caused by a Shrinkage Pore. The analyzed component corresponds to a mill gear of 860 by 1900 mm (34 by 75 in.) external diameter and 1050 mm and 15 teeth, as shown in Fig. 14. The mill gear was purchased in 2003 and fractured in service during the 2003– 2004 harvest (only 3 months working). The aim of this study was to verify the metallurgical properties of the material and the causes that eventually could have contributed to generate the crack nucleation and the component fracture after a short period of working.

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Fig. 13

Fig. 14

Primary and secondary shrinkage cavities

Aspects of the mill gear as received for analysis

A chemical analysis of the studied component was carried out by optical emission spectroscopy, and the results are shown in Table 1. A tensile test was performed according to ASTM E8M-98, an impact test according to

ASTM E23-91, and a hardness test according to ASTM E0-96 (HBS 2.5/187.5/15). The results are shown in Table 2. The fracture, as can be observed, occurred from the bottom of the tooth, propagating toward the internal diameter. The visual aspect of the cracked surface is an indication that the crack occurred by nucleation and propagation of the cracks through cyclic efforts (fatigue) of the unidirectional type. The final crack happened after the longitudinal section had approximately 30% of its area taken over by cracks, as seen in Fig. 15. The area of crack propagation by fatigue indicates that there was propagation during a relatively short period of the milling operation. Most likely the component started the operation already cracked, that is, with casting defects of considerable dimensions near the surface (subsuperficial), causing the cracks to arise and propagate just at the beginning of the harvest, which caused its fast milling rupture.

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Structural analyses were performed on samples from the fracture region and on the tooth radio adjacent to the fractured region (Fig. 16, 17). It can be noted that the colony of subsurface shrinkage cavities/porosity connected, which probably caused the crack nucleation. It could be concluded that the failure cause was related to casting defects, such as connected shrinkage cavities and porosity colonies, associated with tensile loads applied during the mill gear operation, which caused crack nucleation. Figure 18 shows the crack starting point, proving the failure cause. Example 2: Failure Analysis of a Mill Gear with Defect Caused by a Shrinkage Pore. Figure 19 shows a sliced sample of the mill gear with the tooth root used in the failure analysis. Several mill gear presented had fractures on several teeth roots after an intermittent loading time. The chemical composition in weight percent and the mechanical properties of the steel in Fig. 19 are presented in Tables 3 and 4,

respectively, according to the data provided by the manufacturer and compared to the experimental ones. The phosphorus and sulfur quantities are between the maximum limit established by ASTM A148-93B, and the quantity of other elements agreed with the manufacturer specification. The values obtained for the yield strength (0.2%) and for tensile strength agree with the expected ones. The mill gear fractured through the mechanism of crack propagation by fatigue. The cracks nucleated from the casting defects, located mainly in the third component of the width of the mill gear and near the surface of the tooth root. The low cooling rate during the solidification process is probably the main cause of the high susceptibility to casting defect formation, such as shrinkage pores. Inclusions and bubbles represent a small component in the material embrittlement, since they are too small compared to shrinkage pores. Their presence should be neglected.

Table 1 Chemical analysis of the studied material Composition, wt% Specification

ASTM A148 Gr 105-85

C

Si

Mn

P

S

Cr

Ni

Mo

Cu

Al

0.270

0.477

0.859

0.023

0.017

0.946

1.774

0.242

0.058

0.073

Table 2 Properties of the studied material Test

Tensile strength, MPa Yield strength, MPa Yielding at 50 mm, % Reduction in area, % Impact, J Hardness, HB

Specification ASTM A148 Gr 105-85

804 679 6.0 10.0 30–35 257–259

Fig. 15

Aspect of the fracture surface showing that approximately 30% of the longitudinal section had been taken over by the cracks diffused by fatigue. Many subsuperficial casting defects were also observed where the nucleation of the cracks started.

Fig. 16

Structural analysis performed on samples from the fractured region. Etched with 3% nital. Original magnification: 100 ·


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Penetrating liquid analysis was performed on the tooth root surface, and the result is presented in Fig. 20. Localized cracks on the root center have extended to the sides.

Fig. 17

Structural analysis performed on samples from the fractured region. Etched with 3% nital. Original magnification: 200 ·

Figure 21(a) shows the observed microstructural aspect, where it is possible to note the presence of several cracks in the sample interior. These cracks were nucleated in several shrinkage pores and have propagated by the fatigue mechanism to the surface during cyclic loading. The aspect of the fracture is mainly transgranular, which suggests that the material was not embrittled by drawing back. Beyond the casting defects, inclusions were observed in the sample, some of them with sharp forms, as shown in Fig. 21(b). It is important to note that this kind of shape is undesirable since it is a potential cracknucleating site due to stress concentration. The presence of sharp inclusions indicates that the globalization process during casting was not totally efficient. The proposed corrective actions include:

Increase and standardize the extraction heat rate from the casting mold in the region next to the tooth root Improve the degassing process and impurity control of the casting material Increase the thickness of the on-metal along the region of the tooth root, with the goal of increasing the probability of defect elimination during the machining process

Effects due to Decarburization during Microfusion Fig. 18

Region adjacent to the fractured region showing a transgranular crack generated in the casting process and masked by material deformation during the radio machining process, with propagation directed to the internal diameter

Fig. 19

Sample that was analyzed

Among the several kinds of defects that may occur during the casting process and that are detected after heat treatment is the surface

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Effects due to Cold Joints

decarburization layer. It occurs in carbon steel components cast by the lost-wax or microfusion process. The identification of this defect can only be made after heat treatments in controlled atmospheres; otherwise, this identification is impossible, since the decarburization can also come from the heat treatment. This decarburization results from the presence of atmospheric oxygen that remains in the mold as a consequence of the inert feature of the mold and its permeability in relation to the surroundings. Some analysts have measured the thickness of the decarburized layer in carbon steels cast by lost wax and discovered that it increases proportionally to the temperature increase in the mold and the volume/surface ratio of the casting component (Ref 14). Figure 22 shows a microfused decarburized component that underwent heat treatment.

Cold joint is a kind of defect that happens in cast components and normally has a significant effect on the structural integrity of the component. This serious nonconformity happens when: 1) two portions of the metal, each coming from different feeding/distribution canals of the mold, meet and, instead of contributing to the formation of a smooth and homogeneous surface, provoke an undercut discontinuity called a cold junction; 2) the solidification process occurs too far from the metal flow coming from the feeding/ distribution place, where the liquid temperature is lower than the necessary temperature; 3) the pouring, feeding, and distribution channels are underdimensioned and strangle the flow of metal necessary for the filling of the mold; and 4) the molds have voids that need to be filled with such

Table 3 Chemical composition of the studied material Composition, wt% Source of data

Manufacturer Chemical analysis

C

Si

Mn

P

S

Cr

Ni

Mo

Cu

Al

Fe

0.31 0.31

0.51 0.50

0.78 0.80

0.020 0.018

0.013 0.017

0.76 0.76

1.66 1.66

0.23 0.24

0.06 ...

0.046 ...

bal bal

Table 4 Mechanical properties of the studied material Mechanical properties Source of data

st, MPa

se, MPa

e at 50 mm, %

Reduction in area, %

Hardness, HBW

Manufacturer Experimental

777.0 780.1+16.2

626.0 606.7+15.3

19.0 10.2+1.9

38.6 35.9+14.1

228 230+4

Fig. 20

Cracks located in the tooth root revealed by the penetrating liquid technique. Detail of the central region with higher magnification showing the machining imprints. The arrows show the extreme limits of the cracks.


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a thin thickness that the liquid metal, even at the appropriate temperature, cannot fill them completely. The component that shows this kind of defect, depending on the size and location of the joint, must be discarded, since recovery with a weld is not recommended from a metallurgical point of view or, depending on the cost-benefit ratio, is not justified. This defect is usually seen, but

Fig. 21

can occur and be unnoticed initially, in components with complex geometry and abrupt variation of mass, where it is used to obtain a large number of cores that could provide details difficult to be observed by quality control. In these cases, the defect will only be located when there are cracks/disruption in heat treatment or leakage and fracture when the component is in service.

(a) Micrograph showing cracks connecting shrinkage pores (indicated by arrows) in the internal component of the sample. (b) Detail of the box in (a), where an inclusion is indicated by the arrow

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Fig. 22

Surface of a microfused component showing surface decarburization

In summary, to avoid the appearance of cold joints in cast components, it is necessary to control several manufacturing stages of its design; for example, prevent the component from having regions with very thin thickness; appropriate fusion and pouring temperatures for each component; appropriate mold-filling channel system; compatible pouring speed; and well-established necessary amount of liquid metal for filling the mold to avoid temporary interruption in pouring.

Inclusions Inclusions can be defined as nonmetallic and sometimes intermetallic phases embedded in a metallic matrix (Ref 16). They are usually simple oxides, sulfides, or nitrides. In almost all instances of metal casting, they are considered to be detrimental to the performance of the cast component. Sometimes, an intentional introduction in larger quantities can lead to unique dispersion-strengthened materials. There are essentially two classifications for all inclusions:

Exogenous—those derived from external causes Indigenous—those that are native, innate, or inherent in the molten metal treatment process

Slag, dross, entrapped mold materials, and refractories are examples of inclusions that would be classified as exogenous. In most cases, these inclusions are macroscopic or visible to the naked eye at the casting surface. When the casting is sectioned, they may also appear beneath the external casting surface if they have had insufficient time to float out or settle due to the density differences with respect to the molten metal. The presence of these

macroinclusions in steel castings is avoidable, but their presence has plagued all forms of steel casting and is particularly problematic in both foundry processing and in the continuous casting of sheet steels and wire. Macroinclusions are always practice related, and analysis of the size and chemical composition of a macroinclusion can lead to the identification of potential sources of this problem. Once an inclusional source is developed, a clear and effective process change can be made to eliminate such problems in the future. Therefore, the techniques already developed by integrated steel manufacturers can be readily applied to foundries by coupling inclusion identification with an in-depth study of steelmaking and casting practices in the foundry. Horwath and Goodrich (Ref 17) and Svoboda et al. (Ref 18) have studied macroinclusions and identified that these kind of inclusions can result in excessive casting repairs or rejected castings. To reduce these problems, a method was developed to ensure that there are no inclusions in cast materials above a size that results in failure during ultrasonic or visual inspection of the casting. In this method, the macroinclusions should be eliminated; that is, inclusions greater than 100 mm must be eliminated, but more severely, inclusions greater that 50 mm should be eliminated also. Sulfides, nitrides, and oxides are examples of indigenous inclusions that result from chemical reactions of the molten metal and the local environment. They are usually very small and uniformly distributed inclusions, requiring optical microscopy to visualize them. The presence of these microinclusions in castings is generally unavoidable (Ref 9), because they are the natural inclusions that are formed in liquid steels due to the reaction between alloying elements and oxygen; however, it is necessary to minimize these inclusions as a grain-boundary distribution of these inclusions can be damaging to the component mechanical properties. Clean Steel Clean steel is the common name attributed to steel that has low levels of the elements sulfur, phosphorus, nitrogen, oxygen, and hydrogen, as well as residual elements copper, lead, zinc, nickel, chromium, bismuth, tin, antimony, and magnesium and almost no oxide product defects produced during the act of steelmaking, ladle

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metallurgy, casting, and rolling. Because the “clean” concept is not absolute, the cleanliness standard desired by the customer is continuously changing as a function of time and technological improvements. The term clean steel is therefore continually variable, depending on the application and the competition between steel suppliers. Thus, due to the variability of the term clean, it is typical to refer to high-purity steels as steels with low levels of solutes, and low-residual steels as steels with low levels of impurities. For example, there are high-purity, low-residual clean steels, such as ultra-deep-drawing steel sheets for automobiles, that require ultralow carbon contents (530 ppm), low nitrogen contents (530 ppm), and the absence of oxide inclusions with diameters greater than 100 mm; and there are low-residual clean steels, such as those used for drawn and ironed cans, that are a standard low-carbon steel (1006) without highpurity component requirements but are ultraclean, with the requirement that oxide diameters must be less than 20 mm. In addition, in forging and bearing grades, there are clean steels that require strictly controlled inclusion size distributions. The total inclusion content related to the total oxygen content has been correlated with bearing life, and decreasing total oxygen contents (below 10 ppm) improve the bearing life. In addition to total oxygen content, the total length of stringer inclusions after forging is also related to the bearing life, and, at low total oxygen levels, efforts to reduce inclusion clustering lead to very long fatigue life for bearings. Clean steels can be classified as steels with a low frequency of inclusions (55 mm). The major problems in clean steel manufacture are incomplete separation of clustered solid inclusions (45 mm in diameter), the presence of sporadic larger liquid inclusions due to emulsification of covering slags, and the presence of solid materials that originate from the refractories used to contain steels. The equipment used to produce clean steel varies greatly between different steel plants; however, current clean steelmaking and casting practices are based on the following principles:

The oxygen dissolved in liquid steel at the melting stage must be transformed into a solid or a gas and removed before casting. The external oxygen sources that are responsible for the reoxidation of liquid steel

must be eliminated at every step in the process. The physical entrapment of the liquid fluxes used during steel refining and casting must be eliminated. Refractories in contact with liquid steel must be chemically stable and resistant to corrosion and erosion.

These simple principles are based on the importance of maintaining chemical equilibrium between the elements dissolved in liquid steel and the slag and refractory systems that are in contact with the liquid steel. Additionally, it is necessary to control the fluid flow to avoid conditions at liquid slag-steel interfaces that could result in the physical entrapment of the covering slag. Clean steel manufacture is dependent on an understanding of the fundamental steps necessary to produce a clean steel:

Generation of the inclusion Transport of the inclusion to an interface Separation of the inclusion at the interface Removal of the inclusion from the interface

The production of really clean steel depends of the correct application of these principles. The Formation of Macroinclusions There are four major methods of forming macroinclusions, and all problems occur during foundry processing:

Reoxidation Interaction between liquid steel and liquid slags: vortexing, ladle or mold filling, argon stirring, and pouring through a slag layer Erosion/corrosion during steel pouring Inclusion agglomeration due to clogging during steel pouring

Reoxidation. The major cause of macroinclusion formation in casting is reoxidation (Ref 17–19). To understand reoxidation, it is necessary to understand that liquid iron is not thermodynamically stable in the presence of oxygen. The spontaneous reaction that occurs results in the formation of iron oxide. As deoxidizers are added, the steel remains unstable in the presence of oxygen as a gas, but now the inclusions that form include the oxides of the deoxidants. Some deoxidants, such as aluminum, magnesium, and calcium, form very stable oxides that are more stable than some slag and

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refractory chemistries. Thus, the steel reacts with the less stable oxides. Reoxidation can occur by reaction with:

The ambient atmosphere (air) The slag components less stable than the oxide of the deoxidant The refractories that are less stable than the oxides of the deoxidant

Interaction between Liquid Steel and Liquid Slag. Macroinclusion formation can occur by emulsification of liquid slags or scums on the surface of liquid steels. All of these types of defects are practice related and can be solved by practice changes. The issue in understanding emulsification is to understand the source of the energy that allows a buoyant droplet to become submerged. Generally, this energy comes from the interaction of a flowing steel stream and a liquid slag. There are four major sources of this energy:

Open stream pouring onto or through a liquid slag (common during lip pouring) Filling a ladle or mold at too high a fill rate in the presence of slags or scums Vortexing during steel pouring from a ladle Steering in the ladle with gas at too high a stir rate

Vortexing during drainage in a water model of a ladle was studied by Sankaranarayanan and Guthrie (Ref 20, 21). They showed that the initial rotational velocity at the surface of the vessel is extremely important in determining the height at which the vortex will form, and that increased rotational velocities caused increased vortex initiation depth. Entrainment due to fluid flow at the interface has been examined by Noguchi et al. (Ref 22), who attempted to decrease the entrainment of slag in low-carbon titanium-aluminum-killed steels. They noted that entrainment decreased as the casting speed was decreased. In a study conducted by Nakamura et al. (Ref 23), it was found that defects that contained mold slag increased in ultra-lowcarbon grades as the casting rate was increased. They also reported using as low an argon flow rate as possible in their submerged entry nozzles to avoid entrainment. Manabu et al. (Ref 24) have also documented the existence of a critical gas flow rate for entrainment in both a silicon oil-water and a slag-steel system. These authors mention that the slag depth, slag properties, and gas bubble diameter play a role. The oil depth

was found to be directly proportional to the flow rate needed to cause entrainment. The gas bubble size was found to be inversely proportional to the flow rate needed to cause entrainment. Manabu et al. (Ref 24), investigated the effect of oil kinematic viscosities on emulsification and found that although the kinematic viscosity was varied by a factor of 10, very little change was seen in the fluid velocity needed to cause entrainment. Harman and Cramb (Ref 25), documented the effect of interfacial tension and slag viscosity on emulsification phenomena. Erosion-Corrosion during Steel Pouring. This kind of defect is usually associated with the higher corrosivity of some steel grades, because high manganese and grades that are barely killed and have high soluble oxygen contents attack the binder or the mold sand itself, leading to large entrapped sand components. Reoxidation of steel leads to FeO-based inclusions that are very reactive and wet the materials of the mold, leading to erosion of the mold in areas of high fluid turbulence. Of course, sand that is not pressed, sintered, or bonded in any way can easily be entrapped in turbulent fluid flow. Mold binders can also decompose at temperature and release mold components that can be entrapped. Expansion due to the high thermal gradients associated with casting can also cause sand to loosen. Inclusion Agglomeration due to Clogging during Steel Pouring. The formation of clogs when steels containing solid inclusions are cast can result in quite large macroinclusion defects if the clogs are released during teaming. All solid inclusions tend to agglomerate due to surface tension effects. Clogging of pouring nozzles can be the source of large macroinclusion defects when steels are dirty and pouring times are long. The Formation of Microinclusions Microinclusions are formed due to reactions between alloying additions and oxygen in molten steel. Their formation is generally heterogeneous or from highly supersaturated areas during alloy addition. Due to the nature of the formation of these inclusions (nucleation and growth), they are generally small (less than 5 mm), unless they agglomerate due to turbulence or grow under conditions of high oxygen flux. In this study, microinclusions are defined as those inclusions with diameters smaller than 20 mm. In addition, they are defined as having

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diameters greater than 1 mm. Table 5 show typical microinclusions that are found in cast steels. Since microinclusions form due to a reaction, they are driven by thermodynamics; therefore, changing composition or temperature can lead to their precipitation. This means they can form in the ladle, during transport to the mold, or in the mold during solidification.

indicated. Figure 25 shows a micrograph of the fractured surface, near the blade bottom. Several turning gear imprints can be observed, showing the presence of multiple sites of crack nucleation

Case Studies of Defects Caused by Inclusions Failure of a Steam Turbine Rotor Blade. Possible causes were investigated for failure of a rotor blade of a 35 MW steam turbine. One of the rotor blades was fractured after a certain operation time (Fig. 23). The fracture occurred at two different regions: at the bottom and at the top extremity, near the metallic lashing strap. Both regions have the highest stress concentration due to the blade geometry and loading conditions. The blade fracture occurred during the maximum turbine operation. The rotor was working, with new blades mounted in between harvests. The blades were manufactured with steel ingots with the chemical composition presented in Table 6. The specifications for the mechanical properties of the material at room temperature are shown in Table 7. Figure 24 shows the fractured blade compared to an intact one, with the fracture regions

Fig. 23

(a) Turbine stage that had the fractured blade. (b) Detail of the fractured bottom component of

the blade

Table 5 Typical microinclusions found in cast steels Steel type

Microinclusion type

Aluminum killed Manganese-silicon killed Calcium treated, aluminum killed Aluminum killed, with residual magnesium Titanium treated, aluminum killed All steels

Comments

Alumina Manganese silicate or manganese-alumino silicate Calcium aluminate Magnesium aluminate

Formed in liquid steel after deoxidation Formed in liquid steel after deoxidation

Alumina, titania, titanium nitride

Titania forms during reoxidation. Titanium nitride forms during cooling, usually in the mold itself. Forms interdendritically during solidification. Often nucleates on oxides already present in steels

Formed by reaction with alumina, liquid inclusion Formed by reaction with alumina, solid inclusion

Manganese sulfide

Table 6 Nominal chemical composition of FV520(B) steel Composition, wt% C

0.07 max

Si

Mn

Cr

Ni

Cu

Mo

Nb

S

P

0.7 max

1.0 max

13.2–14.7

5.0–6.0

1.2–2.0

1.2–2.0

0.2–0.5

0.06 max

0.03 max

Table 7 Mechanical properties specifications of FV520(B) steel Yielding limit, MPa

680–800

Strain limit, MPa

Elongation, %

Reduction in area, %

Impact energy, J

Hardness, HV

900–1050

20 min

55 min

40 min

270–320


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Fig. 24

Fig. 25

(a) Intact blade. (b) Fractured blade

Micrograph of the blade fracture surface showing several turning gear imprints and the oxidized area

(dotted line)

by fatigue. A darkened region is observed on the fracture surface, indicated by the dotted line, suggesting that this area was more exposed to steam and high temperatures during the turbine operation time, and it occupies a significant component of the fracture surface. Penetrating liquid analysis indicated the presence of secondary longitudinal cracks in the fractured material, normal to the main crack, at

the bottom of the blade. The analysis made in the blade body indicated the presence of a large, longitudinal crack, probably consisting of an extension of the cracks observed at the bottom of the blade, as shown in Fig. 26. Optical microscopy analysis of a cross section of the blade body revealed a different microstructure from the martensitic steel matrix located parallel to the longitudinal crack in the blade body. Because of this different microstructure, electron-dispersive x-ray (EDX) analyses were carried out in the regions around the longitudinal crack in the structural sample. They showed a chemical composition different from the nominal, as much for the central region as for the blade head region. The fracture surface of the longitudinal crack revealed a microstructure rich in silicon, oxygen, manganese, and calcium, suggesting that the material contains a large number of impurities, probably slag from the casting process and certainly introduced during the manufacturing process of the component. The occurrence of these impurities impedes surface welding during the process of forging, creating a surface with a smashed aspect. However, the first region where the nucleation probably occurred was the one near the

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longitudinal cracks detected by the penetrating liquid. Indeed, the fractographic analysis of this region shows a fracture morphology different from the vicinity, with several inclusion components protruding into the fracture surface (Fig. 27). The EDX microanalyses of these components show a chemical composition with a high level of carbon, which suggests that these components are of iron carbide. Moreover, several longitudinal cracks similar to the one found in the blade body were observed. Nonfusible longitudinal cracks exist along the affected area in the blade. The large variety of defects and the excessive mechanical vibration of the blade are probably the main causes of crack nucleation by fatigue in the material near the blade bottom. They culminated in the catastrophic fracture of the component. The recommendation includes a more efficient quality control of the manufacturing process of the blade material and avoiding the occurrence of casting defect formation, slag inclusions, and other impurities. Failure in the Axle of a Reduced Section in a Rotating Component. Possible causes were investigated for failure in the area of an intermediate reduction. The rotating component fractured completely after intermittent loading. Figure 28 shows an outline of the component and the axle region where the cracks developed. The chemical composition (in weight percent) of the fractured axle material is provided in Table 8. The results show that the axle material is a DIN-specified 17CrNiMo6 steel. The specifications of the material mechanical properties at room temperature are given in Table 9. The visual inspection of the fracture surface (Fig. 29) indicated an extremely flat aspect, such as the ones typically displayed in fatigue cracks.

Fig. 26

The flat fracture surface occupied approximately 80% of the cross section (Fig. 29), exactly in the axis of the radius change for the concordance section. Due to the small relative section area of the fracture axis, approximately 20% of its cross section, it was deduced that the stress for the inservice component was relatively low. Ten measurements of Rockwell C hardness were carried out, according to ASTM E18, on the surface of the axle near the fracture region. A mean hardness of 33.9 HRC was obtained. This value is well below the expected one of 43 HRC. Figure 30 shows the microstructure of steel in the reduced section on a longitudinal cut plane in the vicinity of fatigue crack nucleation. The material presents a large amount of globular or granular bainite, in agreement with the relatively low value of hardness of the fractured axle surface. Figure 31(a) shows a general topview of the fracture surface in the region where there was fatigue crack nucleation, indicated by the arrow at bottom. The five clustered arrows point in the direction of fatigue crack propagation advance. The arrow at the top shows a dark region, originated by contamination of the fracture surface with oil or grease. Figure 31(b) shows in detail the fatigue crack nucleating site that probably started at an inclusion located exactly on the circumference surface of the reduced section in a region in the internal concordance radium. The presence of some inclusions in the proximity of the fracture site is pointed out by white arrows in Fig. 31(b). Indeed, fractographic analysis has shown the possibility of the existence of a concentration of inclusions in the nucleation region of the fatigue crack. Figure 32(a) confirms the high level of inclusions in the region, indicated by white arrows, with signs of moving

Longitudinal crack in the blade body revealed by penetrating liquid. The A-A section indicates the approximate position of the cut made for structural observation.

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by second-phase components indicated by black arrows, similar to Fig. 31(b). Figure 32(b), the same image shown in Fig. 32(a) but with backscattered electrons instead of secondary ones, reveals the great amount of inclusions (darker) in the metallic matrix (lighter). The chemical

Fig. 27

(a) General view of the probable initial region of crack nucleation by fatigue crack. (b) Magnification of the region in the box at the left in (a). (c) Magnification of the region in the box at the right in (a)

analysis of the inclusions shows a massive presence of aluminum, sulfur, and calcium elements. It is worth noting that these inclusions act, on a microscopic scale, as metallurgical stress concentrators. The presence of these secondphase components especially near the external axle surface where the maximum tensile stresses are developed during a torsional load (and even flexion) applied to the in-service component, drastically reduces the lifetime in fatigue of the rotating component. This happens through the promotion of both mechanisms of nucleation and fatigue crack propagation in their early stages of growth. It was concluded that crack initiation occurred in the reducer axle by fatigue. A single crack probably was nucleated on a nonmetallic inclusion placed near the finished axle surface, exactly in the internal component of the concordance radius machined in the section change. The combination of the effects of stress concentration generated by both discontinuities, metallurgical (inclusion) and geometric (curvature radius), created sufficient critical conditions for fatigue crack nucleation that grew due to the action of repetitive efforts of torsion (and flexion) imposed in service to the rotating component. Failure of a 52100 Steel Axle. The raw material (52100 steel) used in the manufacture of an axle catastrophically fractured during annealing heat treatment at 350 C. Figure 33 shows the fracture surface along with the circular cross section of the component (one of the samples received for analysis). In the figure, the arrow at left shows the main fracture plane of the axis (i.e., along a longitudinal plane), and the arrow at right points to the starting point of brittle fracture in its cross section. In Fig. 34, this starting site is shown in detail (arrow at bottom). Figure 35 shows the microstructure of the 52100 steel, in the central region of the part in a longitudinal plane, after etching with nital. The massive presence of pearlite and the existence of free cementite in both forms—globulized (inside the pearlitic colonies) and veins (circling the colonies)—is observed. Figure 36 shows the vermiform discontinuities, with an appearance similar to manganese sulfide inclusions, that are invariably present in mechanical construction steels. The presence of a grayish second phase, intermediate to the metallic matrix (lighter) and the voids

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Fig. 28

Component drawing of the intermediate I axle. Highlighted are the section change region where the fracture developed and the crack propagation path for the total fracture of the axle.

Table 8 Chemical composition of the axle material

Table 9 Mechanical properties at room temperature sE, MPa

Composition, wt% C

Mn

Si

P

S

Ni

Cr

Mo

0.17

0.63

0.23

0.10

0.011

1.45

1.59

0.30

742

sR, MPa

AF, %

QF, %

1080

20

57

Fig. 29

Complete cross-sectional fracture surface of the intermediate I axle. The white arrow shows the nucleating site of the fatigue crack. The surface generated by the fatigue crack propagation is identified by “F,” while the final fracture of the remaining section is indicated by “FF.”

Fig. 30

Microstructure of the axle according to a longitudinal cut plane. Etched with 2% nital


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(darker), is observed inside these discontinuities. This material component fills the larger discontinuities, while the smaller discontinuities are almost totally filled by the second phase. Figure 37 shows that the most subtle discontinuities have a rather slim, cracklike aspect and consequently present a great capacity to concentrate high tensile stresses. In these terms, it is possible to assume that these second phases are potential crack nuclei, and that they

also generate a preferential path for crack propagation. It is worth emphasizing that the majority of these discontinuities were found aligned in the direction of the thermomechanical work to which the axle was submitted during its manufacture (i.e., longitudinal direction). The inclusions are disposed on parallel planes to the main fracture of the component during heat treatment. This suggests the possibility that these inclusions played a fundamental role in the catastrophic failure of the 52100 steel axle.

Fig. 31

Fig. 32

Fatigue crack site. (a) General view. (b) Detail. The inclusion that originated the site was removed from the fracture surface. SEM image with secondary electrons

Concentration of inclusions near the fatigue crack site. (a) SEM image with secondary electrons. (b) Backscattered electrons

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Figure 38 shows elongated microvoids, obtained by SEM with secondary electrons, in the vicinity of the tip of one of the cracks that propagated in the fractured component. The

alignment of the discontinuities generate a favorable path for material cracking. The voids that are already interconnected by material cracking are shown by the arrows in Fig. 38. In Fig. 39, the 52100 steel microstructure in the central axle region, cut in the longitudinal

Fig. 33

Cross section of a catastrophically fractured axle. The arrow at left shows the main fracture plane (longitudinal), and the arrow at right shows the starting point of the fracture in the circular cross section.

Fig. 36

Fig. 34 Detail of the starting point of brittle fracture in the circular cross section of the component (arrow at bottom). The clustered arrows show the brittle crack tip front.

Fig. 35

Fractured axle microstructure at the center of the component thickness. Etched with nital. Original magnification: 400 ·

Inclusion-like microdefects detected in the vicinity of a crack in the fractured axle, located at the center of the component

Fig. 37

Slim, cracklike inclusions in the 52100 steel. The inclusions are oriented in the longitudinal direction of the component. No etch

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Fig. 38

Alignment of the elongated inclusions (oriented in the longitudinal direction of the part) act as an easy propagation path in the 52100 steel axle. The main fracture direction, that is, longitudinal, corresponds exactly to the elongation and inclusion alignments. The arrows point to the existence of cracks among the microcavities that compose the inclusions. SEM Original magnification: 100 ·; 20 kV

plane, is shown in greater detail. An essentially pearlitic matrix developed with cementite precipitates (Fe3C) in the globular form (solid arrows). Free cementite exists in the pearlitic colonies contour in the form of veins or platelets (white arrows) that offer an easy path for brittle crack propagation in the material. The absorption spectra obtained by EDX of the 52100 steel confirm that the plate precipitates, shown in Fig. 39, are made of iron carbide or cementite in the free form (Fe3C). However, the absorption spectra obtained in microanalyses of the grayish material inside the elongated microcavities, shown in Fig. 36 to 38, indicated it is made essentially of iron oxide. At first, the hypothesis that this contaminant comes from, for example, the atmospheric oxidation after the fracture event of the component, was discarded, since the inclusions measured by microprobe were completely isolated inside the metallic matrix, without any possibility of reaction with the environment. It was concluded that the raw material used to manufacture the fractured axle was probably contaminated with iron oxide. The contaminant was in the form of elongated inclusions, aligned in the longitudinal part direction, making an easy path for main crack propagation (longitudinal). The elongated format provided the inclusions the capability to concentrate high tensile stresses and then transform them into potential crack nucleation sites. Microanalysis also confirmed the existence of free cementite in the pearlitic grain contours

Fig. 39

52100 steel microstructure in the center of the component thickness. Etched with nital. Solid arrows point to free cementite in the globular form, and white arrows point to Fe3C in the form of platelets in the pearlite contour. Original magnifications: (a) 3000 ·. (b) 10,000·. (c) 18,000·

that form the 52100 steel. The presence of this fragile phase may have contributed, to a certain extent, to the intergranular secondary brittle crack propagation in the catastrophic fracture of the component during annealing heat treatment.

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ACKNOWLEDGMENTS Thanks to the Department of Materials, Aeronautics and Automotive Engineering of the School of Engineering of Saõ Carlos, University of Saõ Paulo, on behalf of Professor Dr. Dirceu Spinelli, for the collaboration on failure analysis case studies.

REFERENCES

1. Defects and Distortion in Heat-Treated Components, Heat Treating, Vol 4, ASM Handbook, ASM International, 1991, p 1320–1325 2. R.E. Reed-Hill, Physical Metallurgy Principles, 2nd ed., Van Nostrand, 1982 3. Casting Design, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 1301–1322 4. Dimensional Tolerances and Allowances, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 1336–1347 5. Quenching and Control of Distortion, ASM International, 1988 6. Residual Stress, Heat Treating, Vol 4, ASM Handbook, ASM International, 1991, p 1325–1343 7. M.T. Milan, O. Maluf, D. Spinelli, and W.W. Bose Filho, Metais—Uma Visaõ Objetiva (Metals A Vision Object) Suprema, 2004, p 148–149, 161–162 8. Gases in Metals, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 175–189 9. Inclusion-Forming Reactions, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 190–211 10. R.D. Pelke and J. Elliott, Trans. TMSAIME, Vol 227, 1963, p 894 11. P.C. Glaws and R.J. Fruehan, Metall. Trans. B, Vol 17, 1986, p 317 12. R.J. Fruehan, B. Lally, and P.C. Glaws, Proceedings of the Fifth International Iron and Steel Congress (Washington, D.C.), Iron and Steel Society of AIME, 1986 13. A. Kagawa and T. Okamoto, Trans. Jpn. Inst. Met., Vol 22 (No. 2), 1981, p 137

14. J. Campbell, Castings, ButterworthHeinenmann, 1993 15. J.G. Davies, Solidification and Casting, Applied Science, 1973 16. AFS Inclusion Atlas Homepage, http:// neon.mems.cmu.edu/afs/afs2/ (Accessed on March 2005) 17. J.A. Horwath and G.M. Goodrich, MicroInclusion Classification in Steel Casting, AFS Trans., 1995, p 495–510 18. J.M. Svoboda, R.W. Monroe, C.E. Bates, and J. Griffin, Appearance and Composition of Macro-Inclusions in Steel Castings, AFS Trans., 1987, p 187–202 19. C.R. Wanstall, J. Griffin, and C.E. Bates, “Clean Steel Cast Technology,” Research Report 106, Steel Founders Society of America 20. R. Sankaranarayanan and R. Guthrie, Slag Entrainment through a “Funnel” Vortex during Ladle Teeming Operations, Proceedings of the International Symposium on Developments in Ladle Steelmaking and Continuous Casting, Aug 1990 (Ontario), CIM, 1990, p 66–87 21. R. Sankaranarayanan and R. Guthrie, A Laboratory Study of Slag Entrainment during the Emptying of Metallurgical Vessels, Steelmaking Conference Proceedings (Ontario), 1992, p 655–664 22. K. Noguchi et al., Zairyo to Purosesu (Curr. Adv. Mater. Process.), Vol 4, 4th ed., 1991, p 1194–1197 23. H. Nakamura, S. Kohira, J. Kubota, T. Kondo, M. Suzuki, and Y. Shiratani, Technology for Production of High Quality Slab at High Speed, Steelmaking Conference Proceedings (Ontario), 1992 24. I. Manabu, S. Yutaka, O. Ryusuke, and M. Zen-ichiro, Evaluation of the Critical Gas Flow Rate Using Water Model for the Entrapment of Slag into a Metal Bath Subject to Gas Injection, Tetsu-to-Hagane (J. Iron Steel Inst. Jpn.), Vol 79 (No. 5), p 33 25. J.M. Harman and A.W. Cramb, A Study on the Effect of Fluid Physical Properties on Droplet Emulsification, Steelmaking Conference Proceedings, 1996, p 773–784

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Sources of Failures in Carburized and Carbonitrided Components Małgorzata Przyłe˛cka and Wojciech Ge˛stwa, Poznan University of Technology Lauralice C.F. Canale, University of Saõ Paulo Xin Yao, Portland State University G.E. Totten, Associaçaõ Instituto Internacional de Cieˆncia and Portland State University

MANY COMPONENTS, such as fasteners, crankshafts, camshafts, bearings, and others, require a differentiated response of the surface and core to external loading. This can be accomplished by surface (case) hardening methods such as induction and flame hardening or by surface diffusion processes such as carburizing and carbonitriding. Raja et al. have reported that case carburizing is one of the most common heat treatments for steel, accounting for 50% of all surface treatments (Ref 1). Case carburizing involves the creation of a gradient that exhibits high hardness, brittleness, and strength in the surface and greater toughness and ductility in the softer core in order to provide optimal (Ref 2):

Wear resistance Resistance to scoring Bending and/or torsional fatigue strength Rolling-contact fatigue strength

These properties are optimized by maximizing surface compressive stresses, and carburizing is one of the most effective and commonly used methods to impart compressive stresses to the surface of a component (Ref 3). The focus of this chapter is on carburized and carbonitrided materials. Gas carburizing, which is the most widely used carburizing process, is a surface diffusion process where the carbon concentration in a surface layer (case) of a steel matrix that is predominantly iron, chromium, and nickel is increased by heating the component at approximately 850 to 950 C with endothermic gas (Endogas), which is a blend of carbon monoxide, hydrogen, and nitrogen (with smaller

amounts of carbon dioxide, water vapor, and methane). Endogas is produced by reacting a hydrocarbon gas, such as natural gas (methane), propane, or butane, with air. After the diffusion process is completed, the component may be quenched from the carburizing temperature or reheated to austenitize the steel, and then quenched. Bainite formation in the case is strongly inhibited by the presence of molybdenum and chromium. Since the surface contains higher carbon content than the core, it is harder than the softer core. Core hardness is most strongly affected by the presence of molybdenum and manganese. Chromium exhibits a moderate effect, and nickel exhibits a weak effect (Ref 3). Core hardness is strongly affected by the quenchant selection and quenching temperature. In addition to strengthening the case, the increased carbon content also provides desirable increased compressive stresses that will inhibit fatigue crack initiation. The lower carbon content in the core also will produce improved fatigue strength. Carbonitriding is similar to carburizing in that it is a diffusion process that involves the simultaneous diffusion of carbon and nitrogen (from ammonia) into the steel surface. To obtain maximum strength, the carbonitriding process produces a surface that is enriched in nitrogen and carbon in the form of an epsilon (e)-carbonitride layer and a diffusion zone containing chromium-iron carbide, (Cr,Fe)7C3; chromium carbide nitride, Cr62C3 5N0.3; chromium nitride, (Cr2N) or [Cr, Fe(2Ni . . . x)]; and Fe2N phases (Ref 4, 5). Typical case thicknesses range from 50 to 200 mm with a hardness

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between 750 and 900 HV. Like carburizing, the case depth of carbonitrided steel is dependent on both the carbonitriding diffusion time and temperature, as illustrated in Fig. 1 (Ref 2, 6). Deeper case hardnesses may be obtained by first precarburizing prior to carbonitriding. Karamis¸ showed that carbonitrided AISI 5115 steel exhibited greater surface hardness and wear resistance than carburized AISI 5115 steel (Ref 4). Carbonitriding processes are typically conducted in either a gas (ammonia) or a salt bath based on trade names such as Tufftride, Nitrotec, and Nitrox. Alternatively, a plasma nitriding process may be conducted. A brief summary comparison of carburizing and carbonitriding processes is provided in Table 1 (Ref 2). Carter has reported that failures of carburized gears are primarily due to service-related causes, such as misalignment, poor lubrication, and overloading, which constitute the greatest source of all gear failures, as shown in Table 2

Fig. 1

Correlation of case depth of carbonitrided steels with varying diffusion times and temperatures

Table 1 Comparison of carburizing and carbonitriding processes Process

Carburizing

Comments

Hard, highly wear-resistant surface (medium case depths), excellent contact load potential, good bending fatigue strength, good seizure resistance, excellent quench cracking resistance, low-to-medium-cost steels required, high capital investment

Carbonitriding Hard, highly wear-resistant surface (shallow case depths), fair contact load potential, good bending fatigue strength, good seizure resistance, good dimensional control, excellent quench cracking resistance, low-cast steels usually satisfactory, medium capital investment

(Ref 7). Heat treatment was the second most often cited cause for failure. However, it is often difficult to detect the root cause of a specific failure under the conditions in which the failure occurred, and many of the service-related failures could have been reduced with more attention to the other potential causes of failure shown, since they are often interrelated. Palaniradja et al. reported that 10 to 12% of carburized parts are rejected due to various process-related defects (Ref 8). To examine this in more detail, they conducted a Taguchi analysis of gas carburization of AISI 8620 and 3310 steels, and their results showed that relative contribution to surface hardness was holding time (20%), carbon potential (20%), carburizing temperature (0%), and quenching time (60%). Similarly, they also studied the effects of process variables on case depth and found: holding time (60%), carbon potential (9%), carburizing temperature (14%), and quenching time (10%) (Ref 8). These results show that an adequate understanding of failure analysis of carburizing and, by implication, carbonitriding must be accompanied by understanding the contribution of process parameters on resulting potential failures. Some of the most common contributors to failure of carburized gears include surface finish, microstructure, excessive or inadequate case depth, incorrect case and/or core hardness, improper carbon concentration and hardness gradients, undesirable surface carbon content, excessive retained austenite, large amounts of globular and network carbides, intergranular oxidation, internal oxidation, residual stress, extremely coarse case or core grain structure, untransformed core with free ferrite, quenching and grinding cracks, surface decarburization, excessive heating during grinding, excessive removal of the case during grinding, microcracking, and so on (Ref 9, 10). This chapter provides an overview of various contributors to failure of carburized and carbonitrided components, with the primary focus on carburized components. Table 2 Survey summary of sources of gear failures Cause of gear failure

Material quality and forming Design Service-related causes Manufacturing Heat treatment

%

0.8 6.9 74.7 1.4 16.2


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Sources of Failures in Carburized and Carbonitrided Components / 179

Design Component design may contribute directly or indirectly to component failure. Deficiencies such as insufficient radii or sudden changes in section size are significant contributors to failure. In addition, the presence of stress raisers, such as those shown in Fig. 2, are among the most common design contributors to quench cracking and fatigue failure. A more comprehensive insight into design is provided by Kuehmann et al., who developed a systems analysis flow chart to describe the effects of case-core hardening in designing a carburizing process/metallurgical structure/ resulting properties and performance for the production of gears produced by three routes: conventional forging, near-net shape casting, and powder metal processing (Fig. 3) (Ref 11). To properly design a component, it is necessary to estimate surface loading, distortion after heat treatment, case depth and carbon profile,

Fig. 2

case and core hardness, and core strength. As an estimate, for hardnesses within the range of 30 to 45 RC, the required case depth can be calculated from (Ref 2): Case depth to 50 HRC=(1:2 · 107 W)=F

where W is the force in pounds pressing the surfaces together, and F is the length of the line contact (inches). Carter has recommended the following general design criteria (Ref 7):

If a component is carburized from both sides, the case depth should not be greater than 20% of the wall thickness. At the base of gear teeth, 30% of the core material should remain uncarburized. Shallow case depths usually require higher case hardness. Case depths should be five times the acceptable wear limit.

Effected of stress raisers on stress concentration and distribution of stress at several changes of form in components. (a) to (c) Progressive increases in stress with decreasing fillet radii. (d) to (f) Relative magnitude and distribution of stress resulting from uniform loading. (g) Stress caused by the presence of an integral collar of considerable width. (h) Decrease in stress concentration that accompanies a decrease in collar width. (i) Stress flow at the junction of a bolt head and a shank. (j) Effect of a single sharp notch. (k) Effect of a continuous thread. (l) Effect of a groove or gauge. Source: ASM Handbook, Volume 11, 2002, p 715

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Although machining is proportional to the case depth, it should be minimal.

Kern and Suess have recommended the following general guidelines for heat treatment of gas-carburized gears (Ref 2):

For forgings, normalize or anneal (as required by the alloy being heat treated) from a temperature at least 28 C (50 F) above the carburizing temperature. Assure that the gears are machined prior to heat treatment. Bring the gear to the carburizing temperature with sufficient circulation of a neutral atmosphere, and then introduce the gas used for carburizing. For deep cases (41.5 mm, or 0.060 in.), adjust the carburizing atmosphere and time to produce uniform carbon diffusion from the surface to the core. A decrease of 0.15 to 0.20%/0.25 mm (0.010 in.) of depth is

Fig. 3

nearly ideal. The gradient may be steeper for shallower case depths. The recommended surface carbon is 0.90 to 1.10% for 4300, 4600, 8600, 8800, and 9400 carburizing steels. Although the same case depth is generally acceptable for grades such as 4800, they are preferably reheated for hardening. The recommended surface carbon is 0.65 to 0.85% for high-nickel steels such as the 4800 series, which is usually direct quenched. To minimize cost and distortion, use direct quenching whenever possible. To assure optimal dimensional control, properly maintain quenching dies and plugs. Quench as rapidly and uniformly as practical, and use spray impingement fixtures on large, solid pinions that are four pitch and coarser. Use hot oil quenching on fine pitch gears.

Kuehmann et al. flow chart to summarize design elements of a carburizing process/metallurgical structure/resulting properties and performance comparison of gas carburizing gears produced by conventional forging, near-net shape casting, and powder metal processing

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Case hardness of the finished gear should be 60 HRC or greater. If possible, test each gear for partial decarburization and/or upper transformation products. To minimize distortion and to permit quieter operation, the surface carbon content should be uniform throughout the production cycle.

Steel Selection and Hardenability Steels typically used for case hardening contain carbon contents of less than approximately 0.25%. The carbon content of the case is usually controlled to between 0.8 and 1% C. The actual surface carbon content is generally limited to 0.9%, because excessively high carbon content may lead to the presence of unacceptably high retained austenite and brittle martensite. Some of the most commonly used AISI grades of steel used for carburizing are shown in Table 3 (Ref 12). Plain carbon steels may be carburized; however, relatively poor hardenability due to the lack of alloying elements reduces the carburizing response of the case. Because of the stabilizing effect of the nitrogen relative to austenite, carbonitriding provides greater hardenability than attainable with carburizing. Therefore, plain carbon steels respond well to carbonitriding. Proper steel selection is a critically important process to provide the desired case depth and microstructure and the required core properties. Typically, the case structure should be fully martensitic, with the exception of allowing for required application design limits on retained austenite content. For example, the steel must

Table 3 Common carburizing grades of steel and their relative processing features AISI steel grade

4620 8620 4320 4820 9310

Note

Lower-cost, chrome/nickel/molybdenum steel where only nominal hardenability and core response is required Most commonly specified grade. Excellent carburizing response, with good hardenability for most section sizes Higher hardenability for improved core response in heavier sections Increased nickel content for improved core toughness; slower response results in longer process times Maximum nickel content for maximum core toughness; slower response results in longer process times

possess sufficient hardenability to provide the desired hardness and microstructure in both the case and the core. After carburizing, the component must possess sufficient toughness without exhibiting brittle failure. Most steels that are carburized are deoxidized by the addition of aluminum (commonly designated as killed steels). Deoxidation will provide finer grain sizes to temperatures of approximately 1040 C. Coarser grained steels may be carburized if grain refinement by double quenching is possible. Double quenching typically involves direct quenching followed by reheating to a lower-temperature quenching a second time (Ref 13). Selection of proper hardenability of steels for both carburizing and carbonitriding is critically important, both of the core and the case, since improper hardenability design can lead to undesirable nonmartensitic transformation products in the case, leading to a potential reduction in static and dynamic fatigue strength of up to 30% and a reduction of impact fatigue of up to a factor of 2.5 times (Ref 14). The hardenability gradient of the case and the core is dependent on a number of factors, including cooling rate during quenching, variability of the chemical composition (alloy content, carbon and nitrogen) of the case, and the carburizing or carbonitriding method being used. Core hardenability is being used increasingly to specify alloy steels used for case hardening where the hardenability of both case and core must be considered. Details for the traditional approach for the experimental determination of hardenability of carburizing steels are provided in Ref 15. Jominy curves for a number of carburizing steel alloys with varying hardenability are shown in Fig. 4 (Ref 16). Procedures have also been described for determining ideal diameter (DI) values and hardenability of carburizing steels from Jominy data using regression equations for composition and grain size (Ref 17, 18). The ideal diameter is defined as the diameter of a cylindrical steel bar that will form 50% martensite at the center when subjected to an “ideal” quench. Hardenability differences may be substantially greater for some case-hardening steel grades relative to others due to the difference in carbon content in the case and core. This is more critical for heavysectioned components that are reheated and quenched. The hardness gradient through the case is due to the relationship between the thermal gradient

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and the carbon gradient during quenching. Therefore, an increase in case hardenability required to produce greater amounts of martensite for a given carbon content will result in an increased case depth. In such cases, a reduced (shallower) carbon profile and shorter carburizing times will be necessary to obtain the desired hardness profile in the carburized component. SAE J1975 standard “Case Hardenability of Carburized Steels” summarizes characteristics of carburized steels and factors involved in controlling hardness, microstructure, and

residual stress. Methods of determining case hardenability are also provided. Parrish reported the following scheme that was developed to classify the case hardenability of steels (Ref 19):

Level 1: Surface carbon contents 40.8% C are martensitic. Level 2: All carbon contents from the surface to 50% C are martensitic. Level 3: All carbon contents from the surface to 0.27% C are martensitic. Level 4: A martensitic case occurs at all carbon levels, including the core material just beneath the case.

Figure 5 illustrates the core hardenabilities for a number of carburizing steels (Ref 19). This figure is used by estimating the equivalent diameter for the critically stressed section of the component of interest, and then the expected level of case hardenability of that steel is determined. Figure 5 indicates that level 4 is attainable only for small section sizes of more alloyed steels, and level 3, depending on the section size, is more readily attainable for most of the steels shown. Level 2 is more typical of the more common case-hardened parts and should represent a minimum target to be attained. Case hardenability may vary widely even for steels with equivalent core hardenabilities. Kern

Fig. 4

Jominy hardenability data for a number of carburizing steels

Fig. 5

Case hardenabilities of a number of carburizing steels with oil quenching

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and Suess provided the following guidelines (Ref 2):

Steel grades in which the case hardenability is due to carbide-forming metals such as chromium (8600 series) are sensitive to microcracking, especially when direct quenched from the carburizing temperature. This can be controlled by restricting the carbon content in the case to 0.9%. Steel grades with relatively high nickel content, for example, 4800 and 9300 series, may form excessive (430%) amounts of retained austenite when direct quenched unless the carbon content of the case is maintained at 50.75%. Carburizing round section sizes greater than 76 mm (3.0 in.) may lead to difficulty in achieving the desired case and core microstructures when quenching in oil. In such situations, consider induction hardening or nitriding or using a highly alloyed steel grade such as AISI 9310. Some standard grades of steel exhibit narrower core hardenability bands than other grades. For example, 8620H exhibits a hardenability band spread of 14 HRC at J 4, and 9310 exhibits only 8 HRC spread at the same J-value. This provides a greater amount of distortion control in addition to some possible application-dependent property advantages as well.

One problem that can arise during the steelmaking process or that may be observed as a millto-mill variant is the presence of segregation effects through the section of the steel billet during a continuous casting process, which results in the presence of a white band (Ref 20). White band is a type of negative segregation often observed in electromagnetically stirred continuous castings. The white banding produces a significant hardness gradient across the billet. After subsequent rolling and forging or machining to produce a component, the resulting grainflow can produce nonuniform hardenability and/or soft spots that can significantly affect distortion. In addition to proper hardenability selection, to achieve maximum core toughness, proper austenitization and quenching to martensite is necessary. These topics are discussed subsequently. Case Depth. The case of a carburized (or carbonitrided) steel alloy is that portion

extending inward from the surface, where the hardness is greater than that of the core. The total case depth is the distance or thickness of the carbon-enriched surface layer. The effective case is the point where 0.4 to 0.5% C (percent is called points in the industry) is present if the part is hardened to 50 HRC (510 HIV). The depth of the case is a function of carburizing time and carbon (carbon potential) at the surface. Genel and Demirkol have reported that the following equation model can be used to predict effective case depth (Ref 21): Effective case depth (mm)= 0:41 ½Carburizing time (h)1=2

The carbon potential of a furnace atmosphere at a specified temperature is defined as the carbon content of pure iron that is in thermodynamic equilibrium with the atmosphere. The carbon potential of the furnace atmosphere must be greater than the carbon potential of the surface of the workpieces for carburizing to occur. The carbon potential is a measure of the ability of a gas to react with the steel surface. It is this difference (carbon content in the gas and at the steel surface) that provides the driving force for carbon transfer to the parts being carburized. The composition of a gas that will produce a given surface carbon content is dependent on equilibrium data for the gas. The amount of carbon transferred will depend on factors that include temperature, time, and steel composition. The alloy composition of the steel will affect the effective carbon potential at the surface. The presence of elements such as chromium, manganese, and molybdenum that form stable carbides of iron will increase the effective carbon potential. Elements such as silicon and nickel form less stable carbides and reduce the effective carbon potential. Alloying elements that stabilize austenite or ferrite also reduce the effective carbon potential. The effect of alloying elements on carbon potential can be calculated from (Ref 22): Log (correction factors)=0:005 (%Si) 0:013 (%Mn) 0:040 (%Cr)+0:014 (%Ni) 0:013 (%Mo)

It is important to note that, except for long carburizing times (3 to 410 h), the surface carbon content is typically not equal to the carbon potential, because the surface of the part

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being carburized does not reach equilibrium with the carburizing atmosphere. This is illustrated in Fig. 6. A different, but similar, correlation exists for various steel alloys and carburizing conditions. From a correlation such as Fig. 6 and given case depth requirement and carburizing time, it is possible to correct for the carbon potential under nonequilibrium conditions to achieve a specific surface carbon content, as indicated by the dashed lines in the figure (Ref 22). For a specific temperature, the case depth (d) will vary with the square root of the carburizing time (t): d=Qt1=2

Values of the depth factor (Q) as a function of temperature are shown in Fig. 7. This equation is reported to be valid for low-carbon steel and some alloy steels (Ref 22). For applications such as automotive gears, typical case depths are 0.8 to 1.4 mm (Ref 23). An equation that is often used to calculate the case depth (in.) as a function of both time and temperature is the Harris equation (Ref 24): Case depth (in:)=½31:6 · t0:5 =10(6700=T)

where t is the time at the carburizing temperature in hours (carburizing time), and T is the

absolute temperature [Kelvin (K) = C+273, Rankine (R) = F+460]. In metric units, the case depth (mm) is (Ref 24): Case depth (mm)=660 · e8287=T · t0:5

At the operating temperature, the carburizing process may be conducted in two parts. Carburizing occurs during the first part of the process in a high-carbon-potential period when the enriching gas is added to the furnace atmosphere to increased the carbon content of austenite (the carburize-boost period) and the carbon potential is greater than the desired carbon potential. This part of the process is typically conducted at a carbon potential close to the solubility limit of carbon in austenite, typically between 1.0 and 1.2% C, which is dependent on the temperature and alloy content of the steel. The time for this part of the process to occur is called the carburizing time. This part of the process is followed by a boost-diffuse period, where the process is operated at the equilibrium carbon potential, which is reduced to a level that will maintain surface content, typically 0.8 to 0.9% C, during which time the carbon will diffuse deeper into the case and provide a gradual case/ core transition. Together, this is called the boostdiffuse cycle. The time for this part of the process is called the diffusion time. When the required case depth is achieved, if the component is direct quenched, the temperature is

Fig. 6

Carbon dioxide content of the atmosphere required to produce certain surface carbon levels at different carburizing times under a given set of carburizing conditions. The dashed lines illustrate alternative times and carbon dioxide contents to produce a single surface carbon content.

Fig. 7

Variation of the depth factor, Q, with carburizing temperature for low-carbon and certain alloy steels

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lowered to 850 C to reduce the distortion and then quenched (Ref 25). Carburizing boost-diffuse cycles conducted over 2 h are advantageous for case depths 40.50 mm (0.020 in). They are also useful when relatively deep cases free of carbides or retained austenite are required. This is important when the carbon content is greater than the eutectoid composition, where there is an increased tendency to form carbides and retained austenite upon quenching. These effects increase with alloy content (Ref 24). Harris also developed equations to compute the carburizing and diffusion times to achieve a specific case depth and surface carbon content (Ref 24): Carburizing time (h)=(C Ci )2 =(CO Ci ) Diffusion time (h)=Total time Carburizing time

where C is the final desired surface carbon content, CO is the surface carbon content at the end of the carburizing cycle, and Ci is the carbon content at the core. The effect of the steel alloy composition on the carbon gradient is illustrated for AISI 1020 plain carbon steel and AISI 8620 after carburizing at three temperatures in Fig. 8 (Ref 24). The alloy content will influence the diffusion rate, but its greatest effect is on the case carbon content. Normal carbon gradients, such as those shown in Fig. 8, can be achieved by maintaining a saturated austenite condition at the surface during the entire boost-diffuse carburizing cycle (Ref 24). It is important to control the ratio between the boost and diffuse times and to carefully control the carbon potential to avoid obtaining a carbon profile such as that shown in Fig. 9 (Ref 24). Although the desired surface hardness was obtained, the lower carbon content at the surface can lead to a transformation that proceeds simultaneously outward from the case-core interface and at the surface and proceeds inward such that the last portion of the case to transform is just below the surface. This will result in an undesirable condition where the surface is in tension relative to the core as well as a corresponding decrease in fatigue strength in addition to an increased potential for cracking (Ref 24). Boyer reports that a maximum tolerable carbon potential for carburizing cycles of up to 10 h at 925 C is 1.3% to avoid excessive soot formation.

If excessively high carburizing temperatures are used, the following situations may occur:

Rapid increase in grain growth and loss of properties Increased energy consumption Increased deterioration of the furnace fixtures and baskets

When high carbon potentials and long carburizing times are used to produce high surfacecarbon content and deep case depths, excessive retained austenite and/or free carbides may be obtained as a result. These microstructural products exhibit adverse effects on residual-stress distribution (which is discussed subsequently). Therefore, although high carbon potentials may be used for short carburizing times, substantial deleterious effects may result if used over prolonged carburizing times. Excessive carbon potentials, gaseous atmosphere composition control problems because of carbon probe malfunctions or air ingression, and inadequate furnace purging can lead to excess free carbon and sooting, which may be sufficiently severe as to leave carbon deposits on the parts (Ref 24, 26). This can lead to corresponding problems in controlling the carbon potential, resulting in nonuniform carburizing and dimensional control problems. Quenching. During carburizing, the steel microstructure consists of polycrystalline austenite. Grain coarsening may occur if the carburizing times are relatively long. The austenitic grain size will determine the size and distribution of martensite that will form as a result of quenching. In addition, phosphorus segregation into the grain boundaries may occur during the carburizing cycle, which has been found to be directly dependent on phosphorus and carbon content. Hyde et al. found that fatigue and fracture toughness decreased slightly when the phosphorus content increased from 0.005 to 0.017%, and when 0.017 to 0.031% P was present, the endurance limit and fracture toughness decreased substantially (Ref 27). Phosphorus also affects the degree of carbon segregation in the form of cementite at the austenite grain boundaries. During quenching, small amounts of cementite form at the austenite grain boundaries in the high-carbon case (Ref 25, 27). This leads to increased sensitivity to intergranular fracture, which is a major cause of fatigue crack initiation in carburized steels (Ref 25).

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Fig. 8

Carbon gradients for gas-carburized 1020 and 8620 steels. The 1020 steel was carburized in a batch furnace, and the 8620 was carburized in a pit furnace.

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After carburizing, the parts are then either quenched directly or air cooled and then reheated and quenched. Quenching is performed to harden the components. It is most desirable to develop a martensitic or bainitic case while controlling the formation of retained austenite to an acceptable level and simultaneously minimizing proeutectoid and pearlitic structures. The challenge is to quench sufficiently fast to produce the desired core structure but not so fast that the higher-carbon and more brittle case cracks. In addition, the desired hardness gradient between the surface and the core is critically

Fig. 9

Carbon profile of an incorrectly carburized steel

Fig. 10

important to achieve the desired wear and fatigue properties. If the carburized gear, such as a spiral bevel gear, is not quenched to achieve the necessary surface hardness and hardness gradient, failures accompanied by micropitting and, ultimately, fracture may occur (Ref 28). The morphology of martensite is carbon dependent, as shown in Fig. 10 (Ref 29). At lower carbon content, a lath martensitic structure forms, while plate martensite forms at higher levels of carbon. The two different morphologies are illustrated in Fig. 11 (Ref 30). Lath martensite exhibits better toughness than the higher-carbon plate martensite. Plate martensite, as the name indicates, forms as lenticular (lensshaped) crystals and is sometimes referred to as acicular (meaning needlelike) martensite or high-alloy martensite. A characteristic of plate martensite is the zigzag pattern of smaller plates, which formed later in the transformation, bounded by adjacent larger plates that formed in the beginning of the transformation (Ref 30). Typically, quenching is performed either directly from the carburizing process after furnace cooling to approximately the Accm temperature and then quenched, or the parts are air cooled and then reheated and quenched. Less commonly, double reheat quenching may be performed to provide high-durability components (Ref 19). Some quenching cycles recommended

Dependence of the martensitic structure on carbon content

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both the case and core grain size. The refined core will be soft and machineable with maximum toughness and resistance to impact. The refined case will be hardened for wear and resistance. In addition to grain refinement, double reheating and quenching is reported to improve fatigue properties by reducing the size and density of microcracks in the structure (Ref 21).

by Crucible Steel for carburized 8620 steel include (Ref 31):

Direct quench from carburizing: Quench in oil directly from the carburizing temperature of 925 C (1700 F). The core is hardened but unrefined. The case is hardened to the extent that it will be fileproof if the carbon content is sufficiently high. Cool, reheat and quench (1): After cooling from the carburizing temperature (925 C, or 1700 F), the carburized (but not yet hardened) steel is reheated to a temperature above the upper critical temperature, Accm, of the core (835 C, or 1535 F) and then quenched in oil. The core will be refined and exhibit maximum strength and hardness. The case will be hardened and somewhat coarsened. Cool, reheat, and quench (2): After cooling from the carburizing temperature (925 C, or 1700 F), the carburized (but not yet hardened) steel is reheated to a temperature above the lower critical temperature, Ac1, of the case (730 C, or 1350 F) and then quenched in oil to harden and refine the case. The core will be unrefined, soft, and machineable, and the case will be hardened. Double reheat and quench: The steel is cooled in the furnace from the carburizing temperature of 925 C (1700 F). The steel is then reheated to above 730 C (1350 F) and oil quenched to refine the core. The steel is again reheated to 730 C (1350 F) and oil quenched to refine the case. This double heating and quenching procedure refines

Fig. 11

Of these methods, the most common is direct quenching. However, there are a number of reasons why reheat quenching is favored for higheralloy, case-hardening steels, including (Ref 19):

To assure grain size and retained austenite control When intermediate subcritical heat treatment is required to condition the carbide structure within the case or to facilitate additional machining When the parts are to be plug or die quenched for distortion control When it is not possible to direct quench, such as in pit carburizing

Ingham and Clarke compared the results obtained for carburized 8620 steel with the carbon gradient shown in Fig. 12, which was oil quenched from a direct quench following carburizing and by a reheat and oil quench cycle (Ref 32). The results obtained showed that the direct quench process yielded a higher hardness than the reheat and quench process, which exhibited a relatively lower as-quenched hardness due to the presence of bainite in the hardened case structure.

Martensite morphology. (a) Lath martensite. (b) Plate martensite. Source: Ref 30

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Quenchants must be selected to provide cooling rates capable of producing an acceptable microstructure and hardness gradient through the case and the core. However, it is not desirable to use quenchants with excessively highheat-removal rates, since the propensity to cause increased distortion or cracking increases with quench severity. Although a reduction of quench severity leads to reduced distortion, it may also be accompanied by undesirable microstructures. Therefore, it is essential to select optimal quenchant and agitation conditions for the required microstructure, hardness, and strength in critical sections of the parts for each steel alloy, section size, and required microstructural and mechanical properties. Actual cooling rates or heat fluxes provided by a specific quenching medium are typically unavailable. However, some illustrative comparative data are provided in Table 4 (Ref 33). Figure 13 illustrates the comparative cooling properties of various oil-quenched steel bars assuming a surface heat-transfer coefficient of 0.019 cal s 1 C 1 cm2 (Ref 32). Quench nonuniformity is a significant contributor to quench cracking. Quench nonuniformity can arise from nonuniform flow fields around the part surface during the quench or nonuniform wetting of the surface. Both lead to nonuniform heat transfer during quenching,

Fig. 12

creating large thermal gradients between the core and the surface of the part. Poor agitation design is a major source of quench nonuniformity, since the purpose of the agitation system is not only to take hot fluid away from the surface and to the heat exchanger but also to provide uniform heat removal over the entire cooling surface of all of the parts throughout the load being quenched. A wide range of quench media can potentially be used when quenching carburized parts. Some comments on quench media selection, provided by Boyer, include (Ref 34):

For carbon steels, the most common quenchants are water and brine. When water is used as the quenchant, bath temperatures of 20 to 30 C with agitation are the most common. In the industry, oil quenchants are the most common, particularly when integral-quench (sealed-quench) furnaces are used at temperatures of 25 to 70 C. The quench oils may be classified as fast, intermediate, or slow depending on the cooling rate, enhancing additive, and quench oil base stock being used. When distortion control is critical, a hot oil that can be used at temperatures as high as 175 C may be used.

Comparison of direct quenching and reheat and quenching of 10 cm (4 in.) diameter AISI 8620 steel after oil quenching

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It is possible to use aqueous polymer quenchants, even in integral-quench furnaces, if appropriate structural conditions are met. The user is advised to consult his furnace manufacturer prior to use. A wide range of quench severities is possible by varying the polymer concentration, bath temperature, and agitation. In one study, it was shown that an aqueous polymer quenchant produced substantial improvements in fatigue properties relative to a conventional quench oil, which was attributed to an improvement in quench uniformity (Ref 35).

Table 4 Comparison of typical heat-transfer rates Maximum surface heat-transfer rate, W m 2 K1

Quench medium

Still air Nitrogen (1 bar) Salt bath or fluidized bed Nitrogen (10 bar) Helium (10 bar) Helium (20 bar) Still oil Hydrogen (20 bar) Circulated oil Hydrogen (40 bar) Circulated water

50–80 100–150 350–500 400–500 550–600 900–1000 1000–1500 1250–1350 1800–2200 2100–2300 3000–3500

When distortion control is especially critical, salt bath quenching may be required. However, parts should never be transferred directly from a carburizing bath containing 45% cyanide to a nitrate-nitrite quench bath, because this will result in a violent reaction and possibly an explosion (Ref 34).

One often-encountered quenching problem that may lead to increased dimensional control problems is contamination. For example, heterogeneous quench media caused by water contamination of oil or oil contamination of water or aqueous polymer solutions can potentially cause cracking problems. Similarly, salt contamination, either from salt baths or hardmetal ion contamination, can lead to problems of cooling rate control. Solid contamination, such as sludge or soot contamination in oil or aqueous media, also may lead to distortion and cracking. Finally, excessive foaming and air entrainment of the quench media will lead to nonuniform cooling, soft spots, increased residual stresses, and cracking. Therefore, it is essential that the quench bath be well maintained to assure optimal distortion control and minimize the potential for cracking. To develop the optimal residual-stress gradient, it is important to use the proper quenching

Soaking temperature

Bar size, in.

Fig. 13

Centerline cooling curves for oil-quenched steel bars of varying section sizes, assuming a surface heat-transfer coefficient of 0.019 cal s1 C 1 cm2

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conditions for the steel grade of interest. Figure 14 shows the development of the temperature distribution through 12.5 mm diameter bars carburized to a depth of 0.9 mm for oilquenched alloy steel and a water-quenched mild steel (Ref 36). This figure shows that the Ms and Mf temperatures are reduced with increasing carbon content in the case. The isochronal lines in the figure illustrate the cooling profile through the case to the core at specified time intervals. The martensitic transformation occurs at the core/case interface first. The case transforms last along with the corresponding expansion of martensite in the case. Since the core is already transformed, this process restrains further expansion during continued cooling, leading to the development of surface compressive stresses and placing the core in relative tension (Ref 36). As this figure shows, since the Mf is less than ambient temperature, retained austenite will accompany this process. Because carbonitriding is similar to carburizing, except that nitrogen and carbon diffusion into the case is involved, both processes exhibit similar microstructural transformation and evolution of residual stresses upon quenching. Process temperatures during carbonitriding are typically lower (800 to 850 C), as are process

times (30 to 60 min), which provides a relatively shallow case, usually 50.5 mm (Ref 36). These results show any factor that affects the cooling profile, such as bath temperature, agitation, or the quenchant selection, will exhibit a corresponding effect on the thermal distribution through the carburized case upon hardening and therefore on the development, type, and magnitude of residual stresses. To assure optimal distortion control, the following variables should be carefully monitored and controlled:

Adequate quality-control procedures of the quench media should be in place. For examples, follow ASTM D6710, “Standard Guide for Evaluation of Hydrocarbon-Based Quench Oils”, for oil quenchants and ASTM D6666, “Standard Guide for Evaluation of Aqueous Polymer Quenchants”, for aqueous polymer quenchants. Carefully control the water content if polymer quenchants are being used. Replace the quench mediate at regular uselevel intervals. Carefully control the quench bath temperature. Monitor fluid flow variation at critical locations in the quench tank. Nonuniform quenching has been reported to lead to quench cracking of a carburized 17CrNiMo6 axle in a reduction gearbox. Monitor hardness and dimensional changes of the parts being processed to look for unexpected variance.

Retained Austenite (Ref 1, 7). The quenchant temperature is a critically important variable in controlling the amount of retained austenite in the carburized steel. This is important because incomplete quenching and the presence of retained austenite will often seriously affect wear resistance and pitting fatigue strength (Ref 9). With carburized steels, the martensite start (Ms) temperature will decrease with increasing carbon content. To determine the impact of the carbon content on the Ms temperature from the steel composition, the Steven and Haynes equation may be used (Ref 37, 38): Ms ( C)= 561 474C 33Mn 17Ni 17Cr 21Mo

Fig. 14

Temperature distribution and martensitic transformation during quenching of carburized 12.5 mm diameter steel bar. The curves (isochronal lines) in the figure indicate time in seconds after immersion of the carburized (0.9 mm case) bar into the quenchant indicated.

where C, Mn, Ni, Cr, and Mo are the percent of the element contained in the steel. This equation is only accurate for steels containing up to 0.5%

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pg 192


C. For steels with higher carbon content, Fig. 15 should be used to determine a more accurate value for Ms (Ref 37). Although the degree of transformation between the Ms and martensite finish (Mf) temperatures is not linear, the difference is essentially constant (Mf is approximately 215 C lower than the Ms temperature) (Ref 37). Nevertheless, if Ms Mf is approximately constant, incomplete transformation can be expected if some part of the transformation occurs at a temperature lower than the quench bath temperature. Therefore, the volume of untransformed austenite (Vc) is related to both the Ms and the quenchant temperature (Tq). This relationship is quantitatively defined by the wellknown Koistenen and Marburger equation: 72

Vc =e71:10 · 10

(Ms 7T q )

Using these equations, Parrish demonstrated the effect of quenchant temperature on retained austenite on a hypothetical steel. The results of these calculations are summarized in Table 5, which show that the amount of retained austenite is expected to decrease with decreasing bath temperature. Trusova studied the formation of retained austenite after quenching and showed that the

retained austenite content was dependent on the carbon content and alloying elements and the content of the carburized case (Ref 39). The higher the steel temperature prior to quenching, the greater the decomposition of austenite but the greater the potential for cracking. To prevent cracking, steel may be quenched in hot oil, but the amount of retained austenite increases. Additional dilatometer examination of an isothermal high-temperature tempering process with subsequent cooling showed there was either a volume shrinkage due to decomposition of tetragonal martensite or an expansion caused by the decomposition of retained austenite (Fig. 16) (Ref 39). These transformations occurred during heating, isothermal holding, and subsequent cooling. As a result of this work, Trusova showed that reducing the carburized steel temperature to 800 C prior to quenching would reduce the potential for cracking and a double tempering at 580 to 600 C for carburized case structures containing i1.2% C. Because of the alloy and high carbon content in many case-hardened steels, the Ms temperatures in the carburized case are typically between 100 and 200 C or lower. These values will vary with the carbon content in the case. Therefore, the Mf temperature, which is approximately 215 C below the Ms, is also below the ambient temperature. Under these conditions, to reach the Mf temperature and therefore minimize retained austenite content, a subzero treatment is required. (This treatment is also known as refrigeration or deep cooling.) To minimize the possibility for the formation of subsurface microcracking, tempering at 150 to 175 C prior to cold treatment is commonly performed to stabilize the retained austenite. Gulyaev reported that the use of cold treatment to reduce retained austenite was most effective if conducted immediately after quenching, as shown in Table 6 (Ref 40). Equipment for achieving temperatures as low as 75 C may be relatively simple, such as dry ice mixed with kerosene or alcohol in a bucket. Temperatures down to 100 C can be Table 5 Effect of quenchant temperature on retained austenite Martensite start (Ms) temperature, °C

Fig. 15

Correlation curves for correcting the Steven and Haynes Ms temperature equation. When the carbon content is less than 0.9%, an 830 C soak of over 2 h is required to produce a fully austenitic structure.

150 150 150 150

Quenchant temperature (Tq), °C

Ms Tq, °C

Approximate retained austenite, %

Estimated hardness, HRC

80 60 40 20

70 90 110 130

45 35 29 25

52 56 57 58


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achieved by mechanical refrigeration. For lower temperatures, down to 195 C, liquid nitrogen can be used (Ref 41).

After cold treatment, the presence of retained austenite can be assessed by comparing the hardness before and after refrigeration. An increase in hardness is expected if retained austenite was present and was transformed to martensite as a result of the cold treatment. Since cold treatment increases martensite as a result of the loss of retained austenite, the carburized steel must be tempered a second time (150 to 200 C for 1 to 2 h) to reduce the potential for cracking. Table 7 shows the effect of subzero treatment after quenching to reduce the presence of retained austenite (Ref 42). Although subzero treatment of carburized parts does provide a reduction of retained austenite, the degree of transformation at a given temperature is variable, depending on the amount of retained austenite at the beginning of the subzero treatment, the elapsed time between quenching and subzero treatment, intermediate thermal treatments such as tempering, the level of compressive stress, any cold working of the material, and part design (Ref 41). However, it has been reported that fatigue resistance is decreased due to localized residual stresses imparted by the subzero treatment (Ref 36). The case ductility also seems to be negatively Table 6 Effect of time delay between quenching and cold treatment on retained austenite reduction Holding time at room temperature before subzero treatment at 183 °C

Fig. 16

Trusova dilatomer curves for tempering of carburized steels quenched from 950 C

Retained austenite, % Steel Kh-12

Steel SKh-12

24 46 ... 48

36.5 54.5 55 ...

2–3 min. 24 h 45 days 60 days

Table 7 Effect of subzero cooling after quenching Heat treatment after carburizing

Oil quenched from 800 C, low-temperature tempered Tempered at 650 C, oil quenched from 800 C, low-temperature tempered Air-cooled from 900–750 C, oil quenched low-temperature tempered

Bending strength

Impact strength

Condition

Retained austenite, %

Hardness, HRC

MPa

kg/mm2

MPa

kg/mm2

As-quenched Subzero treated

62 20

54 62

1530 1442

156 147

25.5 19.5

2.6 2.0


34 10

60 62

1697 1608

173 164

40 ...

4.1 ...


90 20

47 60

1618 1353

165 138

59 19.5

6.0 2.0

Note: Steel is 18Kh2N4VA. Subzero treatment conducted at 120 C.


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affected (Ref 43). In addition, subzero treatment increases the stability and induces microresidual stresses in any remaining retained austenite (Ref 43). Although subzero treatment does decrease retained austenite, in view of these complications, Parrish has suggested that optimization of the following process variables be considered before employing this process: quenchant temperature, surface carbon content, steel composition, and the use of reheat quenching (Ref 43). Increasing amounts of retained austenite will produce corresponding decreases in tensile strength, as shown by Fig. 17, although increasing strains can lead to the transformation of retained austenite to martensite (Ref 37). However, conflicting test results make it difficult to predict if the martensite formed by such straininduced transformations is beneficial or not, since one study reported by Parrish stated that the strain-induced martensite was more ductile, and another stated that the untempered martensite was more brittle. Koistinen showed that the distribution and magnitude of residual stresses in carburized steels was governed by the amount of retained austenite (Ref 44). Figure 18 illustrates the magnitude of residual stress as a function of the amount of retained austenite and position in the case for carburized SAE 8620 and 5140 and carbonitrided SAE 1118. The maximum compressive residual stress occurs at the position where the ratio of martensite/retained austenite is maximum (Ref 37).

Fig. 17

Dependence of stress for first detectable plastic strain (approx. 0.0001) on retained austenite content. AQ, air quenched; OQ, oil quenched; T, tempered

It is assumed, as Fig. 19 shows, that increasing the retained austenite content will reduce the low-stress, high-cycle fatigue limit of carburized steel (Ref 37, 45, 46). In a study using a flexural four-point bending fatigue test and carburized SAE 8620 steel test specimens, it was shown that increasing the

Fig. 18

Residual-stress distribution and retained austenite content in case-hardened steels

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retained austenite in the case resulted in longer fatigue life and that fatigue life is directly proportional to the square root of the grain size (d 0.5) (Ref 47). Jeddi et al. also showed that fatigue strength of carburized 14NiCr11 steel improved with increasing retained austenite content in the hardened case (Ref 48). In this work, it was shown that the level of improvement was related to the microstructure and residual-stress distribution within j200 mm of the surface. In addition, although the improvements in fatigue

Fig. 19

Fatigue limits of plasma- and gas-carburized test specimens as a function of retained austenite content

are attributable to uniform retained austenite content throughout the microstructure, only a maximum of approximately 40% of the retained austenite transformed to martensite at any depth due to the cyclic loading. The impact fatigue resistance is also dependent on the amount of retained austenite and the level of applied stress (Ref 37). Figure 20 shows that impact fatigue resistance actually increases with increasing retained austenite at the highest stress loading, while lowest stress loading exhibited the opposite effect. Increasing retained austenite content resulted in corresponding improvements of carburized SAE 8620 steel using an abrasive wear test utilizing a pin-on-disk tribometer, as shown in Fig. 21 (Ref 47). Sliding wear tests were conducted on carburized SNCM21, which corresponds to SAE 8620 steel, and carburized SCM4, which corresponds to AISI 4140. Although SCM4 is not typically carburized for the study reported, it was used to represent an example of high core strength and case depth. The results are shown in Fig. 22 (Ref 49). According to these results, sliding wear resistance increases with increasing retained austenite at a 40 kg applied load up to a critical retained austenite level, which, for this work, was approximately 30%, at which point the wear resistance decreased. The presence of retained austenite can also exhibit a dramatic effect on the scoring resistance of carburized steels. Kozlovskii studied the scoring resistance of carburized 20Kh2N4A (0.21% C, 0.62% Mn, 0.20% Si, 3.50% Ni,

Fig. 21 Fig. 20

Effect of retained austenite on impact fatigue resistance of a carburized 1.45C-11.5Cr steel

Effect of retained austenite (RA) on abrasive wear. Sample A, HRC = 59.7+1.8, RA = 37; sample B, HRC = 62.7+1.2, RA = 6%; and sample C, HRC = 61.4+1.5, RA = 23%

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pg 196


and propagation. Conversely, tensile stresses reduce desirable mechanical properties such as fatigue, fracture, and wear. Residual stresses are classified as the macrostress or residual stress of the first kind, which acts over a few grains, residual stresses of the second kind, and residual stresses of the third kind. Residual stress of the second kind is the difference between the average residual stress with a grain and the residual stress of the first kind. Residual stress of the third kind refers to stress variations within a grain. Residual stresses of the second and third kind are microstresses (Ref 51). In most engineering materials, such as steel, residual-stress variation between microstructural phases is typically more important than microstresses. The primary focus of this chapter is on macrostresses, residual stresses of the first kind. Carburizing and carbonitriding introduce surface and subsurface compressive stresses as a result of the formation of a carbon-enriched case. The increased carbon content in the case relative to the core significantly reduces the Ms temperature in the case, as illustrated in Fig. 23 (Ref 25). Because of the depressed Ms temperature, austenite-to-martensite transformation begins in the core before the surface, even though the surface temperature is lower. The volumetric expansion of the martensite in the core can be accommodated by the relatively hot untransformed austenite nearer the surface. Upon further cooling, the temperature at the surface is less than the Ms temperature of the carbon-enriched case, and it begins to transform to martensite. The martensite that formed first in the core is cooler and stronger than the austenite that is now transforming, and it resists the expansion of the higher-carbon-containing surface martensite now forming at the surface, which puts the surface in compression relative to the core. To maintain balance, the core is now in tension. This is illustrated by Fig. 23 (Ref 25).

1.42% Cr) using a roller machine test and found that even a relatively small amount of retained austenite could exhibit a large decrease in scoring resistance, even if the hardness was affected only minimally, as shown in Table 8 (Ref 50).

Residual Stress Residual stress is defined as a tensile or compressive force within a material such as steel without application of a thermal gradient or an external force. Residual stresses are produced by phase transformation, plastic deformation, or thermal effects such as contraction upon cooling. Newton’s laws require that compressive residual stresses at the surface of a material are balanced by tensile stresses within the material. Typically, compressive residual stress exhibits favorable effects such as improved fatigue life and stress corrosion by inhibiting crack initiation

Fig. 22

Sliding wear rate (at 200 rpm) as a function of retained austenite content. A, carburized SNCM21, 40 kg load, sliding distance of 864 m; B, carburized SCM4, 40 kg load, sliding distance of 864 m; C, carburized SNCM21, 20 kg load, sliding distance of 1728 m; D, carburized SCM4, 20 kg load, sliding distance of 1728 m

Table 8 Effect of retained austenite on scoring resistance of carburized and carbonitrided case Case carbon content, %

Case depth, mm

Surface hardness, HRC

On surface

At depth of 1.5 mm

Scoring load, kg/cm

Compressive stress at scoring load, kg/cm2

Carburize at 930 C for 8 h, slow cool to 750 C, quench in oil, temper at 180 C for 1.5 h

0.81

1.1

57–59

48

60

930

12,020

Carburize at 930 C for 8 h, slow cool to 20 C, temper at 600 C for 2 h, quench from 830 C in oil, temper at 180 C for 1.5 h

0.84

1.1

60–62

8

16

355

7424

Retained sustenite Treatment


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Fig. 23

Schematic illustration of carbon content, retained austenite, and residual stresses in the case of carburized steels

Fig. 24

An evaluation was made of 70 steels that were gas carburized, oil quenched, and tempered between 150 and 180 C. Case depths were 1 mm or less, and the core carbon contents varied between 0.15 and 0.20%. The range of residual-stress profiles obtained is shown in Fig. 24, where it can be observed that the compressive residual stresses acted throughout most of the case (Ref 25). It is important to consider the history of residual stresses due to dimensional changes that may occur during any stage of the manufacturing process and therefore may contribute to the final residual-stress condition of the heat treated component. These contributions to the stress history are illustrated in Fig. 25, which is a simulation of the production of a carburized 8620 steel cylinder that was subsequently quenched in unagitated water (Ref 52). The development of residual stress was achieved in the case of a 9.5 mm diameter chromium-molybdenum SCM420 (0.23% C, 0.72% Mn, 1.12% Cr, 0.21% Mo) steel test specimen that was gas carburized at 930 C until the case depth shown in Fig. 26 was achieved, oil quenched, and then tempered for 1 h at 200 C (Ref 53). These data show that the compressive residual stress increases with increasing case depth. Generally, the magnitude of the surface compressive stresses will be dependent on the ratio of the case and core thickness. When the core is thicker than the case, the surface

Range of residual stresses obtained for 70 carburized steels

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compressive stresses will be high. When the case is thicker than the core, the surface compressive stresses in the case will be lower, and the tensile stresses in the core will be higher (Ref 36). For automotive gears, the inversion point occurs where the surface compressive stresses become tensile, which was shown to occur at the point where the hardness becomes equivalent to the core value (Ref 54). Furthermore, the formation of surface compressive stresses was shown to be fundamental to the prevention of fatigue cracking. One of the primary reasons for conducting the carburizing process is to improve the fatigue performance of a heat treated component. A surface compressive stress will inhibit the formation and growth of surface cracks, which is important since fatigue failures are typically initiated at the surface (Ref 55). Kanetake, for example, showed an approximately 40% improvement in fatigue strength of the carburized SCM420 steel (Ref 53). Shot peening is also used to create residual surface compressive stresses that will increase fatigue properties of steel components (Ref 56). Therefore, it has been of continuing interest to examine the potential of fatigue property improvements not achievable by either treatment alone (Ref 55, 57–59). For example, Shaw et al. have reported up to a 75% increase in fatigue strength by

Fig. 25

combining carburizing and shot peening of gears prepared from gas-carburized 20MnCr5 steel (Ref 57). Before continuing this discussion, a brief overview of the shot peening process is in order. Shot peening (and shot blasting) involves impinging the surface to be treated with spherical media called shot. (Shot peening should be differentiated from shot blasting. Shot blasting a process in which an abrasive material is accelerated through a pressurized nozzle or centrifugal wheel and directed at the surface of a part to clean or otherwise prepare the part surface for further treatment, (Ref 60)). Shot peening is a cold working process where each individual spherical ball impinging the surface acts as a miniature hammer that plastically deforms and work hardens the surface by creating a small indentation upon impingement, as illustrated in Fig. 27 (Ref 61). The indentation process causes the surface to yield in tension. To balance the tensile forces involved in indentation, the subsurface is in a highly stressed compression state. As the process continues, the indentations overlap, and a uniform layer of metal is in residual compressive stress. The compressive stress is the result of superposition of residual stress formed by surrounding shots (Ref 62). The magnitude of the compressive stresses that are formed is material dependent

Illustration of the tangential stress history over the first 20, of a water-quenched 1.27 cm diameter 8620 carburized steel cylinder. The carbon gradient and retained austenite content are shown in Table 9.

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and is generally at least 1/2 the yield strength of the material (Ref 56). Some of the variables of the shot peening process include the shape and type of shot, size of shot, and impingement velocity. The effect of shot peening on contact fatigue under rolling-sliding conditions was studied using a carburized and hardened 20CrMnTi steel (0.21% C, 1.09% Cr, 1.0% Mn, 0.025% Ti, 0.02% P, 0.008% S), and, as a result of that work, it was shown that failure may produce three kinds of cracks, which were classified as: surface, which initiate from 0.015 to 0.05 mm; shallow sub surface; and deep subsurface, which initiate from 0.3 to 0.5 mm (Ref 63). Depending on the specific contact stresses on the

Table 9 Carbon gradient and retained austenite in the 8620 carburized steel related to Fig. 25 Depth below the surface mm

in.

Carbon, %

Retained austenite, %

0.0 0.5 1.0 1.5 2.0

0 0.020 0.040 0.060 0.080

1.20 1.10 0.80 0.40 0.25

32.0 28.0 15.0 4.0 0

Fig. 26 200

C

Residual-stress profiles of SCM420 steel that was gas carburized at 930 C, oil quenched, and tempered at

shot-peened surfaces, cracking, pitting, shallow spalling, and deep spalling may occur. In addition to providing substantial improvements in fatigue strength, carburizing and shot peening offer other benefits, such as reducing the deleterious effects of internal oxidation. However, in the absence of surface oxidation and oxide inclusions, MnS inclusions will then act as fatigue initiation sites (Ref 57). If a case-hardened surface is shot peened with sufficient intensity, a stress-induced transformation of retained austenite to martensite may be observed, and the surrounding volume constraint may result in a deepening of the surface compressive residual stress (Ref 55, 58). Nakonieczny et al. (Ref 58) have shown that plastic strain induced by shot peening reduces retained austenite in tempered martensite and produces a new e-phase (Fe2C and Fe3C). Peyrac, in another study, also showed significant retained austenite transformation and increases in residual stress as a result of shot peening of gascarburized 18NCD6 steel (Ref 64). Selected results from this study are summarized in Table 10. It is known that surface structure anomalies, also known as soft skin layers, which include internal oxides and nonmartensitic structures, including retained austenite (Ref 65) near the surface, will decrease fatigue properties of

Fig. 27

Illustration of the plastic deformation of the surface and resultant stress distribution after shot peening. Source: Ref 61

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gas-carburized steel (Ref 10). Structural surface anomalies act as preferential zones for fatigue crack initiation under low stress amplitudes. Shear-type crack growth can occur in ductile retained austenite in the near-surface region of the case for gas-carburized steels. Kikuchi et al. have shown that shot peening is very effective in improving the fatigue properties of carburized steels with these surface structure anomalies (Ref 59, 65). In fact, in their study, the fatigue properties were essentially comparable for a 20 mm diameter chromium-molybdenum steel (0.16% C, 0.26% Si, 0.74% Mn, 0.012% P, 0.013% S, 1.01% Cr, 0.18% Mo) bar that was carburized at 950 C for 1 h with and without surface structure anomalies (to a depth of approximately 30 mm) after shot peening. As a result of this work, it was concluded that since internal oxides near the surface can also act as preferential crack initiator sites, it is desirable to avoid the presence of such defects during gas carburizing (Ref 65). Although components such as gears are carburized and shot peened to introduce the desired level of surface compressive stresses, there are other processes that can be used to introduce compressive stresses. One such process is presetting. Presetting involves the introduction of an overload that causes yielding in the area of maximum stress concentration, such as in the root area of a gear tooth. When the load is released, a residual stress is introduced in that area. The induced stress is compressive on the side being loaded and tensile on the other side (Ref 66). Woods et al. evaluated presetting to improve the bending fatigue of carburized AISI 4120 steel spur gear teeth and found that presetting introduces compressive stresses in the area of a gear tooth where fatigue cracks originate. The results of this work showed that presetting provided substantially longer fatigue life.

Dimensional Stability Dimensional stability has two components: size (distortion) and shape (warpage). Distortion is defined as “an irreversible change in the component during heat treatment” (Ref 67). While changes in shape such as straightness (warpage) can be corrected by application of stress or by tempering in the elastic range (reversible), size changes are irreversible and cannot be changed in this way. Metallurgically, distortion may be thermally or transformationally derived. Size distortion typically refers to dimensional variation due to growth or shrinkage that is due to volumetric changes attributable to microstructural phase transformations (Ref 19, 68). Figure 28 shows the effect of the temperature dependence of the specific phase volume of different steel transformation phases (Ref 69). Variables that affect distortion include (Ref 19, 68, 70):

Chemical composition and hardenability— chemical and phase composition as well as hardenability (distortion increases as hardenability increases), as shown in Fig. 29 (Ref 67, 71, 72) Steelmaking—grain size and hardenability Hot working—hot reduction, length and direction of fiber Prior heat treatment—residual stresses, grain size and uniformity of microstructure Geometry—cheese blanks, shafts, rims Heat treatment aspects—heat rate, cooling rate (quench severity), quenching temperature, jigs and fixtures, plug quenching. Plug quenching is used to minimize dimensional change of inside diameters such as roundness and taper distortion of ring-shaped components (Ref 67, 68) (generally, salt bath quenching yields minimum distortion

Table 10 Effect of shot peening on retained austenite transformation and residual stress of gas-carburized 18NCD6 steel Heat treatment(b)

Depth modified by shot peening treatment, mm

Retained austenite on surface, %

Retained austenite converted to martensite, %

None

T1 T2

... ...

24.7 36.8

... ...

300 300

G1

T1 T2

100 100

5.7 14.9

19.0 21.9

1450 1350

G2

T1 T2

200 200

12.1 22.1

12.6 14.6

980 930

Shot peening(a)

smax, MPa

(a) G1, steel shot, BA 300, F 25–30A, overlap rate 150%; G2, steel shot, BA 800, F 55–80A, overlap rate 150%. (b) T1, carburize for 3 h. at 920 C, plateau at 850 C, oil quenching at 60 C, tempering for 2 h. at 150 C; T2, carburize for 3 h at 960 C, oil quenching at 60 C, tempering for 2 h at 150 C


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Fig. 28

Variation of the specific phase volume of different steel transformation phases as a function of temperature. Source: Ref 69

Fig. 29

Effect of steel hardenability on shape distortion

relative to hot oil) (Ref 67). Furnace temperature uniformity, case depth uniformity, prequench temperature, fluid flow during the quench (Ref 73), number of times a part is quenched, carburizing temperature (Ref 70)

Racking—Vashchuk et al. reported that high-temperature deformation of large gears could occur due to their own weight (Ref 73). They recommended the larger diameter of a gear wheel should be on the floor of the furnace (bottom) to support the end face and that the smaller diameter should be the free surface Machining—Parts should be machined as near final dimensions so that outer case will not require grinding after carburizing (Ref 24, 68). Residual stresses due to prior machining exhibit a large effect on distortion, and as shaved thickness increases, potential distortion increases (Ref 67) Method of green part manufacture, for example, parts machined from bar or tubing or forged from bar or tubing and then machined (Ref 70) Growth of surfaces during carburizing Murzin et al. reported that the cooling rate from the carburizing temperature to the prequench temperature exhibited one of the

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strongest effects on gear wheel distortion (Ref 74). Variables that affect warpage include (Ref 19, 68, 75):

Stresses resulting from phase transformations Nonuniform residual stresses in the original blank, such as those due to prior heating, including stress relief Nonuniform heating or cooling—furnace shape, part shape, heat control Insufficient furnace time—undersoaking Creep—method of stacking and fixturing of parts during heating and quenching (hanging versus standing) Internal stresses due to machining

Bulgakov reported that the most effective method of controlling stresses to reduce or stabilize warping was by controlling the hardenability of the steel (Ref 75). Hardenability control will permit control of the phase transformation and reduction of volumetric changes leading to structural stresses and warping. In particular, it was shown that slow cooling of carburized steels in the range of 1100 to 900 C

Fig. 30

after forging was extremely undesirable since this facilitates the formation of hard-to-dissolve particles and austenite grain growth during heat treatment. Optimal warpage control was achieved with grain sizes not greater than grade 7. A comparison was made of the effect of transformation behavior on dimensional stability of machine parts, such as gears, constructed from two different carburized steels: 20KhGR (0.18% C, 0.82% Mn, 0.24% Si, 1.07% Cr, 0.20% Ni, 0.0031% B) and 12KhN3A (0.14% C, 0.45% Mn, 0.21% Si, 0.78% Cr, 0.85% Ni). Since only very low cooling rates occur during carburizing, thermal stresses would be expected to be minimal, and any stresses that result would be due to austenitic transformation, which would be dependent on the steel chemical composition (Ref 76). The isothermal transformation diagrams for both steels before and after gas carburizing are shown in Fig. 30 (Ref 76). These diagrams show that austenite is less stable in the 20KhGR alloy and that relative stability of austenite compared to the 12KhN3A alloy remains after carburizing. Analysis of the kinetics of austenite transformation of alloys, before and after carburization

Isothermal transformation diagrams. (a) 20KhGR and (b) 12KhN3A alloys before carburizing (c) 20KhGR and (d) 12KhN3A alloys after carburizing. Source: Ref 76

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(Fig. 31) over the range of cooling rates used (0.79 to 175 C/min), showed that the noncarburized 20KhGR yielded mostly pearlite. The 12KhN3A alloy yielded predominantly bainite at the higher cooling rates, and at lower

Fig. 31

cooling rates, the pearlitic transformation predominated for both the carburized and uncarburized alloy. When air cooled, the case of the carburized 20KhGR consisted of pearlite and carbides, and

Dilatometric curves for the transformation of austenite in 20KhGR and 12KhN3A steels. Curves (a) and (b) were after carburizing, and curves (c) and (d) were before carburizing. (a,c) 20KhGR. (b,d) 12KhN3A. The cooling rates are: 1, 0.79; 2, 1.46; 3, 4.6; 4, 5.0; 5, 70; 6, 175 C/min. Source: Ref 76

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the core consisted of pearlite and ferrite. Upon rapid cooling, the carburized case had a small amount of martensite with a correspondingly lower amount of pearlite, and in the core, a lower amount of ferrite with the remainder being pearlite was obtained. However, for the 12KhN3A alloy, air cooling produced predominantly martensite and retained austenite in the case and correspondingly less pearlite. The core contained bainite (Ref 76). In practice, it was shown that gears manufactured from 20KhGB steel exhibited an average reduction in diameter (0.7 to 1.4 mm) from 463+0.3 after carburizing. This was attributed to reduced stability of austenite in the pearlite region. Conversely, for the 12KhN3A steel, the gears increased in size (0.2 to 0.5 mm) after carburizing because austenite was more stable in the pearlite region. These data suggest that distortion can be controlled by alloy and cooling rate selection, or improper selection of cooling rate or alloy can result in unacceptable component distortion (Ref 76).

Fig. 32

The effect of phase transformation behavior in the case and the core due to thermal gradients such as those occurring during heating and cooling will affect shape and size distortion of small parts, as illustrated by Fig. 32 (Ref 19). If the thermal and transformational stresses exceed the yield strength of the steel, corresponding distortion will occur.

Quenching and Grinding Cracks (Ref 77) Quenching cracks occur when tensile stresses of the first kind are greater than the material strength. Quenching cracks typically occur during, or in some cases after, quenching at temperatures less than the Ms temperature. Susceptibility to cracking increases with the carbon content of the steel, increasing austenitizing temperature, and cooling rate, especially in the Ms Mf transformation temperature range. The probability of cracking during quenching increases with the presence of stress raisers

Resulting distortion after heat treatment of different steels after quenching in oil and water

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such as notches, abrupt changes in section size, keyways, and holes. Quenching cracks are irreversible but can be minimized by appropriate design modifications, reduction in hardening temperature, and use of lower quench severity. The reasons for the occurrence of quenching defects include:

Poor surface cleanliness, such as residual forging and metalworking lubricants, and so on, which leads to nonuniform quenching, increased thermal gradients, and soft spots Incorrect loading and arrangement of parts in the furnace, which leads to nonuniform heating and related distortion Excessive heating rates, which may lead to warping and cracking Lack of a protective atmosphere to eliminate oxidation and decarburization of the steel surface, which will lead to reduction in mechanical properties after hardening and a decrease of hardness of the superficial layer Excessive cooling rates and incorrect immersion into the quenching bath, causing cracking, warping, and twisting Insufficient cooling rates or undersized quench tanks, which will inhibit the desired martensitic transformation

Quenching cracks, which are characterized by relatively large depth and short length, rarely occur in the case of carburized or carbonitrided components. This can be explained by more

Fig. 33

beneficial compressive surface stresses than those typically formed in higher-carboncontaining through-hardened parts. However, internal cracks may form in the core below the carburized surface or in the transition zone, that is, in the places with the largest tensile stresses in carburized parts. Cracks may also form on the surfaces and corners of carburized parts, which is related to triaxial tensile stresses in these locations. Therefore, to prevent cracking, the case must be sufficiently deep so that stresses developed at any point below the surface are less than the fatigue limit of the material at that point (Ref 1). When steel contains greater than 0.5% C in a martensite matrix, such as in the carburized case, intergranular fracture along prior-austenite grain boundaries may occur. In this situation, the intergranular fracture is due to the presence of both phosphorus and cementite formation on the austenite grain boundaries during austenitizing or cooling from austenitizing temperatures. Krauss has referred to this fracture mechanism as quench embrittlement and has suggested that the mechanism for this to occur is analogous to quench cracking in through-hardened steels, which is due to the formation of tensile surface stresses during quenching, as described previously (Fig. 33) (Ref 78). However, since relatively high surface compressive stresses are present in properly carburized steels, quench cracking should not exist if conditions for potential intergranular cracking are present.

Scanning electron micrographs of overload case fracture surfaces in carburized SAE 8620 steel. (a) Quenched directly after carburizing at 927 C (1700 F). (b) Reheated to 788 C (1450 F). Both specimens were tempered at 145 C (300 F).

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Krauss has reported that carburized steels will fail by intergranular cracking if sufficent bending or tensile stresses are applied to offset the compressive stresses in carburized cases. The fracture map shown in Fig. 34 illustrates three conditions where susceptibility to intergranular cracking can be minimized: carburizing, intercritical austenitizing, and applications where loading is Hertzian or compressive (Ref 79). In some cases, microcracking can occur with higher-carbon lath martensite matrices, which are present in the hardened carburized case. Microcracking is due to contact of martensitic plates with each other or with the austenitic grain boundaries. The potential for microcracking increases with the austenization temperature. Prior-austenite grain size also affects microcrack density, which decreases with decreasing prior-austenite grain size. However, microcrack density is not affected by quench severity (Ref 78, 80). Figure 35 illustrates examples of microcrack formation in the carburized case of SAE 8620 steel (Ref 78). The presence of microcracks can further lead to a surface defect called flaking, which refers to

Fig. 34

flakelike fracture and subsequent peeloff. (Note: Flaking is initiated at microcracks that may also be caused by surface damage due to lubricant contamination by chips, shavings, burrs, or abrasive powder ingression into the lubricating system.) Krauss reported that fatigue resistance decreased with increasing microcrack formation. Crack initiation occurred at the site of the microcrack that acts as a stress concentrator (Ref 78). Although microcracks can be removed by surface grinding and polishing, fatigue failure may still be initiated at prior-austenite grain boundaries, intergranular surface oxides, or surface defects such as scratches, machining marks, and surface asperities due to roughness. Aksenova et al. showed that contributing factors to cracking of case-hardened gear wheels included residual tensile stresses in the case-core interface, grain growth and overheating, supersaturation of the case with carbon, excessively high cooling rates during quenching, too low a quenching temperature, and insufficient residence time in the quench tank after immersion (Ref 81).

Krauss fracture map illustrating conditions where susceptibility to intergranular cracking can be minimized: carburizing, intercritical austenitizing, and Hertzian or compressive contact loading

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McEvily et al., reported the use of fractographic analysis to explain cracking of a carburized AISI 9310 gear at the case-core transition zone. One of the factors was the differential in the Poisson ratio (n) between the case and the core that develops during a monotonic bending test when the core deforms plastically while the case deforms elastically. McEvily used as an example to illustrate this point a 9 mm diameter round bar with a 1 mm case. When the bar is first loaded initially and the deformation is elastic, the Poisson ratio for the case and the core is equal. Upon further bending, the core will deform plastically while the case is deforming elastically. However, because of the difference in the Poisson ratio between the case (n = 0.3) and the core (n = 0.5) at this point, radial tensile stresses will develop. For a tensile strain of 1%, the radial strain in the case (relative to the centerline) would be 0.003, and the radial strain in the core would be 0.005. To remain compatible, the difference in the strain and corresponding displacements must be accommodated by the formation of a tensile stress. However, this radial stress will result in a state of triaxial stress that will promote brittle behavior by inhibiting plastic deformation. The total tensile stress may then be developed sufficient to result in rupture at the case-core interface (Ref 82). Grinding Cracks. Grinding may be used for postprocessing of components to remove growth and distortion that may have resulted from carburizing and carbonitriding. Grinding is also performed to remove such metallurgical features as carbide films, internal oxidation, and hightemperature transformation products that may impart deleterious performance properties. In addition, grinding processes are commonly used

Fig. 35

Microcracks in the martensitic case of a coarsegrained SAE 8620 steel

to create the desired surface finish to improve bending and contact fatigue and lubrication properties. Surface cracks in the carburized case may occur during the grinding process, which can be attributed to microstructural transformations and thermal stresses producing tensile forces. If the tensile forces in the case exceed the material strength, then surface cracks will result. This is due to the difference in the specific volume of the transformational phases present in the case structure, primarily martensite and austenite. Structural defects, such as those caused by inclusions, will also influence the susceptibility of the steel to cracking. Generally, crack creation during grinding is influenced by thermoelastic tensile stresses that are created in the surface cooling zone during grinding. They are dependent on thermal and mechanical properties of the material, maximum contact temperature, grinding feed depth, and cooling rate. One study conducted on the generation of grinding cracks showed that microstructural heterogeneity, such as the presence of carbide inclusions (particularly those with a mean diameter of 6 to 10 mm), which were shown to be associated with large internal residual stresses, were a predominant cause of grinding cracks (Ref 83). Grinding burns arise when excessive heat is generated during the grinding process, and this is characterized by surface discoloration. The term grinding burn refers to localized surface temperature increases at least sufficient to cause tempering of the martensitic surface, resulting in localized soft spots. Furthermore, since carbide precipitation volume contraction accompanies tempering, the burnt areas are in tension and, if the resulting tensile stresses are sufficient, subject to tranverse cracking. However, in other cases, the increase may be in excess of the Ac3 temperature, producing an austenitic surface that, upon rapid cooling, may produce a hard, light-etching, martensitic thin layer at the surface. This induced defect is known as a rehardening burn, which is characteristically surrounded by a layer of tempered steel (Fig. 36) (Ref 84). In this case, the rehardened zone is in compression due to the martensitic volume expansion, and the surrounding areas of tempered martensite are in relative tension. Cracking may occur in the area surrounding the rehardened material or in the interface between the two (Ref 84).

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There are two possibilities to prevent the occurrence of thermal defects in the surface resulting from the grinding process:

Reduction of heating due to the rotation of a grinding wheel by reducing the speed of the grinding wheel will reduce heating and friction. Conversely, increasing the speed of the grinding wheel will increase heat production and the potential for the formation of grinding burns due to overheating. Increasing heat abstraction due to the grinding process and reducing the contact time between the grinding wheel and part will reduce the tendency for grinding burns.

Grinding cracks may exhibit characteristic short, parallel cracks (Fig. 37), or they may exhibit a “chicken-wire” pattern and are typically between 0.076 and 0.13 mm (0.003 and 0.005 in.) deep. The parallel cracks are typically deeper than the chicken-wire pattern. Grinding cracks form perpendicular to the grinding direction. As indicated previously, the potential for grinding cracks is affected by improper heat treatment or a metallurgical structure that is prone to cracking. For example, if the surface temperature exceeds the Ac3 temperature, the steel in this region may transform to austenite, then upon rapid cooling, a hard martensite layer may form. This effect is called a rehardening burn. Grinding cracks may be detected by a magnetic particle test. Some carburizing steels such as chromium and chromium-manganese steels, which include SAE 5120 and 20MnCr5, may undergo overcarburizing with subsequent cracking of the case

Fig. 36

Microstructure of a section through a rehardening burn. Original magnification: 500 ·

upon cooling, which then renders the part more susceptible to the formation of grinding cracks (Ref 1). Overcarburizing leads to the formation of a complex carbide network that is excessively brittle, which causes greater susceptibility to grinding cracks. Severe grinding may lead to the development of residual tensile stress, which can be the initiation point for crack formation. Since cracks will not propagate into layers of compressed stress, it therefore may be advantageous to shot peen the part prior to grinding to prevent the formation of grinding cracks (Ref 86). Parrish reported that the potential for formation of grinding cracks may be minimized by (Ref 84):

The thermal conductivity of the steel is an important design variable, and free carbides and retained austenite have an adverse effect on thermal conductivity. The surface carbon concentration should be between 0.7 and 0.9%. Parts should be tempered immediately after quenching. The tempering temperature should be as high as possible while still achieving the necessary surface hardness.

Improper Case Depth (Ref 77). In a recent study conducted by Bahnsen et al. on carburized SAE 5120 test specimens to rate the relative influence of surface carbon content, case depth, and carburizing temperature on distortion, it was

Fig. 37

Example of grinding cracks on the flank of a worm gear. Source: Ref 85

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reported that of these variables, the most dominant effect was observed for case depth (Ref 87). Case depth of carburized steel is determined by the carburizing time and the available carbon potential at the surface. One of the most common defects of carburized and carbonitrided materials is an insufficient or excessive case depth. For example, when prolonged carburizing times are used to produce a deep case depth, a high carbon potential will produce a high carbon content on the surface and the possible corresponding formation of excessive retained austenite or free carbides, which may lead to an improper residual-stress distribution in the casehardened part. Therefore, a high carbon potential may be suitable for short carburizing times and shallow case depths but not for prolonged carburizing times and deep case depths. Furthermore, the fatigue limit of carburized SAE 8620 steel was related to case depth and also microstructure, distribution of retained austenite, depth of internal oxidation, and nearsurface compressive residual stresses (Ref 10, 21). Bending-fatigue strength decreased with the increasing case depths due to the presence of increasing internal oxidation and nonmartensitic transformation products at the surface. Wear behavior of carburized 8620 steel is also related to case depth (Ref 88). To assure optimal quality during the production of case-carburized parts, the following are essential:

The processing temperature should be accurately controlled. Maintain temperature uniformity throughout the load. Rack the parts to assure uniform gas flow throughout the load. Use uniform circulation of atmosphere throughout the load in the furnace To properly control case depth, either use shim stock or sample the parts periodically during the carburizing cycle. Conduct the process at the lowest acceptable temperature and time. The components constituting a load should possess uniform size and surface area with respect to each other, using empirically established carburizing conditions. Avoid carburizing and carbonitriding in the same furnace. Use a separate furnace for each process, if possible.

The optimal case depth for a specific component and steel alloy is based on the design and

service conditions of the component. Typically, the case depth is designed to provide the necessary residual-stress distribution for the wear requirements for the part (Ref 89). Typically, the greater the case depth, the greater the fatigue strength (Ref 23). For carburized steel gears used in the automotive industry, for example, SAE 8620, hardened case depths are generally 0.8 to 1.4 mm. Improper case depth may be caused by establishing an unnecessarily restrictive case thickness specification that is not appropriate for the process or the particular furnace in use, which leads to decreased and nonuniform hardness and unacceptable material properties.

Insufficient Case Hardness and Improper Core Hardness (Ref 77) One reason for insufficient case hardness is the presence of incorrect microstructure, such as bainite. The appearance of bainite in a carburized case in even small amounts will significantly decrease fatigue strength, especially contact fatigue strength. It is an important microstructural defect to be avoided. The presence of bainite in the carburized case is particularly problematic because it cannot be detected by hardness measurements and by the severity of the quenchant used to harden the steel after carburizing. Figure 38(a) illustrates a bainitic case microstructure of carburized SAE 8620. Figure 38(b) shows the core structure. In the case of carbonitrided components, this is less important because nitrogen increases hardenability of steel more than carbon. Insufficient as-quenched case hardness is caused by:

Insufficient carbon content in the entire case or in the superficial zones Increased retained austenite content Insufficient case hardenability Insufficient case depth

On the other hand, core hardness is dependent on carbon content, steel alloy hardenability, and section size. Case carbon content less than 0.4 to 0.5% C is easily detected by conducting a spark test. Insufficient case carbon content resulting from the gas carburizing process occurs when the process is conducted with a low carbon potential, inadequate furnace pressure, or cooling the

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load from the diffusion temperature to the quenching temperature without a protective atmosphere or improper atmosphere composition. These problems may be rectified by additional soaking in a carburizing atmosphere with the proper carbon potential. Hardness reduction due to an increase in retained austenite content may occur with a direct quenching after the carburizing process. This is even more critical with steels containing increased chromium or nickel, which may produce retained austenite levels as high as 80 to 100% in the surface after quenching. As the case thickness increases, the zone containing retained austenite may be sufficient to cause a significant reduction in hardness. To reduce the retained austenite content from a direct quenching process, it is important to select the proper steel alloy for carburizing and to use the appropriate quenchant and quenchant conditions, including the use of a subzero treatment, if necessary. Reduced hardenability usually does not occur throughout the entire case but only in the 0.001 to 0.01 mm depth from the surface. Common problems leading to reduced hardenability include internal oxidation and overcarburization of surface zones. To improve case hardenability, internal oxidation and overcarburizing should be prevented. When the case depth is too shallow, the observed hardness is dependent on the load applied. For example, for thicknesses of 0.3 to 0.4 mm, if the surface hardness is determined using a 60 kg load, the value will be

approximately 12 HRC units higher than if the hardness is determined using a 150 kg load. In practice, it is possible for the hardness to differ by 1 to 2 HRC units from published values for the material. Hardnesses that are i2 HRC units less than published values may be due to the use of an incorrect steel grade or an insufficient austenitizing temperature for the steel hardenability and section size in use. Hardness values iHRC units higher than published values may be due to the use of an incorrect steel grade, excessive case depths, or puncturing a copper layer or paste on surfaces that are protected from carburization. Core Microstructure. The design material properties of case-hardened steels are not only dependent on a martensitic case but also on the microstructural composition of the core. An important design criterion is the ultimate tensile strength, which is dependent on the microstructure of the core. For example, soft cores (5770 N/mm2, or 50 ton/in.2) are suggestive of a core with high ferrite content, as shown in Fig. 39(c) (Ref 19), and a hard core (41240 N/ mm2, or 80 ton/in.2) would be expected for a predominantly martensitic structure, as shown in Fig. 39(a) (Ref 19). Intermediate structures would be bainitic structures, such as those illustrated in Fig. 39(b) (Ref 19). The effect of core microstructure on ultimate tensile strength is illustrated by Fig. 40 (Ref 90). The approximate relationship between the core microstructure and hardness for a Ni-Cr-Mo steel is illustrated in Fig. 41 (Ref 19).

The case and core microstructure of carburized SAE 8620 test specimens (0.95% C potential); carburized at 955 C (1750 F), quenched into a 50:50 mixture of sodium nitrate and potassium nitrate at 250 C (480 F), held 120 min, then air cooled and tempered at 250 C (480 F) for 240 min. The case is lower bainite (56 HRC), and the core (42 HRC) is lath martensite. These images were made by etching with 10% Na2S2O5 (sodium metabisulfite). The magnification bar is 20 mm. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

Fig. 38

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Influence of the Transition Zone. There is a transition zone between the case and the core, and the thickness of a transition zone is dependent on carburizing time (transition zone thickness increases with carburizing time), carburizing medium, and carburizing temperature (excessive carburizing temperature will lead to increased pearlite content in the core). After subsequent hardening, if there is too great a transition between the case and the core, there is increased potential for peeling and chipping, resulting from the presence of a martensitic carburized case and an uncarburized core containing sorbite. Fatigue cracks occur most often in the transition zone, which subsequently propagate into the core and into the carburized case. When the gradient between the case and the core changes too rapidly, operations such as grinding or when a component is subjected to bending due to heavy loadings, could lead to peeling and chipping failures. For example, deep grinding could not be performed. Increasing the depth of the transition zone will increase the strength of adhesion of the case to the core. Then, if the transition zone is sufficiently large, deep grinding operations may be performed. Figure 42 provides the microstructure of two transition zones. Figure 42(a) illustrates a gradual transition between the case-core microstructure. A more rapid transition between the case-core microstructure is illustrated in Fig. 42(b). Typically, a transition zone such as that illustrated in Fig. 42(a) is desired, since it will exhibit a lesser tendency for chipping.

Influence of Surface Carbon Content

Fig. 39

Microstructures obtained by cooling a 0.16%C3%Ni-Cr steel from 920 C. (a) Fast cool (920– 200 C in 30 s), giving low-carbon martensitic structure of 1590 MPa ultimate tensile strength (UTS). Original magnification: 800·. (b) Intermediate cooling (920–250 C in 200 s), giving bainitic structure of 1360 MPa UTS. Original magnification: 800·. (c) Slow cool (920–250 C in 104 s), giving a ferrite/ pearlite structure of 740 MPa UTS. Original magnification: 800·

Overcarburizing or Overcarbonitriding. Important microstructural defects related to carburized or carbonitrided case structure include overcarburization or overcarbonitriding of the case and coarse grain structure. Excessive carbon content (carburizing) or carbon and nitrogen (carbonitriding) is typified by the presence of carbides or carbonitrides in the case, which creates an almost continuous nonetching area. Case-hardened steels with these microstructures characteristically exhibit increased brittleness, a propensity for chipping during grinding and use, and decreased fatigue strength and pitting resistance.

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Fig. 40

Fig. 41

Approximate effect of microstructure on the ultimate tensile strength of low-carbon, low-alloy steels

Approximate relationship between core microstructure and hardness of a Ni-Cr-Mo carburizing steel (approximately 4% alloy content) with approximately 0.16% C. The alloy content/carbon content extension (upper right corner of the figure) permits phase percentage plots to be adjusted in relation to the fixed hardness scale to approximate core strength for other steels. Below 250 HV represents slow-cooled (normalized) and annealed steels, and bainite can be read as bainite, pearlite, or spheroidized carbides. Above 250 HV refers to quenched steels. For the 180 C tempered condition, there will be zero change at 360 HV and below, but there will be a 20 HV loss at 100% martensite.

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Figure 43 illustrates the effect of carbon content on the hardness of martensite in carbon and alloy steels. Increasing carbon content to 0.5% increases hardness from 20 to 65 HRC. However, increases in carbon content to approximately 1% do not produce a corresponding increase in hardness above 65 HRC. Hardness and mechanical properties are related not only to carbon content but also the composition of carburized steels, which is illustrated in Fig. 44. This figure shows the corresponding carbon gradients of the cases of three carburized steels: chromium-molybdenum, carbon, and nickel (Ref 25). The data in Fig. 44 show that the presence of chrome and molybdenum increases the case carbon content, and nickel decreases the case carbon content. The case carbon content is increased due to the presence of carbide-forming elements; their structures in the carburized case influence the mechanical properties of the steel. Overcarburization may also lead to quench cracking (Ref 91). The influence of steel composition on the microstructure of a carburized and quenched 20H steel is shown in Fig. 45. Figure 45(a) shows a martensite-retained austenite microstructure with some carbides dispersed throughout the

Fig. 42

structure. Fig. 45(b) shows a martensiticretained austenite structure with network carbides. This structure is due to the chrome content in the steel and leads to varying carbon content during carburizing. Additional examples of the case microstructure of SAE 8620 steel and core microstructures of SAE 1524 and 8115 steel are shown in Fig. 46 to 48 respectively. Decarburization is the opposite of carburizing. While carburization is performed to increase carbon content in the surface of steel, decarburization is the process by which carbon is lost from the surface of steel. Decarburization can lead to catastrophic failures of components (Ref 92) and must be minimized because of fatigue failure such as bending and contact fatigue (Ref 93). Figure 49 illustrates decarburization of a poorly carburized SAE 8620 steel, and Fig. 50 shows the microstructure at higher magnification. Decarburization occurs at temperatures in excess of 700 C in the presence of gases that act as decarburization agents, which include carbon dioxide (CO2), water vapor (H2O), hydrogen (H2), and oxygen (O2). The decarburization process involves the following chemical reactions of molecules with carbon in the steel surface (CFe) until there is an equilibrium

Micrographs illustrating transition zones between the carburized case and the uncarburized core. (a) Illustrates a gradual transition. (b) Illustrates a rapid transition between the case and the core microstructures. Original magnification: 500·

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Fig. 43

Fig. 44

Martensitic hardness as a function of carbon content in carbon and alloy steel. Source: Ref 80

Dependence of the carbon gradient as a function of case depth for three carburized steels that were carburized under the same conditions: 925 C and 10 h. 1, chromium-molybdenum steel (0.56% Cr, 0.16% Mo); 2, carbon steel; 3, nickel steel (3.5% Ni)

Fig. 45

Microstructures of the carburized case structure of two different samples of 20H steel that were carburized in the same load. Carburizing temperature: 930 C for 7 h; hardening temperature: 860 C for 0.5 h. (a) Surface. (b) Core. Etchant: 3% HNO3. Original magnification: 500 ·

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Fig. 46

Microstructure of SAE 8620 case of a mold taken just below the surface. Etchant: alkaline sodium picrate boiling (60 s), area just below the surface. Original magnification: 500·. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

Fig. 47

Illustration of lath martensite in the core of carburized SAE 1524 steel; water quenched from 925 C (1700 F). Etchant: nital. Scale = 10 mm. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

Fig. 49

Micrograph of poorly carburized SAE 8620 mold showing decarburization at the surface (note patches of ferrite and pearlite). Below this zone is where the grainboundary carbides are seen. Original magnification: 500 ·. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

Fig. 50

This is a higher magnification of the decarburized microstructure shown in Fig. 49 of the surface of a poorly carburized SAE 8620 mold (note patches of ferrite and pearlite). Original magnification: 1000 ·. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

established between the gaseous atmosphere and the steel surface: CFe CFe +H2 O CFe +2H2

Fig. 48

Illustration of lath martensite in the core of carburized SAE 8115 steel; water quenched from 925 C (1700 Etchant: nital. Scale = 10 mm. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL F).

2CO CO+H2 CH4

When the reactions proceed from left to right, decarburization will occur. These are the reverse of the carburization process. Some of the more commonly reported causes of decarburization include a malfunctioning endogas generator, such as soot accumulation hindering catalyst activity; excess moisture in

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the furnace atmosphere; air contamination and leakage (Fig. 51); heating in aged (deoxidized) salt baths; and improper carbon potential selection (Ref 94). If a carburized steel is at a temperature above the Ac3 (approximately 900 C) in a decarburizing atmosphere, the carbon potential will be low and the surface carbon content will also be low, since the carbon in the steel surface and the related gaseous reactions shown previously will be driven to this equilibrium condition. This will result in a decarburized layer being produced, and the depth of the decarburized layer will depend on residence time in the furnace under these conditions. If the temperature of the steel is below Ac3 and above Ac1 (800 to 840 C), there is a different decarburization condition. In this case, the carbon content rapidly decreases from “A” to “B,” as shown in Fig. 52 (Ref 94). Further decreases in carbon content will result in a material of carbon content “C” in equilibrium with material of carbon content “B”. Therefore, further loss of carbon by decarburization must result in the formation of ferrite containing carbon content “C”. If the atmosphere carbon content is controlled to carbon potential “D”, then ferrite cannot form. Instead, a gradient is formed between carbon contents “A” and “D”.

Fig. 51

Effect of air ingression into the carburization atmosphere (N2/4% natural gas) on the decarburization of SAE 8620 after 2 h at 850 C. Source: Ref 94

Decarburization is typically classified as total or partial. Figure 53(a) illustrates a case of total decarburization of 1018 steel and is characterized by a ferritic surface layer (Ref 94). Usually, there is a gradient from total to partial decarburization with increasing depth from the surface. Partial decarburization is illustrated in Fig. 53(b) and is often characterized by grainboundary ferrite at the surface (Ref 94). In this case, nital etching will reveal a structure more gray in color than would be achieved with a higher-carbon martensite. The quenched surface of partially decarburized steel is typically bainitic or martensitic. The final microstructure produced by a specific level of decarburization is dependent on the steel alloy and cooling rate. Decarburization is accompanied by surface hardness reduction. However, partial or shallow decarburization may not necessarily be detected by macrohardness determination. If decarburization is suspected, microhardness determinations, in view of their sensitivity to the presence of such microstructures, should be used. Decarburization can exhibit dramatic effects on the residual stress of a component, as illustrated by Fig. 54 (Ref 94). In this example, the surface carbon content of a carburized 3.5Ni-1.5Cr steel would be approximately 1%, and the surface residual stresses would be compressive at 4392 MPa. After decarburization reduces the carbon content to 0.64%, the surface residual stresses were found to be nearly 0. Finally, when decarburization reduced the surface carbon content to 0.35%,

Fig. 52

Iron-carbon equilibrium diagram to explain decarburization

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the surface residual stresses were tensile at 226 MPa. Various reports have shown that decarburization can result in large decreases in bending and contact fatigue strength (Ref 94). Since wear-resistance properties are typically dependent on achieving optimal hardness, reduction of carbon content and thus surface hardness will have a correspondingly adverse effect.

Fig. 53

Micrographs illustrating total and partial decarburization. (a) Total decarburization of 1018 steel caused by a furnace air leak. Etchant: 1% nital. Original magnification: 500·. (b) Illustration of partial decarburization. Original magnification: 190 ·

Influence of Grain Size Grain size is one of the most important and characteristic features of steel. Grain size influences mechanical and plastic properties, especially impact resistance and also steel hardenability. Grain size is characterized by the size of the austenite grain, and it is dependent on various factors, such as degree of cooling and the deoxidation process during steelmaking (Ref 95). Generally, an ASTM grain size of 6 to 8 is specified. Figure 55 illustrates grain sizes of ASTM No. 6 to 9. Grain growth is affected by temperature. Typically, grain growth increases with temperature and time at temperature. Aluminum may be added to steel to provide resistance to grain growth (grain refined). Alloying elements such as nickel and molybdenum also provide greater resistance to grain coarsening at typical carburizing temperatures than plain carbon steels. Although grain coarsening is usually not a problem for carburizing temperatures up to 925 C, carburizing at temperatures greater than 1000 C is typically accompanied by some grain coarsening, yielding a mixed-grain structure. Reheating at 820 to 860 C can be performed to refine the mixed-grain structure. Fine-grained steels are less hardenable than coarse-grained steels with the same composition, and this generalization is true for case structure also. The slower the cooling of the steel during steelmaking, the larger the grain, since there are fewer nucleation sites formed. Steels

Fig. 54

Effect of decarburization on the residual stresses of carburized and hardened 3.5Ni-1.5Cr steel. The carbon content at 0.002 mm was approximately 1% for curve 1, 0.64% for curve 2, and 0.35% for curve 3.

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are typically deoxidized by silicon and manganese prior to aluminum addition. This will inhibit the undesirable formation of AlN or Al2O3 particles, which provide nucleation sites for coarse-grained carbides that lead to the

Fig. 55

Comparison of nominal ASTM No. 6 to 9 grain sizes. Etchant: nital. Original magnification: 100 ·. Source: Ref 96

formation of coarse ferrite/carbide (pearlitic) grain structures on cooling. Grain size is possible by control of the steel composition during the steelmaking process. Subsequent to this, control is by proper heat treatment. Heating the steel to the upper critical temperature, Ac1, will typically produce an average minimum grain size. Heating to higher temperatures will increase the grain size. Also, quenching from the Ac1 temperature will produce fine grain size, and quenching from a higher temperature would yield a coarser grain size. The austenite grain size at the onset of transformation during the quenching process will influence the martensite platelet size and thus will affect microcracking potential, the amount of retained austenite formed, and the frequency and depth of internal oxidation (Ref 95). Coarse grain structure is observed on etched microsections, most often in the form of coarse grain structure of martensite relative to finer retained austenite structure. Of the various factors affecting grain size, the primary factor is furnace treatment. Although coarse-grained steels exhibit better machinability, they generally possess lower toughness and ductility and exhibit a greater tendency for distortion and cracking than fine-grained steels. Coarsegrained steel also exhibits a more limited range of thermal treatment temperatures, and they possess better hardenability with higher asquenched hardness. Additionally, coarsegrained steels typically possess lower impact resistance and a lower yield point. Carburized steel case structures with a coarse-grained structure, along with significant amounts of overcarburizing, are characterized by decreased mechanical properties. In this case, the grain-boundary structure contains a continuous network carbide structure that is difficult to remove, and within the grain structure there are typically acicular carbides. Such microstructures are fatigue sensitive, with cracking throughout the brittle network carbide structure. The presence of these defective microstructures can be prevented by controlling the carburizing temperature and carbon potential or by using a steel with characteristically fine-grained structure. An example of grain size formation during austenitization of steel 40 is shown in Fig. 56. The irregular austenitic grain boundaries result from the short heating time and also are related to a prior normalizing and annealing process.

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This thermal history produces the variable grain size shown. Isogai et al. reported that grain sizes of approximately 8 were required for carburized transmission gear steel to achieve the required high fatigue strength, pitting strength, and impact strength in severe-use environments (Ref 97). To assure a fine grain size, after the carburizing process, the chromium-containing steel, such as SCr420, was reheated to 820 to 870 C for 20 to 60 min in an atmosphere containing the carburizing gas (carbon potential of 0.75%), quenched in oil until the steel was 120 C, and then tempered at 120 to 200 C. Treatment of the carburized steel in this way permits substantial reductions in grain size and corresponding improvements in fatigue strength and impact strength. Similar improvements were achievable with carbonitrided steel (Ref 97).

temperature, as well as the steel alloy chemistry. The total depth of internal oxidation can be calculated from (Ref 100): Xi 2 =½(2D0 C0 )=(n CM ) t=kp t Xi =½kp t0:5

The depth of internal oxidation that originates the carburizing atmosphere may vary from 1 to 30 mm (Ref 93, 98, 99). Internal oxidation consists of a continuous oxide layer on the surface, on the order of 0.01 mm, due to oxygen reaction with the carburized steel surface. In this region, the oxygen content can be 10 to 20 times that of the core. During the gas carburizing process, oxides will not only form on the surface but also penetrate into the steel surface. Since internal oxide formation is a diffusion process, the depth and extent of oxide penetration is dependent on the square root of the total carburizing time and

where Xi is the depth of oxygen penetration, D0 is the diffusion coefficient of oxygen in the alloy, C0 is the oxygen concentration at the steel alloy surface, CM is the concentration of the base metal in the alloy (e.g., silicon, chromium, manganese, titanium, vanadium), and n is the stoichiometric factor. Internal oxidation appears on the polished metallographic specimen in the form of very small inclusions concentrated in an austenite grain or within the grain boundary. The probability of oxidation within the grain increases as the grain size decreases (Ref 101). Figure 57 shows grain-boundary oxidation of carburized 20MnCr5 steel. Figure 58 also shows grainboundary oxidation, but it is accompanied by nonmartensitic transformation products. Internal oxidation occurs in two zones: an inner zone and an outer zone. Oxides of chromium-manganese are typically formed in the outer zone, both within the grain and in the grain boundaries. In the inner zone, silicon-rich oxides are typically formed exclusively within the grain boundaries (Ref 102). Figure 59 also illustrates intergranular oxidation of a gas-carburized steel (Ref 79). Typically, the greater the case depth, the greater the degree of oxide formation at the

Fig. 56

Fig. 57

Internal Oxidation (Ref 77)

Prior-austenite grains formed in hardened steel 40, which were due to abnormal growth during the austenitizing process, Etched: S. Bechet and L. Beaujurda. Original magnification: 500·

Illustration of grain-boundary oxidation of carburized 20MnCr5 to a depth of 30 mm. Unetched. Original magnification: 200·. Courtesy of Fluidtherm Technology P. Ltd., Ambattur, India

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surface and the greater the depth of internal oxidation. The thickness of the internal oxidation zone typically is approximately 5% of the carburized layer thickness and, on occasion, may be as high as 10%. The critical cooling rate is greater in the internal oxidation zone. As a result, an otherwise normal hardening produces a greater amount of bainite, which will lead to

Fig. 58

Illustration of grain-boundary oxidation with nonmartensitic transformation products to a depth of approximately 30 mm. Etchant: nital. Original magnification: 200·. Courtesy of Fluidtherm Technology P. Ltd., Ambattur, India

Fig. 59

lower surface hardness and poorer abrasive wear resistance. Typically, the oxides in the outer surface region are globular in form, while intergranular oxides were formed further from the surface. In one analyis of a carburized chromium-manganese steel, larger globular oxides were formed in the region closer to the surface (1.9 mm depth) and intergranular oxides in the region farther from the surface (2.49 mm) when the steel was heated for 16.6 h (diffusion at 2 h at 800 C, followed by 3 h at 930 C and a boost cycle at 930 C for additional heating times, in this case 11.6 h) (Ref 103). For this work, glow discharge optical emission spectroscopy, in which a sputter erosion process using ionized argon gas with a voltage of 600 V and 25 mA was used to quantify the degree of oxidation and elemental distribution, was performed using energy-dispersive x-ray analysis. When the carburized samples were subjected to shorter heating times, intergranular oxides formed relatively farther from the surface. Transmission electron microscopy was used by An et al. (Ref 103) to identify oxide type and morphology. Chromium and manganese globular oxides formed nearer the surface after a total heat carburizing cycle of 5.8 h (5+0.8 h). An agglomerated internal oxide of chromium

Intergranular oxidation of the surface along prior grain boundaries in a carburized steel. Original magnification: 1000 ·. Source: Ref 78

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and manganese oxides in the core surrounded by silicon oxides formed after heating for 16.6 h (5+11.6 h) (Ref 103). Internal oxide formation within grain boundaries provides sites for crack initiation (Ref 10). This was shown in a study by Laue et al. (Ref 104), who evaluated the fatigue behavior of case-hardened SAE 5115 steel with internal oxidation of the case structure. Fatigue studies were conducted on this steel, and it was shown that fatigue crack initiation occurred along the oxidized grain boundaries of the steel (Ref 104). Internal oxides form as a result of oxygen diffusion into the surface, with subsequent formation of metal oxides at carburizing temperatures. The formation of these oxides is enhanced by the presence of metals, chromium and manganese and Cr-Mn-Ti, which possess a greater affinity for oxygen than iron. The susceptibility for internal oxidation increases with increasing concentration of these oxide-forming elements. Lohrmann et al., referring to earlier work by Kozlovskii and co-workers, reported that the depth of internal oxidation was dependent on the total oxidation potential (TOP) of the steel alloy (Ref 105, 106): TOP=4:87 Si+3:7 Mn+1:47 Cr 3:24 Ni 1:82 Mo

Fig. 60

where the elemental compositions are given in weight percent. Figure 60 provides a correlation of the TOP and depth of internal oxidation (Ref 106). In addition, Kozlovskii et al. reported that most steels, when subjected to the gas carburizing process, will undergo internal oxidation with a corresponding surface formation of troostite to a depth of 0.01 to 0.03 mm. If troostite is formed at a depth greater than 0.014 mm, there is a substantial decrease in fatigue strength. However, the potential for internal oxide formation can be reduced by the addition of 5 to 10% of ammonia to the furnace for 10 min before the carburizing process is completed or by using steels containing 0.5% Mo and not more than 0.5% Cr (Ref 106). Lohrmann showed that the form and type of internal oxide obtained was dependent on the alloy composition of the steel. For example, depending on the steel alloy, spot-, liner-, or lattice-type internal oxides could be formed (Ref 105). Kehr and Seese examined the effect of internal oxidation during carburizing of investment-cast ingot-iron test specimen steels containing various amounts of chromium (0.20, 0.45, 0.91, 1.85, and 4.74% Cr) (Ref 107). It was shown that steels containing approximately 0.50% Cr, for

Correlation of calculated total oxidation potential (TOP) and average depth of internal oxidation

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example, SAE 8620, were more susceptible to internal oxidation than either plain carbon steel or steels with greater amounts of chromium, such as AISI 5120 (0.70 to 0.90% Cr). Figure 61 shows that the depth of oxide formation increases as the chromium content increases up to 0.45% (Ref 107). The presence of internal oxidation on carburized steels greatly reduces bending and contact fatigue strength and wear resistance (Ref 93). However, surface oxides that are formed may be removed by grinding or shot peening. The potential for internal oxidation can be reduced by heating steel to the carburizing temperature under a nitrogen/hydrogen gas mixture. During carburizing, these gases are replaced by a carburizing atmosphere whose oxygen activity is less than that required for the formation of manganese II oxide or chromium III oxide, and, in some cases, in the presence of ammonia (Ref 100). Finally, internal oxidation (and decarburization) is known to lead to variations in the surface compressive stresses in a carburized component (Ref 108). There is one report of the presence of internal oxidation leading to undesirable surface tensile stresses that then led to subsequent cracking of a carburized idler gear when used in a diesel engine gearbox.

Carbides and Carbide Structure Carbides formed during carburization are treated as undesirable products to be avoided.

Fig. 61

Effect of chromium content of steel on the depth of oxidation

There are three types of carbides to be discussed here: globular (or massive) carbides, network carbides, and surface-film or flake carbides. Carbides in steel are hard and brittle ceramiclike interstitials with a high compressive strength but low tensile strength (approximately 35 MPa, or 5000 psi). Carbides in steel basically form when carbon levels exceed the solubility limits of carbon in the iron crystal structure. The allotropic nature of iron also has different phase structures (i.e., crystal) with different solubility limits for carbon. For example, the maximum solubility of carbon in the body-centered cubic (bcc) structure of ferrite is approximately 0.025 wt% at 723 C on the iron-carbon phase diagram (see Appendix 9). For the face-centered cubic phase of austenite (c), the maximum solubility limit of carbon in c is approximately 2.06 wt% C at 1147 C. At still higher temperatures, another type of bcc solid phase is d-ferrite. The maximum solubility of carbon in d-ferrite is 0.09 wt% C at 1493 C. A peritectic also occurs at 0.16% C at 1493 C. The iron-carbon system has eutectic transformation at 1147 C during soldification, with steel carbon levels of 2.06 to 6.67 wt% C. The eutectic carbon concentration is 4.3%. Solidstate transformations in steel include the wellknown eutectoid transformation at 733 C, with a carbon concentration of 0.83%. At 733 C, austenite transforms to pearlite. Pearlite is a eutectoid mixture containing 0.83% C and is characterized by a fine ferrite-cementite structure that forms upon austenite decomposition during slow cooling. The upper critical temperature (A3) is the temperature below which ferrite starts to form due to ejection from austenite in the hypoeutectoid alloys. The lower critical temperature (A1) is the temperature where the austenite-topearlite eutectoid transformation occurs. Austenite does not exist below this temperature. A2 is the temperature below which ferrite is ferromagnetic. Cementite (Fe3C) is the principal carbide of iron and carbon, with an orthrhombic crystal structure. Cementite is harder and more brittle than ferrite. Steel hardness increases with increasing cementite content. Increasing carbon content increases the amount of cementite but does not affect the amount of ferrite present, since ferrite is saturated with 0.22% C. If there is less than 0.83% C, iron and carbon will combine to form Fe3C until no carbon remains. The

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cementite formed in this manner will combine with the required amount of ferrite to form pearlite, and any remaining ferrite will be in the structure as free ferrite (proeutectoid ferrite). Pearlite will form if the carbon content in the austenite is 40.83%, and excess carbon will form cementite. Excess cementite (proeutectoid cementite) will deposit in the grain boundaries. Because cementite (Fe3C) contains a specific amount of carbon and iron, pearlite also contains a specific amount of cementite and ferrite. As a phase in steel, the chemical composition of cementite will contain carbides of other carbide-forming elements, such as chromium and manganese Alloying elements of chromium, manganese, nickel, and other elements are, of course, commonly used in alloy steels for property improvement. They also impact the properties of ferrite and cementite, because they partition differently in the phases. For example, chromium and manganese partition in cementite instead of ferrite. However, nickel and silicon tend to favor partitioning in ferrite. Chromium, manganese, molybdenum, and titanium are thus cementite stabilizers in steels, while nickel and silicon are ferrite stabilizers. Interestingly, while chromium, manganese, molybdenum, and vanadium show no negative effect on cementite formation, titanium, nickel, and silicon exhibit a negative effect on cementite formation (Ref 109). Cementite develops different morphologies and distributions depending on the process of cementite formation. Figure 62 illustrates three microstructural forms of cementite: lamellar, mixed, and granular (Ref 110). Cementite may also be classified as reticular, acicular, or granular (Ref 110). Reticular cementite, also known as shell-type cementite, possesses a cracksensitive network or platelet structure. Acicular cementite, or needlelike structure, refers to a

Fig. 62

lamellar structure of cementite in ferrite, shown in Fig. 62(b). Finally, cementite exhibits a granular or grainy appearance, as shown in Fig. 62(c). Cementite forms during soldification from a liquid or during solid-state transformations. When cementite originates by crystallization from a liquid melt, it is referred to as primary cementite (Fe3CI). Secondary cementite (Fe3CII) is formed from austenite by hypereutectoid alloys (carbon40.8%). Tertiary cementite (Fe3CIII) is formed at temperatures below 723 C by precipitation in the grain boundaries (which become richer in carbon with the decreasing carbon content in a-iron). Globular Carbides. Slowly heating a steel to the carburizing temperature in the presence of a carburizing atmosphere through the ferrite-toaustenite temperature transformation region will lead to unconnected globular carbide formation either within the ferrite grains or at the former ferrite grain boundaries, as shown in Fig. 63. This process is favored by high carbon potentials and also by reduction of normal heating rates typically involved during carburizing, by excessive furnace loading, or by a furnace malfunction. Globular carbide formation may also be enhanced by austenitic nuclei or by localized concentrations of carbide-forming elements. The problem of globular carbides may also coexist with other problems, such as retaining austenite or quench cracking. When carburizing steels (0.15 to 0.25% C), which are commonly ferritic with localized areas of spheroidal carbides due to prior normalizing and subcritical annealing, are heated through the Ac1 temperature, the high-carbon regions begin to transform to austenite, resulting in the formation of localized regions of carbon and carbide-forming elements in addition to undissolved carbides in the presence of the

Cementite structures of CT60 steel with (a) lamellar, (b) mixed, and (c) granular cementite. Original magnification: 500 ·

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carbon-rich carburizing atmosphere. As the temperature increases further to Ac3, additional ferrite will transform to austenite by combining with carbon from the furnace atmosphere, preferentially over carbon transfer from the spheroidal carbides already present until the austenitic transformation is complete, at which point there will be an equilibrium between the austenite and carbides (Ref 112). Those steels containing alloying elements such as chromium and manganese that will reduce the eutectoid carbon content and Ac1 temperature are more likely to develop globular carbides. These elements will increase the solubility limit of carbon (Acm), thus shifting the equilibrium diagram to the left, as shown in Fig. 64 (Ref 112). It should be noted that although the terms spheroidal carbides and globular carbides are often used interchangeably, these carbides may, in fact, possess a round, angular, or even needlelike appearance. Also, although the structures are typically designated as M3C, their actual ratio of the element/carbon composing the carbide being observed is not only dependent on the time and temperature of heat treatment but primarily on the elemental availability in the steel during formation. There is a three-step process of high-density carburizing that is conducted to increase case hardness by aggressive precipitation of cementite (Fe3C) to improve surface fatigue strength (Ref 113). Generally, grain-boundary carbide precipitation will form network carbide structures that are susceptible to quench cracking and

Fig. 63

Bad globular carbide formation in the case of a carburized 9310 steel. Etchant: boil in alkaline sodium picrate solution (45 s). Scale = 10 mm. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

result in reduced fatigue strength. Interestingly, it is possible to obtain desirable surface fatigue strength by vacuum or plasma carburizing a chromium-containing steel such as SCr420H or a chromium-molybdenum steel such as SCM420H (Ref 114). This process involves the formation of a M23C6-type globular carbide that is approximately 1 mm in diameter. These microstructures are reported to exhibit excellent surface fatigue strength and rolling fatigue strength under high bearing loads (43 GPa) at relatively high temperature (100 to 300 C), which are unachievable in the presence of larger (i3 mm) M3C carbides. Reheating a carburized steel to a temperature below Acm causes spheroidized carbide particles to form. (Both austenite and cementite are stable at this temperature.) Since these carbides bind some of the carbon in the case, there will be an increase in the Ms temperature. In addition, grain-boundary migration is reduced (Ref 115). Carbides are typically very hard. For example, the microhardness of globular carbides in a carburized case of a plain carbon steel has been measured to be 41000 HV, and carbides in a 2% Ni-Cr steel have been measured to be approximately 800 HV (Ref 112). Therefore, the presence of free globular carbides is often assumed to improve component wear, abrasion, and scuffing resistance. However, in a study reported by Parrish on the effect of globular carbides on contact fatigue with carburized 2% Cr-Mn steel, it was shown that the presence of massive

Fig. 64

The saturation surface carbon content (Acm) of various carburizing steels as a function of carburizing temperature as related to the iron-carbon equilibrium diagram

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carbides was detrimental, while a beneficial effect was obtained if the carbides were finer and better distributed, as shown in Fig. 65 (Ref 112). The presence of globular carbides may also lead to grinding cracks, which are also related to surface residual stresses. Network Carbides. Globular carbides are most typically formed when the carbon content is less than the Acm. However, under some conditions, it is possible for globular carbides to still be formed when the carbon content is greater than the solubility limit (in excess of the eutectoid composition) and austenite is supersaturated with respect to carbon. Most typically, these are conditions where carbon will precipitate in the grain boundaries as cementite (Fe3C) during slow cooling from the carburizing temperature, leading to the formation of network carbides. Figure 66 illustrates network carbide structure observed in a broken carburized AISI P5 steel tool. The equilibrium diagram shown in Fig. 67 illustrates the conditions for the formation of network carbides when excess carbon is precipitated from austenite as Fe3C (Ref 112). When the carbon steel supersaturated with carbon content of C1 is cooled from t0 to t1, Fe3C will begin to precipitate. As cooling continues to t2, additional carbon will have precipitated as Fe3C until the carbon content is C2. At the

Fig. 65

eutectoid temperature t3, carbon precipitation as Fe3C will have stopped, and the austenite carbon content C3 will transform to a pearlitic eutectoid microstructure. The relative proportion of Fe3C to austenite can be determined from Fig. 67 using the lever rule (CxC1/C1B). Although austenite is supersaturated with respect to carbon during the carburizing process, and carbide precipitation at austenitic grain boundaries will occur during cooling, if the steel is quenched from the carburizing temperature, the excess carbon can be retained by the resulting as-quenched martensitic/retained-austenite microstructure. Typically, during carburizing, the load is cooled in the furnace from the carburizing temperature to the temperature from which the steel will be hardened. Since the Fe3C migrates to the grain boundaries during slow cooling, faster cooling by quenching will allow retention of the carbon in the martensitic/austenitic structure. Thus, cooling rate can be used to control the amount of network carbide formation. Alternatively, carburized steel can be cooled to ambient temperature and reheated to 820 to 860 C and quenched. However, it has been shown that traces of the network carbides remained in the microstructure even after heating to 900 C (Ref 112). Parrish also has reported that if steel containing excess carbon is subcritically annealed prior to quenching, a fine

Contact fatigue strength of carburized 25Kh2GHTA steel (tempered at 180 to 200 C)

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Fig. 66

Micrographs of a broken carburized P5 tool steel die. Note the layer of cementite along the surface and the heavy grain-boundary network in (a). The case was 59.5 HRC, and the core was 22 GRC. (a) The case is shown at an original magnification of 100· (100 mm bar). (b) The case at an original magnification of 1000 · (10 mm bar). (c) The core at an original magnification of 400· (25 mm bar) in Nomarski differential interference contrast (note the “old” and “new” ferrite, as in dualphase steel). The austenitizing temperature for the case is approximately 1475 F, which is in the two-phase field for the core. New ferrite formed due to limited hardenability in the quench. Courtesy of G. Vander Voort, Buehler Ltd., Lake Bluff, IL

dispersion of relatively unharmful spheroidized carbides will be obtained instead of network carbides (Ref 112). Karpov studied the effect of cooling from the carburizing temperature on the nature of network carbides formed when quenching gascarburized 11 by 11 by 56 mm 07Kh16N6 steel test specimens. The test specimens were carburized, cooled, then reheated to austenitize to 1020 C, and then cooled to the quenching temperature at 0.036 C/s, which is less than the critical quenching rate that leads to network carbide segregation. Quenching was conducted in water at room temperature. From Table 11, it is evident that network carbide formation begins at 900 to 850 C for this alloy and is completed at 600 C. From the phase diagram for this alloy, the Acm temperature is 860 C (Ref 116). During the course of this work, Karpov found that a nondestructive electromagnetic flaw detector could be used to rank the network carbide size (Ref 116). Network carbides have been reported to reduce surface fatigue (pitting) resistance of carburized steel used for bearing applications (Ref 114). Like globular carbides, complex carburized networks in an overcarburized case will reduce the potential for carbide redissolution during reheating, which will lead to increased brittleness and grinding cracks (Ref 1). The presence of network carbides in a Kaplan turbine blade constructed from improperly carburized 17CrNiMo6 steel was reported to be a major contributor to failure by an intergranular microcracking mechanism (Ref 117). The cracks seemed to follow the path of the network carbide structure. Parrish has summarized various studies and concluded that continuous network carbides do reduce fatigue properties, leading to premature cracking failures by a stiffening mechanism (Fig. 68) However, other studies with partial nertwork carbides showed no deleterious effects (Ref 112). The presence of network carbides also is not expected to produce an adverse effect with respect to wear under heavily loaded conditions or scuffing (Ref 112). The three most common failure modes of carburized steels are ductile fracture, cleavage, and intergranular fracture (Ref 23). Ductile fracture is caused by nucleation growth and coalescence of voids that are initiated at inclusion sites and second-phase particles. Cleavage fracture occurs by separation at crystallographic planes by a transgranular pathway. Intergranular fracture, such as that involving cementite

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Fig. 67

The use of the iron-carbon equilibrium diagram to illustrate network carbide formation

Table 11 Effect of initial cool-down period on network carbide formation of carburized 07Kh16N6 steel Cool-down temperature(a), °C

1020 (initial temperature) 950 900 850 800 700 600

Network carbide severity rating(b)

1 1 1–2 1–3 1–3 3–4 5–5

(a) The cooldown occurs from 1020 to the temperature shown at a rate of 0.036 C/s. (b) 1, very fine carbide network; 3–4, failure rating; 5, largest

deposition at grain boundaries, involves cracking on grain boundaries and is due to (Ref 23):

Precipitation of a brittle phase (such as network carbides) on the grain boundary Hydrogen embrittlement Environmental-assisted cracking Intergranular corrosion

Grain-boundary cavitation temperature cracking

and

high-

Film or Flake Carbides. Surface-film or flake carbides are composed of a continuous or discontinuous carbide film with typically little or no penetration into the case structure, which is caused by cooling of the carburized steel in the furnace with high carbon potential. Parrish summarized various previously published reports that stated such carbide films contain approximately 19% Fe3C, 16% austenite, with the balance being martensite. These films cover a nonmartensitic layer of approximately 30 mm thickness, with carbides penetrating into the grain boundaries (Ref 112). Koistinen showed that such films have high tensile surface (0.025 mm depth) residual stress, as shown in Fig. 69 (Ref 44). Until now, no further information relating to the presence of these carbide types on carburized steel properties has been reported.


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Noncarbide Inclusions Two of the primary causes of fatigue failure are inclusions and surface defects. Inclusions may be metallic impurities or metallic oxides

(Ref 118). Metallic element inclusions (impurities), although typically in trace quantities, may be traced to the scrap used in the steelmaking process. These elements cause intergranular segregation, which may lead to crack

Fig. 68

Comparison of bending fatigue of carburized 12Khn3 gears showing adverse effect of network carbides

Fig. 69

Residual-stress distribution of carburized SAE 1018 steel with a film-carbide layer formed due to a high carburizing potential. The surface layer consisted of 16% Fe3C, 16% retained austenite, and the balance was as-quenched martensite.

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formation, detrimental precipitate formation, and are often observed as “slivers” in the final product. Metallic oxide inclusions vary in morphology and composition. Some sources of these oxides include:

Deoxidation products, such as alumina inclusions, that are formed by the reaction of dissolved oxygen and the added deoxidant, such as aluminum. Alumina inclusions are dendritic when formed in the presence of high oxygen concentration. Reoxidation products are generated when the aluminum remaining in the liquid steel is oxidized by FeO, MnO, or SiO2 and other oxides in the slag or refractory materials or by exposure to the atmosphere. Slag entrapment occurs in metallurgical fluxes entrained in the steel. This occurs during transfer between steelmaking vessels. These inclusions are typically spherical. Exogenous inclusions from other sources include dirt, broken refractory brickwork, and ceramic lining materials. They act as sites for heterogeneous nucleation of alumina and may include a central particle. Chemical reactions may produce oxides from inclusion modification when calcium treatment is imperfectly performed. Inclusions containing CaO may also originate from entrained slag.

All steels contain various noncarbide inclusions. It is well known that inclusions, in addition to surface defects and inhomogenities such as retained austenite, nonmetallic inclusions, and inhomogeneities, can reduce fatigue

Fig. 70

strength and negatively influence ductility and toughness (Ref 119). Although harder inclusions such as metallic inclusions are more harmful than softer inclusions, nonmetallic inclusions are still generally harmful. In addition, the deleterious effects of inclusions increase with size. In one study conducted by Bomas and Schleicher on the effect of inclusions on the fatigue strength of carburized 16MnCrS5 (SAE 5115) steel, it was found that subsurface fatigue crack initiation was initiated by nonmetallic inclusions up to depths of 1.4 mm (Ref 120). Similar results were obtained for a study of MnS-induced bending fatigue failure of carburized EN39B steel (Ref 121). Even in clean steels, oxide and sulfide inclusions exist. From a study of the effect of defects such as inclusions on the fatigue bending strength of carburized SCM20 steel, a model of fatigue crack initiation during fatigue bending, such as would occur in a gear tooth, was developed (Ref 122). This model is illustrated schematically in Fig. 70. Figure 70(a) illustrates the chromium and manganese oxides and grainboundary segregation of silicon oxides present in a decarburized surface layer. When the surface is loaded, the grain boundary or oxides act as a stress raiser for microcrack formation, as shown in Fig. 70(b). Although most of the cracks remain in the decarburized layer, the most critical crack penetrates deeper into the steel, as shown in Fig. 70(c). The crack that initiates fatigue failure was thought to be due to a combination of the microcracks shown in Fig. 70(b). In this case, the decarburized layer in combination with the presence of inclusions was projected to be the source of fatigue failure (Ref 122).

Model of fatigue crack initiation due to the presence of inclusions in a nonmartensitic (decarburized) steel layer. Source: Ref 122

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Micropitting Micropitting fatigue usually occurs on heavily-loaded surface-hardened components and is characterized by a frosted or gray-stained appearance under thin-film lubrication conditions (Ref 87, 123). Numerous small cracks in the surface may exceed the depth of the micropits (Ref 124). Under magnification, small pits approximately 10 mm deep will be observed. The surface will appear etched with a pattern that sometimes follows the slightly higher ridges left by cutter marks or other surface irregularities on the finished component. Micropitting is influenced by high surface loads, frictional heat generation due to poor lubrication caused by insufficient film thicknesses in the partial elastohydrodynamic lubrication regime, excessive retained austenite, tangential speed, and lubricant additives and oxidation (Ref 123). Micropitting is strongly influenced by the relative lubricating oil film thickness and can be quantitatively related to both the surface condition (roughness) and the thickness of the lubricating film by the lambda (L) value (Ref 124, 125): h L=h=s= qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 (s1 +s22 )

where s1 and s2 are the root mean square surface roughness of the two opposing wear surfaces, and h is the lubricating oil film thickness. When Li3, there is full film lubrication with no asperity contact. When 0.85L53, there is partial elastohydrodynamic lubrication. When L50.8, there is a boundary lubrication condition. When L51, micropitting will occur, and once micropitting occurs, pitting fatigue (macropitting) will be accelerated (Ref 124). However, if macropitting does occur, it is often characterized by an arrowhead or fan shape (Ref 123). There have been reports of substantial improvements in fatigue lives, such as with carburized 9310 steel, with corresponding reductions in surface roughness (Ref 125). For gears, there is a critical temperature where pitting fatigue and scuffing are likely to occur. This is called the critical scuffing temperature (Tc), which is calculated from: Tc =Tb +Tf

where Tb is the equilibrium temperature of the gears before meshing, and Tf is the flash

temperature, which is the instantaneous temperature rise due to localized friction heat at the point where the gear teeth mesh. The value of Tb is controlled by gear geometry design, and the value of Tf is controlled by the lubricant viscosity and surface roughness. To minimize micropitting:

Use higher operational speeds and smooth material surfaces. Use the recommended amount of clean, dry lubricant with the highest viscosity permissible. Reduce the lubricating oil temperature and surface loading. Use the optimal amount of case carbon content in carburized gear materials.

Contact Fatigue Pitting (Macropitting) Pitting failures occur when fatigue cracks are initiated on the tooth surface or just below the surface. Usually, fatigue pits are the result of surface cracks caused by metal-to-metal contact of asperities or defects due to insufficient lubricant film thickness. They are dependent on the Hertzian contact surface stress and the number of stress cycles (Ref 124). Surface asperities of the harder material of a wear contact will lead to damage of the softer surface, sometimes by a work-hardening mechanism, leading to the creation of microcracks that then become fatigue pits as the wear process continues (Ref 124). Pitting damage is commonly encountered with rolling element bearings, gears, and machine components subject to cyclic rolling-sliding motion under a load. Initially, fatigue pits may occur in localized areas and may range in size from 0.38 to 0.76 mm (0.015 to 0.030 in.) in diameter (Ref 85). Vinokur et al. examined the effect of case carbon content of carburized 18KhGNMFL steel (1.3% Mn, 1.3% Cr, 0.8% Ni, 0.25% Mo, 0.1% V) on the contact endurance of fatigue tests conducted with a wear test with an applied stress of 3500 MPa using a lubricating oil. Contact fatigue was the average of 10 tests to determine the number of cycles until pitting was observed. The case depth was approximately 1.8 mm. The carburized steel was hardened from the intercritical range and tempered at 170 C. The results of this study are summarized in Fig. 71 (Ref 126). The contact endurance increases with carbon content up to approximately 0.9% C and

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then decreases. The optimal hardening temperature was at approximately 850 C, which is in the intercritical temperature range, just below the upper critical temperature for the steel. Although fatigue pits are usually initiated at the surface, subsurface initiation is relatively common in case-hardened rolling element bearings with serious inclusion problems. In these situations, failures usually do not follow the case-core interface (Ref 125, 127). The potential for micropitting, pitting, and spalling phenomena may be assessed from the lambda (L) value (see the section “Micropitting” in this chapter); however, even relatively smooth surfaces and lubricating film thickness, such as those used for high-speed gears, may exhibit pitting failures due to the presence of subsurface cracks. The subsurface cracks may be caused by the presence of inclusions that act as stress concentrators, causing the crack to propagate parallel to the surface and subsequently break through the surface. Spalling failure occurs when the wear process causes several pits to join together. These larger craters are usually caused by more severe overloading conditions. As the number of stress cycles increases, the pitting process will continue in an effort to relieve stresses. The rule is that spalling cracks initiate where the ratio of shear stress to Vicker’s hardness is maximum (Ref 124). However, this relationship is not correct when there are excessive amounts of retained austenite. Generally, it is assumed that spalling will occur when L41.

Fig. 71

Effect of case carbon content and hardening temperature on the contact endurance limit of carburized 18KhGNMFL steel (1.3% Mn, 1.3% Cr, 0.8% Ni, 0.25% Mo, 0.1% V). Source: Ref 126

To prevent pitting fatigue, either the surface loading must be decreased to a level below the endurance limit of the material or the hardness must be increased to increase the endurance limit (Ref 28). Pitting may also be reduced by instituting a break-in period at reduced loads and speeds to improve gear tooth contact (Ref 85, 124]. Other potential causes of fatigue pitting include hydrogen embrittlement due to water contamination of the lubricant and particle contamination of the lubricant, which act as surface stress-concentration points that lead to pitting failure.

Case Crushing If the case depth is too deep, case-core separation may occur due to the tips of the gear teeth becoming too brittle and possibly breaking, if the case depth is too thin, the strength of the gear teeth will be reduced, causing premature pitting, or it may lead to a condition called case crushing (Ref 131). Case crushing occurs in heavily loaded case-hardened components such as gears. Case crushing occurs by a subsurface fatigue process where the high-cycle contact stress exceeds the endurance limit. This will occur when subsurface stresses exceed the strength of the core. Case crushing failures may have a similar appearance to pitting, although it often occurs as longitudinal cracks on the surface of only one or two gear teeth, where sections of the tooth surface may subsequently break away. However, the case material may appear to have chipped away from the core in large flakes (Ref 85). The observed cracks will move toward the case-to-core boundary and then to the gear surface (Ref 129). Adequate case support is provided by proper core structure to not only prevent case crushing but also to transmit torque, support bending loads, and provide adequate toughness to prevent brittle fracture. The presence of any ferrite will contribute to reducing the toughness of the core. Case crushing may be prevented by increasing the case depth and possibly the core hardness. For general applications where core hardnesses of 30 to 45 HRC are specified, the required case depth can be estimated from (Ref 128): Case depth to 50 HRC=½1:2 · 107 (W)=F

where W is the force in pounds pressing the surfaces together, and F is the length of the line

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contact in inches. The strength of the core can be determined from hardness. Since case crushing is promoted by shear, the shear strength of the core must be determined and can be estimated from Fig. 72 (Ref 128). Generally, the subsurface stress/strength ratio should not be greater than 0.55. Eryu et al. have studied case-crushing fracture mechanisms of carburized 20CrMnTi by scanning electron microscopy (Ref 130). They showed that the surface features of the primary cracks exhibited scaly features and that there were two features of the fracture surfaces of the branching cracks. The fatigue steps and dimples were analogous to materials of higher strength, and the fragmentation pattern was analogous to brittle material. In addition, spherical particles were observed that were composed of a-iron, which were proposed to be caused by the movement of the faces of the primary crack due to shearing and compressive stresses.

Pitting Corrosion Pitting corrosion is a localized penetrating corrosion attack of typically corrosion-resistant steel resulting in a mass loss of the steel (Ref 131). Pitting corrosion is related to localized discontinuities of a passive layer caused by mechanical imperfections, inclusions, surfacelocalized chemical attack of the passive layer by salts such as chlorides, or by overaggressive lubricant additives. After the corrosion pit is

Fig. 72

Shear strength of carburizing steels as a function of hardness

created, the localized chemical surrounding is much more aggressive than the surrounding area of the uncorroded material. Initiation of the pitting process is dependent on temperature and on the steel surface, including the presence of sediment. In some applications where pitting corrosion is more prevalent, such as steel in concrete structures, the pitting corrosion process is characterized by the temperature or narrow range of temperatures above which pitting will nucleate. The creation of corrosion sediments will lead to a temperature decrease. Pitting corrosion will only occur above this critical temperature. Therefore, to increase the lifetime of steel used in reinforced concrete by reducing the rate of pitting corrosion, frequent sediment removal (cleaning) is recommended. Most often, pitting corrosion is initiated by the presence of chloride salts, and the rate of corrosive attack is steel alloy dependent. The critical concentration of chlorides for different steel alloys cannot be defined, because corrosivity is dependent on other chemicals that may be present, which will affect the rate of corrosion attack. However, since pitting corrosion is typically relatively fast, it should be prevented. Resistance of steel to pitting corrosion is dependent on the alloy composition (chromium, molybdenum, tungsten, nitrogen). Relative corrosion resistance of steel alloys may be empirically quantified by: Relative corrosion resistance=%Cr +3:3 ½%Mo+0:5 (%W)+16 (%N)

Chromium and molybdenum are also useful alloy additions to minimize the potential for stress-corrosion cracking. Corrosion pitting may also be caused by chemical attack of the steel surface by lubricant additives such as extreme-pressure additives, particularly in the presence of acid, water, or contaminants. Also, during use, the oil itself will oxidatively degrade, producing acidic by-products that may lead to corrosion pitting. In addition to pitting, corrosive attack may occur at the grain boundaries of the carburized case. It is particularly important that components exposed to saltwater, liquid chemicals, or other foreign materials during use should be sealed from their operating environment (Ref 85). Pitting corrosion may also occur during heat treatment, particularly those processes involving salt baths (Ref 132). Heating of steel with scale on a surface not only accelerates decarburizing,

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pg 233


but pitting corrosion is also accelerated. The corrosive attack is increased with temperature and holding time. If the scale is not uniform, then pitting corrosion is localized to those areas where scale is present. If residual salt from the bath crystallizes on the surface of the steel, violent boiling may occur during subsequent quenching in oil, which may result in blister formation on the surface. After cleaning, localized pitting corrosion will then occur. To avoid pitting corrosion during furnace heating, parts should be thoroughly cleaned. In some cases, stainless steel machine parts are carburized to reduce wear. However, carburizing a stainless steel (1Kh16N2AM) reduces corrosion resistance where machine parts are used in humid environments (Ref 133). The potential for pitting corrosion increases as the amount of d-ferrite increases. Corrosion resistance is reduced as the quenching temperature prior to carburizing is increased. Although increasing the quenching temperature after carburizing to 950 to 1100 C does not affect corrosion resistance, decreasing the quenching temperature to 800 to 850 C reduces corrosion resistance. Increasing the tempering temperature decreases corrosion resistance, as shown in Fig. 73 (Ref 133). The corrosion resistance of carburized steel is greatest after stress relieving at 250 C.

Partial Melting Partial melting occurs when there is nonuniform heating of the surface of the steel, such that some areas are heated to the liquation temperature (the partial melting temperature of an alloy) (Ref 77, 131). Corners and edges are particularly susceptible to partial melting. Microscopically, the presence of partial melting is typically observed as black spots containing retained austenite in a large cluster of carbides. Macroscopically, partial melting is accompanied by the formation of tiny surface cracks. Partial melting occurs when the carburized steel is heated to an excessively high temperature, resulting in incomplete or selective carburizing of the surface. For example, partial melting may occur during stray current flow into the load from electrodes used to heat salt pot furnaces or if a load is placed too close to the furnace hearth, so that some areas of the load are heated to an abnormally high temperature. To avoid this defect, heating in salt baths with appropriate composition and at appropriate austenitizing temperature should be conducted by keeping the load at a recommended distance from the heating electrodes. Similarly, when heating in a conventional furnace, the load should be properly placed to facilitate uniform heating. It is also important to be aware of the liquation temperature (beginning of melting) of the alloy being heated. Some typical examples of the approximate liquation temperatures for different steels are provided in Table 12, where the soaking temperature is 2 to 3 min. Precise definition of the partial melting temperature range is typically a difficult task, because of the relatively large data scatter due to potential compositional variation within the alloy, variation of carbon content, and, in some cases, relatively large carbide segregation. These structural variations favor the potential for partial melting to Table 12 Approximate liquation temperature for various steel alloys

Fig. 73

Effect of tempering temperature on corrosion resistance of carburized stainless steel 1Kh16N2AM. The corrosion test was conducted in a humidity cabinet. Source: Ref 133

Steel alloy

Approximate liquation temperature, °C

SW14 SW18 SW7Mo SK5 SKC SK5V SW12C SK10V

1320 1330 1280 1350 1280 1270 1260 1250


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occur or the creation of ledeburite networks. Improvement in temperature control will reduce the potential for partial melting defects.

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44. D.P. Koistinen, The Distribution of Residual Stresses in Carburized Cases and Their Origin, Trans. ASM, Vol 50, 1958, p 227–241 45. J.L. Pacheco and G. Krauss, Microstructure and High Bending Fatigue Strength of Carburized Steel, J. Heat Treat., Vol 7 (No. 2), 1989, p 77–86 46. O. Asi, An Investigation of Retained Austenite Contents in Carburized SAE 8620 Steel, G. U. J. Sci., Vol 17 (No. 4), 2004, p 103–113 47. V.F. da Silva, L.F. Canale, D. Spinelli, W.W. Bose-Filho, and O.R. Crnkovic, Influence of Retained Austenite on Short Fatigue Crack Growth and Wear Resistance of Case Carburized Steel, J. Mater. Eng. Perform., Vol 8, 1999, p 543–548 48. D. Jeddi, H. Sidhom, D. Ghiglione, and H.-P. Lierade, Role of Cyclic Stability of Retained Austenite in Fatigue Performance of Carburized 14NiCr11 Steel, J. Mater. Eng. Perform., Vol 14, 2005, p 37–49 49. H.-J. Kim and Y.-G. Kweon, The Effects of Retained Austenite on Dry Sliding Wear of Carburized Steels, Wear, Vol 193, 1996, p 8–15 50. I.S. Kozlovskii, S.E. Manevskii, and I.I. Sokolov, Effect of Retained Austenite on the Resistance to Scoring of Steel 20Kh2N4A, Met. Sci. Heat Treat. (USSR), Vol 20 (No. 1), Jan 1978, p 78–80 (translated from Russian) 51. C. Kim, Fracture of Case-Hardened Steels and Residual Stress Effects, Case-Hardened Steels: Microstructural and Residual Stress Effects, D.E. Diesburg, Ed., The Metallurgical Society of AIME, 1984, p 59–87 52. J.A. Burnett, Prediction of Stresses Generated during the Heat Treating of Case Carburized Parts, Residual Stress for Designers and Metallurgists, L.J. Vande Walle, Ed., American Society for Metals, 1981, p 51–69 53. N. Kanetake, Chapter 17—Effect of Carburising on Mechanical Properties of Steels, Quenching and Carburizing, Proceedings of the Third Seminar of the International Federation for Heat Treatment and Surface Engineering, The Institute of Materials, London, U.K., 1993, p 281–291

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85. G. Parrish, Chapter 8—Postcarburizing Mechanical Treatments, Carburizing: Microstructures and Properties, ASM International, 1999, p 199–230 86. M. Lawerenz and I. Ekis, Creating an InHouse Shot Peening Specification for Gears—Part 2, Paper 1993118, Shot Peener, Vol 7 (No. 1), 1993, p 3–5 87. C. Bahnsen, B. Clausen, F. Hoffmann, P. Mayr, and H.-W. Zoch, Influence of Carburising on Distortion Behavior, Proceedings of First International Conference on Distortion Engineering, Sept 14–16, 2005 (Bremen, Germany), p 235– 242 88. M. Izciler and M. Tabur, Abrasive Wear Behavior of Different Case Depth Gas Carburized AISI 8620 Gear Steel, Wear, Vol 260 (No. 1–2), 2006, p 90–98 89. M. Boniardi, F. D’Errico, and C. Tagliabue, Influence of Carburizing and Nitriding on Failure of Gears—A Case Study, Eng. Fail. Anal., Vol 13, 2006, p 312–339 90. G. Parrish, Chapter 6—Core Properties and Case Depth, The Influence of Microstructure on the Properties of CaseHardened Components, American Society for Metals, 1980, p 131–165 91. L. Faqian, L. Shuanlu, l. Helin, Q. Changyi, Z. Jie, T. Jiping, and C. Jun, Analysis of Causes of Casing Elevator Fracture, Eng. Fail. Anal., Vol 14, 2007, p 606–613 92. Z. Yu, X. Xu, and H. Ding, Failure Analysis of a Diesel Engine Piston-Pin, Eng. Fail. Anal., Vol 14, 2007, p 110–117 93. Z.Ya. Nostislavskii, Yu.N. Griboedov, and S.I. Markov, Effect of the Atmosphere During Heating of Carburized Cr-Ni Steel to Quenching Temperature on Decomposition of Austenite, Met. Sci. Heat Treat. (USSR), Vol 18 (No. 5–6), May–June 1976, p 443–445 94. G. Parrish, Chapter 2—Decarburization, Carburizing: Microstructure and Properties, ASM International, 1999, p 37–49 95. G. Parrish, Chapter 5—Influential Microstructural Features, Carburizing: Microstructure and Properties, ASM International, 1999, p 99–133 96. Atlas of Microstructures of Industrial Alloys, Vol 7, Metals Handbook, 8th ed., American Society for Metals, 1972, p 4 97. T. Isogai, R. Uchino, M. Sato, and Y. Kono, Method for Heat Treatment of

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Steel, U.S. Patent 6,149,734, Nov 21, 2000 E. Gianotti, I. Calliari, M. Zanesco, and E. Ramous, Grain Boundary Oxidation in Endothermic Gas Carburizing Process, Proceedings of the Seventh International Tooling Conference, Tooling Materials and Their Applications from Research to Market, Vol 2, May 2–5, 2006, Politecnico Di Torino, Torino, Italy, p 481–490 D.K. Matlock, K.A. Alogab, M.D. Richards, and J.G. Speer, Surface Processing to Improve the Fatigue Resistance of Advanced Bar Steels for Automotive Applications, Mater. Res., Vol 8 (No. 4), 2005, p 453–459 H.-P. Schmidt, Process for Avoiding Surface Oxidation in the Carburization of Steels, U.S. Patent 5,498,299, March 12, 1996 V.D. Kal’ner and S.A. Yurasov, Internal Oxidation During Carburizing, Met. Sci. Heat Treat. (USSR), Vol 12, (No. 6), June 1970, p 451–454 (translated from Russian) J.D. Verhoeven, N. Chen, and A.J. Bevolo, Scanning Auger Microscopy Characterization of the Internal Oxidation Produced on Carburizing, J. Heat Treat., Vol 9 (No. 2), 1992, p 73–80 X. An, J. Cawley, W.M. Rainforth, and L. Chen, A Study of Internal Oxidation in Carburized Steels by Glow Discharge Optical Emission Spectroscopy and Scanning Electron Microscopy, Spectrochim. Acta B, At. Spectrosc., Vol 58, 2003, p 689–698 S. Laue, H. Bomas, and F. Hoffman, Influence of Surface Condition on the Fatigue Behaviour of Specimens Made of SAE 5115 Case-Hardened Steel, Fatigue Fract. Eng. Mater. Struct., Vol 29, 2006, p 229–241 M. Lohrmann and W. Lerche, Investigation of Process and Steel Influences on the Extent of Internal Oxidation during Gas Carburizing, Heat Treat. Met., Vol 27 (No. 4), 2000, p 85–90 I.S. Kozlovskii, A.T. Kalinin, A. Va. Novikova, E.A. Lebedeva, and A.I. Feofanova, Internal Oxidation during Case-Hardening of Steels in Endothermic Atmospheres, Met. Sci. Heat Treat (USSR), Vol 9 (No. 3), March 1967, p 157–161 W.D. Kehr and H.E. Seese, Surface Attack in Chromium-Iron Alloys, IBM J., Sept 1966, p 420–421

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108. Z. Yu and X. Xu, Failure Analysis of an Idler Gear of Diesel Engine Gearbox, Eng. Fail. Anal., Vol 13, 2006, p 1092–1100 109. M. Umemoto, Z.G. Liu, K. Masuyama, and K. Tsuchiya, Influence of Alloy Additions on Production and Properties of Bulk Cementite, Scr. Mater., Vol 45, 2001, p 391–397 110. S.N. Saltykov and N.V. Tarasova, Anodic Dissolution of Ferrite Phases of IronCarbon Ferrite-Cementite Alloys with Different Forms of Cementite, Prot. Met. (Russia), Vol 42 (No. 5), 2006, p 501–506 111. V.Ya. Zubov, Patenting of Steel Wire, Met. Sci. Heat Treat., (USSR), Vol 14 (No. 9), Sept 1972, p 793–800 112. G. Parrish, Chapter 3—Carbides, The Influence of Microstructure on the Properties of Case-Hardened Components, American Society for Metals, 1980, p 39–60 113. T. Ichihara and H. Fujiwara, Pullu Unit Having One-Way Clutch, U.S. Patent Application US2003/0085091 A1, May 8, 2003 114. T. Nakamura, S. Nakamura, T. Kimura, T. Yamaguchi, K. Otani, N. Uchiyama, N. Kino, and T. Nanbu, High Bearing Pressure-Resistant Member, U.S. Patent 6,569,267 B1, May 27, 2003 115. G. Krauss, Heat Treatment, Microstructures, and Residual Stresses in Carburized Steels, Conf. Proceed. Quenching and Distortion Control, G.E. Totten, Ed., ASM International, 1992, p 181–191 116. L.P. Karpov, Effect of Quenching Conditions on the Carbide Network in Steel 07Kh16N6, Met. Sci. Heat Treat. (USSR), Vol 35 (No. 1), Jan 1993, p 20–23 (translated from Russian) 117. G. Diego, M. Serrano, and A.M. Lancha, Failure Analysis of a Multiplier from a Kaplan Turbine, Eng. Fail. Anal., Vol 7, 2000, p 27–34 118. L. Zhang and B.G. Thomas, State of the Art in Evaluation and Control of Steel Cleanliness, ISIJ Int., Vol 43 (No. 3), 2003, p 271–291 119. Y. Murakami and M. Endo, Effects of Defects, Inclusions, and Inhomogenities on Fatigue Strength, Fatigue, Vol 16, April 1994, p 163–182 120. H. Bomas and M. Schleicher, Application of the Weakest-Link Concept to the Endurance Limit of Notched and Multiaxially Loaded Specimens of Carburized

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(No. 5), May 1970, p 393–395 (translated from Russian) 133. V.I. Belyakova, M.F. Alekseenko, L.Ya. Gurvich and V.L. Erofeeva, The Corrosion Resistance of Carburized Stainless Steels, Met. Sci. Heat Treat. (USSR), Vol 15 (No. 2), Feb 1973, p 143–146 (tanslated from Russian)

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Fatigue Fracture of Nitrided Layers Aleksander Nakonieczny, Institute of Precision Mechanics

THE DURABILITY of products depends strongly on surface conditions on the order of 1 mm to several millimeters, depending on the type of technological process applied. The condition of the surface layer is critical to wear resistance during the process of friction and to corrosion resistance. In the case of mechanical loading (especially fatigue resistance) and corrosion, a critical role is played by the substrate, its condition and properties, as well as the atomic relationship between the surface layer and the substrate. For this very reason, one should take into account the substrate-surface layer system when considering the life expectancy of machine components and assemblies. Attributes of such a system are the thickness of the surface layer, the ratio of this thickness to the entire cross section, the ratio of surface hardness to core hardness, and the state of residual stresses, usually compressive, situated within the surface layer relative to the state of stresses in the substrate, which are usually tensile. An incorrectly applied surface layer may cause the formation of a structural flaw in the transition zone of the layer and may be the location of crack initiation, especially by a fatigue mechanism (Ref 1). Surface engineering encompasses various process technologies and also the service properties of products, surface-layer investigation methodologies, and design aspects of the substrate-layer system. In terms of service properties of products, the functionality of surface treatment may be assessed by defining the fatigue limit, wear resistance, or corrosion resistance. Such evaluations are usually performed on specimens in laboratory conditions. However, the most valuable information is to be gained from actual service trials, which also unfortunately may be costly.

Fatigue Resistance Fatigue resistance of machine components is a function of their design, material and

technological parameters, as well as the type of loading in service conditions. When discussing the problem of fatigue resistance, one should consider in detail the effect of these parameters and, in the case of loads, define fatigue characteristics (e.g., plots for different types of loading, such as bending, tensile, and torque). These problems have been sufficiently dealt with in the technical literature. In the process of searching for methods to increase fatigue resistance, there are some constant elements that have a favorable effect, including enhancing treatments such as thermal, thermochemical, as well as surface work hardening. In order to increase fatigue resistance, it is not sufficient to apply a chosen enhancement treatment. Rather, it is important to select the appropriate initial volume heat treatment prior to successive surface, thermal, and workhardening treatments. The problem of enhancing fatigue resistance of machine components by technological methods involves the application of not one chosen treatment but a cycle of successive treatments. The appropriate selection of these treatments affects the structural flaw formed in the process of enhancement, which has a decisive influence on fatigue resistance (Ref 1). A structural flaw occurs in all locations where, as the result of heat treatment (e.g., induction hardening), thermochemical treatments (carburizing, nitriding, etc.), or work-hardening treatments (burnishing, shot peening) of machine components, the layer formed in these processes has different physical-chemical properties than that of the core due to a large gradient of property changes. The value of the structural flaw coefficient, bs, depends on the type of material and the parameters of the technological processes that cause this structural flaw to form. In other words, it depends on heat treatment and surface hardening. Thermal and surface work-hardening treatments used industrially cause enhancement of fatigue resistance. Based on research carried out

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by the Institute of Precision Mechanics, it can be accepted that the fatigue limit (s 1) rises 15 to 30% on average as a result of implementation of such treatments. Enhancement of fatigue resistance is obtained by structural changes, strengthening, and favorable distribution of residual stresses, which are formed as a result of thermal and surface treatments. Due to physical processes taking place within the material during the application of surface, thermal, and work-hardening treatments, changes in microstructural and mechanical properties arise between the surface layer and the core of the material. The gradient of changes of physical-chemical properties depends on the selected technological process and its parameters. Numerous examples have been noted where the fatigue crack origins were traced to the transition zone between the hardened surface layer and the substrate. Figure 1 shows a fatigue fracture with the origin located under the hardened layer at the point where stresses mounted. The mounting of stresses occurs as the result of residual stresses created during heat treatment combined with external stresses. Fatigue Resistance of Steel after Nitriding and Related Nitriding Treatments. The significance and detailed assessment of the effect of a structural flaw are explained, using investigations of the effect of variable core conditions on fatigue resistance as an example. Reference 2 defines the effect of tempering temperature and time of nitriding on the

rotational-bending fatigue resistance of 40HMgrade steel (AISI 4140). This structural steel is commonly used for the manufacture of various types of machine components (e.g., gears, crankshafts). Specimens used for the study were quenched and tempered to the following hardness levels: 30 to 32 HRC, 33 to 34 HRC, and 35 to 36 HRC, applying tempering temperatures within the range of 550 to 620 C. Nitriding was carried out in a controlled process at a temperature of 530 C for 4 to 16 h. Investigations encompassed metallurgical characterization of nitrided layers as well as determination of fatigue resistance (s 1). Moreover, the yield strength (R0.2) was determined in conditions of shear bending (Rg0.2). For the metallurgical investigations pertaining to surface hardnesses, hardness traverses, layer thickness, and microstructure, the Neophot 30 metallograph and the Zwick microhardness tester were used. Investigation of fatigue resistance was carried out with the aid of the PUNZ machine, manufactured by Schenk. The loading frequency was 100 Hz, and the investigations covered 107 cycles. The value of the fatigue limit was calculated by the Dixon-Mood method. Results of fatigue tests of nitrided specimens were compared with results obtained for the same steel (40HM grade) quenched and tempered to a hardness of 30 to 32 HRC. Fractographic investigations of fatigue fractures and determination of the chemical composition of visible inclusions on these fractures were carried out by scanning electron microscopy. Yield strength values were determined in cases of static bending with the aid of the Instron TT-DM machine. The results of metallurgical and strength investigations are shown in Tables 1 and 2. Table 1 Metallurgical characteristics and corresponding values of fatigue limits for the investigated versions of technological processing Case depth

Fig. 1

Fatigue source located in the transition zone between case and core

Surface hardness

Temper Total Effective at core temper- Core ature hardness, Nitriding at 500 +50 HV °C mm HRC time, h HV, mm

HV1

HV10

s 1 MPa

... 4 16 4 8 16 4 16

... 686 743 752 777 786 772 778

... 642 657 695 707 701 698 699

550 735 745 720 725 777 820 840

620 620

30–32 30–32

590

33–34

550

35–36

... 0.13 0.26 0.14 0.19 0.27 0.16 0.30

... 0.24 0.46 0.22 0.29 0.43 0.2 0.42


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Investigations of fatigue resistance (Table 1) showed that nitriding, independently of process time (case depth and core hardness), caused an increase in the fatigue limit in comparison with quenching and tempering. The smallest increase in fatigue limit value (by 30%) was obtained on specimens tempered at 590 C and nitrided for 4 h. The greatest increase in the fatigue limit (49 to 53%) was obtained on specimens tempered at 550 C and nitrided for 4 and 16 h. For the remaining versions of technological processing, the increase in fatigue limit was almost identical and amounted to 31 to 35%. Investigations of quenched and tempered and subsequently nitrided specimens showed a variation in the value of the fatigue limit (s 1), depending on the tempering temperature. Nitriding time did not affect the value of the fatigue limit, both after tempering at 550 C as well as at 620 C. Some variation was observed in the case of versions in which the tempering temperature was 590 C. For this tempering temperature, the highest value of the fatigue limit (777 MPa) was obtained on specimens nitrided for 16 h. At the same time, it was established that increasing the time of nitriding from 4 to 8 h did not affect the fatigue limit. The difference of approximately 1% is within experimental error. Extending the time of nitriding from 4 and 8 h to 16 h caused an increase in fatigue limit value by approximately 50 MPa, that is approximately 6.5%. A comparison of specimens tempered at 550 and 620 C shows that lowering the tempering temperature causes an increase in the fatigue limit by approximately 12%. As has already been mentioned, decreasing the nitriding time fourfold, from 16 to 4 h, for a given tempering temperature does not cause any significant changes of the s 1 value. Further comparison with results obtained on specimens tempered at 590 C indicates that temperatures of 550 and 620 C cause a clear variation in the effect of lowering the tempering

temperature, from the point of view of the fatigue limit. A tempering temperature of 590 C is intermediate between the two aforementioned temperatures, at which this effect is manifest only after the application of a longer nitriding time. It should therefore be emphasized that by the appropriate selection of the tempering temperature, it is possible to achieve an increase in the fatigue limit with shorter nitriding times. On the other hand, selection of inappropriate tempering temperatures may cause the inability to achieve an increase in the s 1 value when the nitriding time is too short. An analysis of angular coefficients of simple regressions of the Wo¨hler plots (Table 2) indicates that, for almost all versions, the angle of inclination of the regression plots is similar. By the same token, the sensitivity of the material to a change in loading within limited fatigue strength is not connected with the time of nitriding. Only specimens tempered at 550 C and nitrided for 16 h were characterized by a slightly higher sensitivity to a change in the level of loading. By the same token, an increase in the level of loading within limited fatigue strength causes a more significant decrease in the number of loading cycles to failure, compared with other technological versions. Thus, shortening the nitriding time coupled with a lower tempering temperature is favorable even when the loading level exceeds the fatigue limit. Fractography of specimen fatigue fractures showed they are of a fine-grained character, with the exception of the middle zone, which has a differing, coarse-grained structure (catastrophic failure zone). The clearly observed fatigue sources occur in the form of fisheyes (Fig. 2).

Table 2 Equations of regression for the investigated technological versions Tempering temperature, °C

Nitriding time, h

620

4 16 4 16 4 16

590 550

Regression equations

s 1 = 141.2 lgN+1567.9 s 1 = 62.4 lgN+1146.6 s 1 = 58.6 lgN+1092.7 s 1 = 67.7 lgN+1201.5 s 1 = 137.2 lgN+1658.3 s 1 = 27.6 lgN+1025.8

Fig. 2

Appearance of fatigue source. Original magnification: 100·

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Initiation of fatigue fractures occurs at the border between the layer and the substrate, in a zone characterized by hardness higher than that of the core by 50 HV. The features initiating fatigue fractures are nonmetallic inclusions of calcium sulfides (Fig. 3), originating from the metallurgical process. Results of investigations of yield strength in bending indicate that variation of both the tempering temperature as well as the nitriding time affects its value insignificantly, not exceeding 7%. The values of Rg are all within the range of 1850 to 1980 MPa. Results of metallurgical evaluations of nitrided layers (Table 1) show that extension of the nitriding time from 4 to 16 h causes an almost doubled growth of the effective minimum tendency to increase the case depth with a lowering of the tempering temperature. A similar but clear trend is observed in the surface hardness of layers, where lowering of the tempering temperature from 620 to 550 C with a 4 h nitriding time cycle causes a rise in HV1 surface hardness by approximately 100 units and in HV10 hardness by 50 units. Figure 4 shows the changes taking place in the character of the microhardness profile as a function of process time, with tempering temperature at 590 C. As can be seen, extension of the nitriding time causes a drop in the angle of inclination of the microhardness profile. Similar changes in the microhardness profile versus process time occur for the remaining tempering temperatures.

Fig. 3

Precipitation on the fracture surface of a specimen which served as the source of the fatigue fracture. Original magnification: 500 ·

Fatigue Evaluation of Nitrided Steels An evaluation of fatigue properties requires an understanding of the mechanism. This mechanism also depends on the condition and mutual relationship between the substrate and the surface layer. This requires determination of the strength condition of the system: substratesurface layer, as a function of external loading. Among the most important parameters describing the condition of the surface layer are microstructure, degree of strengthening, state of stresses, and roughness. Other important parameters include texture, surface energy, and chemical composition (Ref 2, 3). In engineering practice, usable properties such as tensile strength, Rm and fatigue limit, may be determined as functions of mechanical parameters, that is, hardness, H, or tensile strength, Rm, or surface roughness (Ref 4). Such correlations maintain their validity for a material homogeneous throughout its cross section. For heterogeneous materials (e.g., ones that have been surface treated), such correlations cannot be applied directly. The character of distribution

Fig. 4

Microhardness traverses for different nitriding process times

Fig. 5

Fatigue characteristics for 1, a homogeneous material, and 2, a heterogeneous material

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of basic mechanical properties (hardness, residual stresses, as well as their significant effect on fatigue properties) vary in different ways for homogeneous or heterogeneous materials (Fig. 5). Homogeneous materials (curve 1, Fig. 5) have constant properties through the cross section. Heterogeneous materials (curve 2) have variable properties. The fatigue characteristic, which is shown by the distribution of the fatigue limit, s, is a function of hardness, H, and residual stresses, sr: s=f ðH,sr Þ

(Eq 1)

The method of designing a usable characteristic for the fatigue limit distribution has been described in publications (Ref 3–5). The distribution of stresses from extraneous loading constitutes a significant characteristic, because it enables the determination of the strength condition for the surface layer. Loading characteristics are typical distributions of stresses across the section of a component or specimen for the investigated types of loading: tensile-compressive, bending, or torque. For smooth specimens, their determination does not present any problems. Some difficulties may arise when determining the distribution of stresses in a notched specimen. Distribution of stresses from extraneous loading for notched specimens in conditions of bending is (Ref 3): 2 x 3a2 smax ð xÞ=sn a 17 d

(Eq 2)

where smax(x) is the value of local stress at a distance of (x) from the surface, sn is the nominal stress, a is the coefficient of stress concentration, and d is the cross-sectional dimension.

Knowledge of the fatigue characteristic, as well as the loading characteristic, allows the determination of the fatigue strength condition for any location on the component cross section: s1 is0z i

(Eq 3)

where s l is the fatigue limit at any location on the specimen cross section, and si0z is the value of stresses from extraneous loading at the given location i. The investigations were conducted on structural steels 40HM (4140) and 38HMJ (Nitralloy 135M). The steels were hardened and tempered prior to nitriding at two temperatures: 550 and 620 C (Ref 1). Nitriding was carried out using two types of atmosphere, that is, NH3-NH3(diss) and NH3-N2. In the nitriding processes, the atmosphere gas composition was varied, as were the time of nitriding (4 and 16 h) and the nitriding potential, KN (from 1.65 to 4.8). Fatigue resistance tests were carried out on the PUNZ machine (manufactured by Schenck), applying rotational bending stresses with a notch (a = 1.02). The specimen diameter was W = 5.88+0.02 mm. The results of the fatigue resistance tests, metallurgical evaluation and process parameters are shown in Table 3. Figure 6 shows microhardness traverses in the nitrided case for 40HM-grade steel, while Fig. 7 shows the same for the 38HMJ grade. Test results show that the tempering temperature has an effect on the properties of the nitrided case. The effect of the tempering temperature on the basic properties of the nitrided case depends on the steel grade. A higher increase in hardness (by approximately 50%) as well as in case depth is observed for the lowalloy chromium steel 40HM. As can be seen in Fig. 7 for the 38HMJ steel, the effects of tempering temperature on the

Table 3 Technological parameters and test results HV0.5 hardness Tempering temperature, °C

40HM (4140)

38HMJ (Nitralloy 135M)

Steel grade

Nitriding time, h

Core hardness HV0.5

Max on cross section

On surface

550 550 620 690

4 16 4 16

402 396 343 343

677 642 715 343

550 550 620 620

4 16 4 16

356 343 318 296

1030 1030 1030 1030

Residual stresses Fatigue limit (s 1), MPa

At surface, MPa

Depth at which stress changes sign, mm

757 757 826 642

820 840 735 745

600 650 600 900

0.32 0.52 0.37 0.55

1373 1227 1273 1304

805 785 766 810

900 600 450 800

0.25 0.48 0.30 0.45


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changes in hardness and case depth are smaller than for the 4140 steel. It was found that lowering the tempering temperature raises the fatigue resistance of the core-case system. Shorter nitriding time, following a low tempering temperature, does not affect the volume fatigue resistance (Ref 6). The effect of heat treatment of the core on fatigue resistance is shown in Fig. 8. From the illustration, it is seen that raising the core hardness moves the fatigue resistance characteristic, that is, the distribution of the speed limit value, across the section in the direction of higher stress values. Data to determine the characteristics in Fig. 8 are shown in Table 4. To calculate the fatigue limit, the following formula was used: s1 =1:98H70:0011ðHVi Þ2

(Eq 4)

where HV is Vickers surface hardness, and HVi is Vickers hardness at i location on the cross section. The relationship is valid for a hardness range of 340jHVj900. A significant increase in the fatigue limit value (to 820 MPa) with a tempering temperature of 550 C and up to 735 MPa with a tempering temperature of 620 C, relative to prior values of 618 and 550 MPa, determined at the location of fatigue crack initiation (Fig. 9) (on an average 0.5 mm from the surface) should be interpreted as the favorable effect of compressive stresses in the nitrided case (Ref 7). It follows from Fig. 8 that fatigue crack initiation of nitrided cases (compare with Fig. 7) takes place under the surface because stresses

Fig. 6

Microhardness traverses across a nitrided case on 40HM (4140)-grade steel. 1, tempering temperature 550 h; 2, tempering temperature 550 C, time 16 h; 3, tempering temperature 620 C, time 4 h; 4, tempering temperature 620 C, time 16 h C, time 4

from extraneous loading locally exceed the value of the fatigue limit, and, in accordance with curve 3 in Fig. 8, material decohesion must occur.

Fatigue Property Characteristics after Carbonitriding In most modern methods of manufacturing, it is recommended that the design stage consider the different manufacturing technologies. In connection with that, there is an urgent need to determine material characteristics, especially of materials after the application of modern technological property-enhancing treatments. There also exists the need to develop calculation methods of fatigue resistance after thermal and thermochemical treatment. This, however, is the next stage of activity and is possible to carry out only when the basic fatigue properties of the steel following heat treatment are known. The carbonitriding treatment is used for components exposed to lighter loads and subjected to wear as well as bending (Ref 8–10). For those components, which are subjected to contact fatigue during service, the case depths are designed deeper. For the present series of tests, a

Fig. 7

Microhardness traverses across a nitrided case on 38HMJ (Nitralloy 135M)-grade steel. 1, tempering temperature 550 C, time 4 h; 2, tempering temperature 550 C, time, 16 h; 3, tempering temperature 620 C, time 4 h; 4, tempering temperature 620 C, time 16 h

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Fig. 8

Distributions of fatigue limit, curves 1 and 2; residual stresses, curve 3 (550 C) and curve 4 (620 C); extraneous loading, curves 5 and 6, 40HM (4140)-grade steel

Table 4 Fatigue limit values across the specimen section Experimental results Fatigue limit (s 1), MPa

Core At surface At location of initiation of fracture

Theoretical results

Tempering temperature 550 °C

620 °C

550 °C

620 °C

613 ... 820

550 ... 735

618 852 618

550 819 550

case depth of 0.7 mm was selected. The optimal microstructure of carbonitrided components is fine acicular martensite with a small amount of retained austenite and containing no coarse carbide precipitations (Ref 8–10). Specimens prepared to meet the aforementioned conditions were subjected to rotationalbending and one-point bending fatigue tests. Such types of loading were selected based on the premise that bending is the most common method of loading during service, as well as the fact that it is during bending that one can observe

Fig. 9.

Location of fatigue crack initiation on nitrided 40HM (4140)-grade steel. Original magnification: 100·

the most favorable and strongest effect of surface strengthening. Simplified Smith curves were plotted to determine the fatigue resistance for at least three

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methods of bending (the Wo¨hler curve), as well as static strength and yield strength for a given type of loading and materials. The fatigue tests were carried out with the following coefficients of cycle asymmetry: R = 1, R = 0.1, and R = 0.3. The coefficients of 0.1 and 0.3 were selected to ensure the possibility of running the tests only within the range of one-sided bending stresses. The fatigue characteristic was developed for a material in the quenched and tempered condition as a reference and, for materials with a diffusion case, heat treated to the same condition as the heat-treated-only version. Once these values were known, the surface coefficient of strengthening was determined from the equation: m=

sww 1 s1

(Eq 5)

ww where s 1 is the fatigue limit of the enhanced specimen, and s 1 is the fatigue limit of the reference specimen. Rotational-bending fatigue tests were carried out on the Schenck fatigue machine with a constant distribution of the bending movement along the length of the specimen. The frequency was 100 Hz. One-point bending tests were carried out on the Amsler machine. In order to obtain the bending effect on this machine, a prototype addition was designed that, through a lever, allows loading of the tested section of the specimen under a constant bending moment (Fig. 10). The frequency was 150 Hz. Tests were carried out to NG = 107 cycles. The material used in these tests was the 18HGT grade, normalized, with a fine-grained

Fig. 10

Schematic of equipment for fatigue testing in the rotational-bending mode

ferritic-pearlitic microstructure. The chemical composition is given in Table 5. Tests were also carried out on carbonitrided specimens. Carbonitriding of specimens from 18HGT-grade steel was carried out at a temperature of 860 C in an endothermic atmosphere enriched with ammonia and natural gas. Metallurgical evaluations were carried out on 18HGT steel in the quenched and tempered only and carbonitrided condition. In the normalizedonly condition, the specimens showed a ferriticpearlitic microstructure with very fine-grained pearlite (Fig. 11). The microstructure of specimens with diffusion cases was determined based on micrographs and microhardness measurements. Specimens made from 18HGT steel, after carbonitriding and quenching and tempering, exhibit a microstructure of tempered martensite in the subsurface zone (Fig. 12) and a bainiticmartensitic microstructure in the core (Fig. 13). The microstructure of the core was 550 HV0.1. To determine the fatigue resistance, a static bending test was carried out. For the heat treated Table 5 Chemical composition of specimens prepared from steel grade 18HGT Specimen No./diameter, mm

1ø 12 2 ø 12 3 ø 12 1 ø 14 2 ø 14 3 ø 14

Fig. 11

Chemical composition, % C

S

Mn

Cr

Si

Ni

Cu

Ti

0.187 0.185 0.189 0.216 0.225 0.215

0.014 0.015 0.014 0.009 0.009 0.009

0.89 0.87 0.89 1.00 1.02 0.98

1.07 1.12 1.12 1.00 1.02 1.02

0.26 0.28 0.32 0.35 0.34 0.33

0.06 0.06 0.06 0.13 0.13 0.12

0.13 0.12 0.12 0.13 0.12 0.14

0.05 0.05 0.05 0.07 0.08 0.08

Microstructure of heat-treated-only 18HGT-grade steel. Etched with nital. Original magnification: 500·


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(normalized)-only steel, it was not possible to obtain a fatigue resistance value because of the ductility of the material. Only the yield strength was determined, and it amounted to 822.8 MPa on specimens. The static bending test carried out on carbonitrided specimens is shown in Fig. 14. In this case, it was not possible to determine the yield strength, and only the relative bending strength was established as 2289.8 MPa.

In order to obtain a full characteristic of the properties of surface-enhanced materials and of the reference version, a static tensile test was performed. The average value of tensile strength Rm for specimens without a diffusion case was 599.8 MPa, and yield strength was determined as Re = 430 MPa. Similarly, as in the case of the technological bending test, it was not possible to determine the yield strength for the carbonitrided material (Fig. 15). An analysis of static test results delivers new data. The tensile plot for the carbonitrided material is characteristic of brittle materials. There is no necking and no elongation of the specimen. Similar behavior was noted when the bending strength test was performed. In neither case was it possible to determine the yield strength. Based on fatigue tests for rotational bending, which were performed on specimens made from carbonitrided and heat treated (normalized)only material, it was possible to determine the coefficient of surface strengthening, that is, the ratio of m = s 1 (with diffusion case) to s 1 (with no case). For 18HGT steel after carbonitriding, this coefficient was 2.48. With the aid of results obtained in static and fatigue tests for the case of two- and one-side

Fig. 12

Microstructure of carbonitrided case on specimen made of carbonitrided 18HGT-grade steel. Etched with nital. Original magnification: 500 ·

Fig. 13

Microstructure of core of specimen made of carbonitrided 18HGT-grade steel. Etched by nital. Original magnification: 500 ·

Fig. 14

Plot of static bending test of carbonitrided 18HGTgrade specimen

Fig. 15

Plot of static tensile test of carbonitrided 18HGTgrade specimen

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bending, simplified Smith curves were plotted for the heat-treated-only material (Fig. 16) and for the carbonitrided material (Fig. 17). To plot the chart, values of unlimited fatigue resistance were used from the Wo¨hler curves. The upper limit of the chart for the heat-treated-only material is the yield strength obtained from the bending test. The Smith curve for materials after thermochemical treatment differs from that for the reference heat-treated-only material, because its upper limit is determined by the bending yield strength, Rg (Fig. 17). The designer of the component, basing his design on the presently available tables containing data of the ultimate properties of the steel after hardening and fatigue properties for alternating stresses, creates a design that consumes large amounts of material. As the result of using values of s 1 taken from catalogues, the strength of the assembly is compromised. The values for fatigue resistance, sgR, are much higher, which can be seen from the Smith plot (Ref 11). Based on the results of fatigue resistance tests for carbonitrided 18HGT steel, shown for comparison in Table 6, it can be concluded that carbonitriding ensures good strength properties. Summarizing the results of the tests presented in the form of Smith plots, the designer, using

the characteristics of the steel achieved after thermochemical treatment, will be able to reduce the amount of material used and to choose an optimal technological version, thus ensuring high parameters and longer life of the designed component.

Summary Testing of properties of structural materials is essential for development and manufacturing

Fig. 17

Simplified Smith plot for 18HGT-grade steel after carbonitriding

Table 6 Comparison of fatigue properties of 18HGT-grade steel after different types of thermal and thermochemical treatment Heat Treatment Strength parameter

Fig. 16

Simplified Smith plot for 18HGT-grade steel after normalizing

Yield strength (Re), MPa Tensile strength (Rm), MPa Relative elongation (A5), % Necking (Z), % Bending yield strength (Rg0.2), MPa Bending strength (Rg), MPa Bending fatigue limit (s 1), MPa Surface-strengthening coefficient, m

Normalizing

Carburizing

Carbonitriding

...

...

1302.3

1303.2

30.45

...

...

70.06 822.8

... ...

... ...

...

2604.1

2889.8

358.0

745.0

887.1

430 599.8

1

2.08

2.48


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as well as for the practice of design and construction of machines. A large number of assessment criteria for materials properties, as well as methods and testing equipment, stem from the multitude of different mechanisms for failure of structural materials from which components, assemblies, and whole machines are made. Among the most important of these are failures due to the action of static loading. There is, however, an entire spectrum of different types of dynamic loading, as well as volume fatigue, impact loading, contact stresses, destruction by wear of mating surfaces in friction, and others that affect the performance of engineering components. Processes connected with the calculation and design of machine components and assemblies call for a database of materials properties. Modern computational methods and their constant development force the necessity of determining an ever growing number of parameters that describe the properties of materials. A good example of this is the attempt to describe the mechanism of material failure through the coefficient of stress intensity (KIC), both when the failure takes place under variable loads (fatigue) as well as when it occurs due to wear in processes involving friction. Computation of the life of machine components in conditions of variable loading calls for information not only about the value of the fatigue limit but also about the angle of inclination of the straight line in the range of limited fatigue strength, and also about the parameters of the bend point of the fatigue curve. Modern structural material does not need to be homogeneous throughout its section. A great number of steels, plastics, and other metallic materials call for enhancement of the surface, due to the constant quest for decreasing material and energy consumption as well as increasing properties. Surface layers are, in the majority of cases, superficially hardened layers formed by thermal and thermochemical treatment or other enhancement technologies, such as surface work hardening and anticorrosion coatings. Testing of materials properties after heat treatment shows that the achievement of desired service properties is connected with the appropriate selection of parameters not only in the final thermochemical treatment but also in the prior volume heat treatment. In the case of nitriding of machine components, this technology is usually preceded by quenching and tempering at a minimum temperature 20 C above

the subsequent nitriding temperature. Initial hardening by quenching and tempering is critical to core hardness and to the properties of the nitrided case and affects the fatigue resistance of the material after nitriding. An analysis of fractured surfaces of nitrided specimens, exposed to service in conditions of rotational bending, revealed that the weakest location on the specimen cross section is the zone of transition of the nitrided case to core. The method of designing surface cases enables an explanation of the root cause of fatigue crack initiation under the nitrided case. The fatigue limit in the cross section of the specimen was described as a function of microhardness and residual stresses. The initiation of fatigue cracks takes place in the location where stresses from extraneous sources exceed the value of the fatigue limit, which is obviously in agreement with conditions of strength. The favorable effect of core hardness on fatigue resistance was observed. The investigations showed good fatigue properties of cases obtained by carbonitriding, as well as a lowering of ductility of these cases. Results of investigations, presented in the form of Smith plots, confirm the necessity of further pursuing investigations in this field. In designing practice, the application of obtained results is the least expensive method of lowering material consumption and enhancing the life of machine components. It was established that fatigue resistance is significantly affected not only by compressive stresses but also by tensile stresses. Among the parameters describing the state of residual stresses, the distribution of stresses was of more significance, followed by the value of the residual stresses. Models of surface layers described in literature are difficult to implement in industrial practice. There appears to be a need for the creation of such a model of the surface layer. This could be described by parameters that can be used in strength calculations and that would allow its application in instances of different types of extraneous loading, depending on the type of service of the component. This model could become the basis for predicting the state of the surface layer, based on required usable properties of machine components. Work on such a model is carried out in two directions:

Based on experimental description of the state of the surface layer through hardness

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traverses, residual-stress distributions, and distributions of element concentrations, for example, carbon and nitrogen Based on a description of the material through the theory of elasticity and solving constitutive equations by numerical methods

Surface layer design for the criterion of fatigue failure is based on a comparison of the local fatigue resistance with local stresses occurring at critical locations in the investigated component. Contemporary machines and designs should be characterized by required life and reliability, featuring a sufficient life between overhauls, depending on the type of service, while at the same time fulfilling the requirements for ecology and ergonomics. Such parameters should be attained concurrently with a reduction of material and energy consumption during manufacture and service. This task may be achieved only when modern computational methods are implemented along with modern technology and proper service conditions at each stage of the product life, that is, study phase, design, manufacture, service, and recycling. The implemented computational methods enable the design of products according to strength criteria, somewhat less often according to tribological criteria, and least often pertaining to corrosion. Contemporary machine components and assemblies are subjected in service to the joint interaction of strength, tribological, and corrosion hazards. On the other hand, the implemented computational methods enable the design of products with one selected mode of failure. In the construction of machine components, there are many parts (crankshafts, threaded joints, springs) that are concurrently exposed to different types of failure hazards during service: mechanical, tribological, or corrosive. Similar elements of construction (bridges, masts, cables, earth-moving and mining machines) are exposed to concurrent hazards of fatigue-type stresses and corrosion. Classical strength or tribological calculations do not take into account the factor of time. During service, due to the processes of fatigue, tribological, or corrosive deterioration, there occurs a change in the properties of the system being evaluated. Tribological and corrosive processes cause a change in the geometry and surface condition of the component. This, in turn, causes a change in the state of stresses in working systems, affecting their life and reliability.

Therefore, the development of failure criteria, taking into account the joint effect of an accumulation of damage due to the working of alternating loads, wear by friction, and the action of corrosion, is a very important task, because the determination of the criteria for failure will enable proper selection of surface layers for the given service condition.

REFERENCES

1. A. Nakonieczny, Podwyz˙szenie wytrzymałosći zme˛czeniowej cze˛sći maszyn przez obro´bke˛ cieplna˛ i powierzchniowa˛ obro´bke˛ plastyczna˛ (Enhancement of Fatigue Strength through Heat Treatment and Surface Work Hardening), Proc. XXIV Seminar IMP, XI on Metallurgy and Heat Treatment, Oct 23–24, 1984 (Warsaw), IMP (translated from Polish) 2. T. Babul, A. Nakonieczny, and J. Tacikowski, Wpływ umocnienia podłoz˙a na wytrzymałosć´ zme˛czeniowa˛ azotowanej stali 40HM (The Effect of Core Strengthening on the Fatigue Resistance of Nitrided 40HM Grade Steel), Proc. III Polish Scientific Conference on Surface Treatment, Cze˛stochowa-Kule, Politechnika Lodz, 1996 (translated from Polish). 3. A. Nakonieczny and J. Tacikowski, Analiza pe˛kania zme˛czeniowego stali azotowanych (An Analysis of Fatigue Fracturing of Nitrided Steels), Proc. First Polish Scientific Conference on Modern Technology in Surface Engineering, Sept 1994, Politechnika Lodz (translated from Polish) 4. A. Nakonieczny, The Effect of Residual Stresses and Hardness on Fatigue Behavior of Surface Treated Materials, Proc. MATTEC 91, Technology Transfer Series, A. Niku-Lari, Ed., 1991 5. B. Winderlich, Das Konzept der lokalen Deuerfestigkeit und seine Anwendung auf martensitische Randschichten, in bensondere Laserhartungsschichten, Mater. wiss. Werkst. tech., Vol 21, 1990, p 378–389 (in German) 6. V.P. Kogaev, N.A. Machutov, and A.P. Gusenkov, Rascˇety detalej masˇin i konstrukcij na procˇnost´ i dolgovecˇnost´, Masˇinostroenie, Moskva, 1985, p 150–182 (in Polish)

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7. J. Tacikowski and A. Nakonieczny, Report 114,01,0163, IMP, Warsaw, 1992 8. A. Nakonieczny, Dissertation, Russian Academy of Science, Moscow, 1991 9. W. Olszan´ski, More Important Problems Pertaining to the Austenitic Carbonitriding Process, XVI Seminar, VII on Metallurgy and Heat Treatment Book 2, IMP, Warsaw, 1977

10. J. Wyszkowski, Nowoczesne tendencje w zakresie nawe˛glania i we˛gloazotowania gazowego, (Modern Trends in the Field of Gas Carburizing and Carbonitriding), IMP, Warsaw, 1974 (translated from Polish) 11. W. Olszan´ski, I. Sułkowski, J. Tacikowski, and J. Zys´k, Obro´bka cieplno-chemiczna, (Thermochemical Treatment) Book 5, IMP, Warsaw, 1979

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Steel Heat Treatment Failures due to Quenching L.C.F. Canale, Universidade de Saõ Paulo G.E. Totten, Associaçaõ Instituto Internacional de Cieˆncia and Portland State University

QUENCHING is one of the more important heat treating processes, because it is so closely related to dimensional control requirements and control of residual stresses. Quenching is often attributed to many distortion and cracking problems, whether the quenching process is the actual root cause or not. Approximately 20% of the problems in heat treating relate to heating processes, while as much as 80% of the problems relate to cooling processes. This chapter provides an overview of the fundamental material- and process-related parameters of quenching on residual stress, distortion control, and cracking. This overview is followed by various selected case histories of failures attributed to the quenching process.

Phase Transformations During Heating and Quenching Properties such as hardness, strength, ductility, and toughness are dependent on the microstructural products that are present in steel. Typically, the first step in the transformation process is to heat the steel to its austenitizing temperature. The austenitized steel is then cooled rapidly to avoid the formation of pearlite, which is a relatively soft transformation product, and to maximize formation of martensite, a relatively hard transformation product, and to achieve the desired as-quenched hardness. The most common transformation products that may be formed in quench-hardenable steels from austenite are, in order of formation with decreasing cooling rate, martensite, bainite, pearlite, ferrite, and cementite. The formation of these products and the proportions of each are dependent on the time and temperature cooling history of the particular alloy and the elemental

composition of that alloy. The transformation products formed are typically illustrated with the use of transformation diagrams that show the temperature-time dependence of the microstructure formation process for the alloy being studied. Two of the most commonly used transformation diagrams are the time-temperature transformation and continuous cooling transformation diagrams. Time-temperature transformation (TTT) diagrams, also called isothermal transformation diagrams, are developed by heating small samples of steel to the austenite transformation temperature, followed by rapid cooling to a temperature intermediate between the austenitizing and the martensite start (Ms) temperature, and then holding for a fixed period of time until the transformation is complete, at which point the transformation products are determined. This is done repeatedly until a TTT diagram is constructed, such as that shown for an unalloyed steel (AISI 1045) in Fig. 1 (Ref 1). The TTT diagrams can only be read along the isotherms. Continuous Cooling Transformation Diagrams. Alternatively, a given steel may be continuously cooled from the austenitizing temperature at different specified rates. The proportion of transformation products formed after cooling to various temperatures intermediate between the austenitizing temperature and the Ms temperature is used to construct a continuous cooling transformation (CCT) diagram, such as the one shown for an unalloyed carbon steel (AISI 1045) in Fig. 2 (Ref 1). The CCT curves provide data on the temperatures for each phase transformation, the amount of transformation product obtained for a given cooling rate with time, and the cooling rate necessary to obtain martensite. The critical cooling rate is dictated by the time required to avoid formation of

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Martensite

Fig. 1

Time-temperature transformation diagram of an unalloyed steel containing 0.45% C. Austenitizing temperature: 880 C. Source: Ref 1

Martensite

Fig. 2

Continuous cooling transformation diagram of an unalloyed steel containing 0.45% C. Austenitizing temperature: 880 C. Source: Ref 1

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Steel Heat Treatment Failures due to Quenching / 257

Fig. 3

Crystal structures. (a) Austenite, face-centered cubic. (b) Ferrite, body-centered cubic. (c) Martensite, body-centered tetragonal. Source: Ref 1

pearlite for the particular steel being quenched. As a general rule, a quenchant must produce a cooling rate equivalent to or faster than that indicated by the nose of the pearlite transformation curve to maximize the martensite transformation product (Ref 1). The CCT diagrams can only be read along the curves of different cooling rates, and a continuous cooling curve can only be superimposed on a CCT but not on a TTT diagram. Metallurgical Crystal Structure. When steel is slowly cooled, it undergoes a crystal structure (size) change as it transforms from a less densely packed (face-centered cubic) austenite to a more densely packed body-centered cubic structure of ferrite. At faster cooling rates, the formation of ferrite is suppressed, and martensite, which is an even less densely packed body-centered tetragonal structure than austenite, is formed. Illustrations of these crystal structures are provided in Fig. 3 (Ref 1). This results in a volumetric expansion at the Ms temperature as shown in Fig. 4 (Ref 1). Figure 5 shows that the crystal lattice of austenite expands with increasing carbon content (Ref 2). It has been reported that typically when a carbide-ferrite mixture is converted to martensite, the resulting expansion due to increasing carbon content is approximately 0.05 mm/mm (0.002 in./in.) at 0.25% C and 0.18 mm/mm (0.007 in./in.) at 1.2% C (Ref 2). The fractional increase in size when austenite is converted to martensite is approximately 0.36 mm/mm (0.014 in./in.) for eutectoid compositions. This illustrates the effect of carbon structure and steel transformation on residual stresses and distortion leading to dimensional changes. Estimation of Volumetric Change due to Steel Transformation upon Quenching. Various microstructures are possible upon

Fig. 4

Steel expansion and contraction upon heating and cooling. Source: Ref 1

quenching of steel, and the potential microstructural transformations that are possible for a given steel are illustrated by their CCT or TTT diagrams. Furthermore, dimensional changes depend on carbon content and the microstructural transformation product formed. Table 1 summarizes the atomic volumes of various microstructural components as a function of carbon content (Ref 3). Table 2 provides an estimate of volumetric changes as a function of carbon

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content for various metallurgical transformations (Ref 4, 5). Thelning reported that volumetric expansion occurring as a result of quenching could be estimated from (Ref 6): DV=V · 100=(100 Vc Va ) · 1:68C +Va ( 4:64+2:21C)

(Eq 1)

where (DV/V) · 100 equals the percentage change in volume, Vc equals the percentage by volume of undissolved cementite, (100 Vc Va) equals the percentage by volume of martensite, Va equals the percentage by volume of austenite, and C equals the percentage by weight of carbon dissolved in austenite and martensite. Berns reported that if the value of (DV/V) is known or can be computed, internal stresses that are developed in a part due to temperature differences (DT) arising from either

one-dimensional heating or cooling could be estimated from (Ref 7): s=E e=E 1/3 (DV=V)=E a DT

(Eq 2)

where E (modulus of elasticity) = 2 · 105 N/mm2 and a (coefficient of thermal expansion) = 1.2 · 10 5. Relative volume changes due to phase transformation are illustrated in Fig. 6 (Ref 7). Kunitake and Susigawa (Ref 8) reported that the tendency for cracking decreases as the start of the martensite transformation temperature (Ms) increases. The Ms temperature was approximated from: Ms ( C)=521 353C 225Si 24:3Mn 27:4Ni 17:7Cr 25:8Mo

(Eq 3)

The correlation between the occurrence of quench cracks and Ms temperature is shown in Fig. 7. A similar study produced a poor Table 2 Volumetric changes with various steel transformations Steel transformation

Pearlite?austenite Austenite?martensite Austenite?acicular lower bainite Austenite?feathered upper bainite

Volumetric change

4.64+2.21 C(a) 4.64 0.53 C(a) 4.64 1.43 C(a) 4.64 2.21 C(a)

(a) Percent carbon. Source: Ref 4, 5

Fig. 5

Carbon content versus lattice parameters of (retained) austenite and martensite at room temperature. “a” at the top of the graph is the lattice parameter of face-centered cubic austenite. a and c in the lower half of the graph are the two lattice parameters of tetragonal martensite. The ratio of c/a for martensite as a function of carbon content is also given. Source: Ref 2

Table 1 Atomic volume of various microstructural constituents of ferrous alloys Phase

Ferrite Cementite Ferrite+carbides Pearlite Austenite Martensite (a) Percent carbon. Source: Ref 3

˚3 Apparent atomic volume, A

11.789 12.769 11.786+0.163 C(a) 11.916 11.401+0.329 C(a) 11.789+0.370 C(a)

Fig. 6

Specific volume (DV/V) of carbon steels relative to room temperature. Source: Ref 7


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correlation between grain size and quench cracking, as shown in Fig. 8 (Ref 8). Kunitake and Sisigawa (Ref 8) developed a relationship to interrelate the combined effect of both carbon content and elemental composition on cracking propensity. This was designated as the carbon equivalent (Ceq), and it is calculated by: Ceq =C+Mn=5+Mo=5+Cr=10+Ni=10

(Eq 4)

Figure 9 shows a good correlation between the carbon equivalent and steel cracking. In general, steels are classified as crack sensitive if the Ceq value is greater than 0.52 to 0.55% (Ref 8). Another measure of cracking tendency is the difference in the start and finish temperatures of martensite formation (Ms Mf) (Ref 9).

Fig. 7

A summary of the Ms and Mf values for some common steels is provided in Table 3. The correlation between cracking sensitivity and the transformation temperature range is due in part to the low Mf caused by high-carbon steels (which expand more) and to the fact that wide transformation ranges may result in cracking of the brittle untempered martensite formed at higher temperatures in the transformation range. Fujio et al. (Ref 10) showed that the volumetric expansion caused by martensite formation can be estimated from the maximum cooling rate in a particular type of steel, as shown in Fig. 10. Similar correlations were evaluated for both cooling time and cooling rate at the Ms temperature. However, these correlations were dependent on the cross-sectional size and thus could not be used for gears or other parts with complex shapes. Volumetric expansion can be estimated for various crosssectional sizes by a correlation between the volume fraction of martensite versus the cooling

Relationship between quench cracking frequency and martensite start (Ms) temperature. Source: Ref 8

Fig. 9

Relationship between carbon equivalent (Ceq) and quench cracking frequency. Source: Ref 8

Table 3 Martensite start (Ms) and martensite finish (Mf) values for selected steels AISI No.

Fig. 8

Relationship between quench cracking frequency and austenitic grain size. Source: Ref 8

1065 1090 1335 3140 4130 4140 4340 4640 5140 8630 8695 9442

Austenitizing temperature, °C

Ms, °C

Mf , °C

815 885 845 845 870 845 845 845 845 870 845 860

275 215 340 330 375 340 290 340 330 365 135 325

150 80 230 225 290 220 165 255 240 280 ... 15


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rate at the Ms temperature for the steel (Fig. 11). It is well known that retained austenite can substantially affect distortion. Geller and Brimene (Ref 11) published a nomogram that can

Fig. 10

Fig. 11

be used to predict dimensional changes caused by the total carbon concentration in the martensitic transformation product and the amount of retained austenite. Steel chemical compositions

Relationship between maximum cooling rate and volumetric fraction of martensite. Source: Ref 10

Relationship between maximum cooling rate and the martensite start (Ms) temperature and volumetric fraction of martensite. Points in the same curve are related to different positions in the bar and therefore with the degree of martensitic transformation. Source: Ref 10

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Table 4 Steel chemical compositions listed in Fig. 12 Russian steel designation(a)

Composition, wt% Mo

V

Ti

W

Co

Cu

S

P

U8

0.75–0.84 0.17–0.33 0.17–0.33 0.15 max

...

...

...

...

...

...

...

KhVG

0.90–1.05 0.10–0.40

...

...

...

...

1.20–1.60

...

ShKh15

0.95–1.05 0.17–0.37 0.20–0.40 1.30–1.65

...

...

...

...

...

7KhG2VM

0.68–0.76 0.20–0.40 1.80–2.30 1.50–1.80

0.30 max ...

0.50–0.80 0.10–0.25

...

0.5–0.9

...

7KhG3V Kh12M

0.68–0.76 0.20–0.40 3.0–3.5 1.50–1.80 1.45–1.65 1.10–1.40 0.15–0.45 11.0–12.5

... ...

... 0.10–0.25 0.40–0.60 0.15–0.30

... ...

0.5–0.9 ...

... ...

0.30 max 0.25 max 0.30 max ... ...

4Kh5V2FS

0.35–0.45 0.80–1.20 0.15–0.40 4.50–5.50

...

...

0.60–0.90

...

1.60–2.20

...

3Kh2V8F

0.30–0.40 0.15–0.40 0.15–0.40 2.20–2.70

...

...

...

...

7.50–8.50

...

4Ch5W2FS

0.35–0.40 0.80–1.20 0.15–0.40 4.50–5.50

0.30 max

0.60–0.90

...

0.30–0.40 0.15–0.40 0.15–0.40 2.20–2.70

0.50 max

0.20–0.50

7.80–8.50

...

7.0 5.0

... ...

0.03 max 0.03 max ... ...

1.60–2.20

3Ch2W8F

0.35 max 0.35 max ... ...

14.0 ...

25.0 9.0

0.03 max 0.03 max 0.02 max 0.03 max ... 0.03 max 0.03 max 0.03 max 0.03 max 0.03 max ... ...

0.03 max 0.03 max 0.03 max 0.03 max ... 0.03 max 0.03 max 0.03 max 0.03 max 0.03 max ... ...

R14M7K25 N18K9M5

C

1.0 1.0

Si

... ...

Mn

Cr

0.8–1.10 0.90–1.20

... ...

... ...

Ni

0.30 max 0.30 max 0.30 max 0.03 max ... ...

(a) These compositional data were provided by Dr. Dmitry Wainstein, Surface Phenomena Research Group, Physical Metallurgy Institute, CNIICHERMET, Moscow, Russia.

are found in Table 4. The following comments will assist in interpreting the nomogram shown in Fig. 12:

The shaded line represents zero distortion. Steels with martensitic carbon contents and retained austenite levels falling on the line will exhibit essentially no distortion. Martensitic steels with carbon contents and retained austenite levels that fall bellow the shaded line will exhibit shrinkage upon quenching. Martensitic steels with carbon contents and retained austenite levels that fall above the shaded line will exhibit expansion.

This nomogram was developed for various construction and tool steels. Therefore, it should be used with caution for other steel grades (e.g., high-speed tool steels). Basic Distortion Mechanism. Shape and volume changes occurring during heating and cooling can be attributed to three fundamental causes (Ref 12):

Residual stresses will cause shape change when they exceed the yield strength of the material. Stresses caused by differential expansion due to thermal gradients will increase with the thermal gradient and cause plastic deformation as the yield strength is exceeded.

Volume changes due to transformational phase changes will be contained as residualstress systems until the yield strength is exceeded.

When parts are heated during heat treatment, a thermal gradient exists across the cross section of the component. If a section is heated so that a portion of the component becomes hotter than the surrounding material, the hotter material expands and occupies a greater volume than the adjacent material and will thus be exposed to applied stresses that will cause a shape change when they exceed material strength. These movements can be related to heating rate and section thickness of the component. Volume Changes During Phase Transformations. When a steel part is heated, it transforms to austenite with an accompanying reduction in volume. When it is quenched, the structure transforms from austenite to martensite, and its volume increases. If these volume changes cause stresses to be set up that are constrained within the strength of the material, a residual-stress system is created. If the stresses cannot be contained, material movement will occur, which will cause cracking under extreme conditions. The expansion is related to the composition of the steel. Figure 13 shows the relative volume increase of two steels as a function of austenitizing temperature and specimen dimensions (Ref 13).


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Fig. 12

Changes in linear dimensions during quenching relative to carbon concentrations in martensite and retained austenite. Source: Ref 11

Fig. 13

Volume increase of 90MnV8 and 15CrV6 steels as a function of austenitizing temperature and specimen dimensions. Source: Ref 13

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While each of these phenomena is a wellknown physical change, the situation is made more complex when all three events occur simultaneously. In addition, other events, such as heating rate, quenching, and inconsistent material composition, further complicate the process. Relief of Residual Stresses. If a part has locked-in residual stresses, these stresses can be relieved by heating the part until the lockedin stresses exceed the strength of the material. A typical stress-strain curve obtained from a tension test is shown in Fig. 14 (Ref 12). Initial changes in shape are elastic, but under increased stress, they occur in the plastic zone and are permanent. Upon heating, the stresses are gradually relieved by changes in the shape of the part due to plastic flow. This is a continuous process, and as the temperature of the part is increased, the material yield stress decreases, as shown in Fig. 15 (Ref 14). It is a function not only of temperature but also of time, since the material will creep under lower applied stresses. It is apparent that the stresses can never be reduced to zero, because the material will always possess some level of yield strength below which residual stresses cannot be reduced.

Effect of Materials and Quench Process Design on Distortion Quenchant selection and quenching conditions are critically important parameters in quench system design. For example, one study

compared the distortion obtained with quenching of a 0.4% medium-carbon plain steel bar of 200 mm diameter by 500 mm long in water or oil from 680 C (Ref 15, 16). The results, shown in Fig. 16(a and c), show essentially equivalent variation in diameter and length with both cooling processes, which was due to thermal strains within the steel. Interestingly, the wellknown diameter variations at the end of the bar, known as the end effect, were observed, which is attributable to heat extraction from both the sides and ends of the bar (Ref 1). If the same steel bars of the same dimensions are heated to 850 C to austenitize the steel and then are quenched in water or oil, the results shown in Fig. 16(b and d), respectively, are obtained (Ref 15, 16). Considerably greater dimensional variation and lengthening of the bar (for the oil quench) was obtained due to both thermal and transformational strains within the steel. Thuvander and Melander modeled the dimensional changes of a 70 mm steel (0.15% C, 1% Mn, 0.75% Cr, 0.85% Ni) cube after austenitizing and then quenching in water and oil (Ref 15, 17). The results of this work are shown in Fig. 17. They show that the edges and faces shrink (becoming concave) and the effect is greater when quenched in water than when quenched in oil (Ref 1). Various factors may affect distortion and growth of steel during heat treating. These include component design, steel grade and condition, machining, component support and loading, surface condition, heating and atmosphere control, retained austenite, and the quenching process (Ref 18). Component Design One of the overwhelming causes of steel cracking and unacceptable distortion control is

Fig. 14

Various features of a typical stress-strain curve obtained from a tension test. Source: Ref 12

Fig. 15

Variation of yield strength with temperature for three generic classes of steel. Source: Ref 14

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component design. Poor component design promotes distortion and cracking by accentuating nonuniform and nonsymmetrical heat transfer during quenching. Component design characteristics that are common to distortion and cracking problems include (Ref 19, 20):

Parts that are long (L) with thin (d) cross sections. Long and thin parts are defined as greater than L = 5d for water quenching, L = 8d for oil quenching, and L = 10d for austempering, where L is the length of the parts, and d is the thickness or diameter. Parts that possess large cross-sectional areas (A) and are thin (t), which are defined as A = 50t

Parts that exceed these dimensions must often be straightened or press quenched to maintain dimensional stability (Ref 20). If possible,

Fig. 16

materials with sufficient hardenability should be oil or salt quenched. Design symmetry is also an important variable to minimize distortion. For example, the unsymmetrical gear design shown in Fig. 18(a) may typically undergo distortion, as shown in Fig. 18(b) (Ref 19). (The load on a gear tooth increases by the 4.3 power of the taper, Ref 19). The solution to the gear design problem shown in Fig. 18 is to provide greater symmetry, as shown in Fig. 19. If this is not possible, press quenching or tooth-by-tooth induction hardening may be the only solutions (Ref 19, 20). Another common design problem is parts with holes, deep keyways, and grooves. One illustration of this problem is hardening of a shaft with a lubrication cross hole, as illustrated in Fig. 20 (Ref 19). Preferred alternative designs are also shown in Fig. 20. If a radial cross hole is

Dimensional variation of a medium-carbon (0.4%) steel bar (200 mm diam by 500 mm) after the indicated heat treatments. These bars were quenched vertically with one end down (marked 0 in the figure). (a) and (c) show no transformation, only thermal strain after water quenching from 680 C. (b) and (d) show thermal and transformation strains after quenching from 850 C. Source: Ref 1

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mandatory, the use of a carburized steel with oil quenching would be preferred. Kern and Suess have reported that the size of the tapped holes can be maintained by the insertion of SAE grade 8 set screws or bolts (Ref 21). Prequenching can be used to control the taper of plain holes during heat treatment. Some hole distortion problems may require oil quenching (conventional or hot oil) or austempering. The distortion encountered when quenching a notched part, such as a shaft with a milled slot, is illustrated in Fig. 21 (Ref 20). In this case, nonuniform heat transfer results. The metal within the notch is affected by the shrinkage of the metal around it due to slower cooling within

the slot, caused by vaporization of the quenchant. Therefore, upon cooling, the metal on the side with the shaft is too short, pulling the shaft out of alignment. 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 (Ref 20). Flat plates are also susceptible to distortion upon quenching. If the material is flat and stress free, round or nearly square, and free of decarburization, Kern and Suess have reported a guide (Table 5) to maintain a flat surface (within 0.025 mm, or 0.001/in., of size) if parts are racked and quenched edgewise (Ref 21). Parts exceeding these limits may require press quenching.

Fig. 19

Design solutions to the distortion problem shown in Fig. 18. Source: Ref 19

Fig. 17

Dimensional changes in a 70 mm steel (0.15% C, 1% Mn, 0.75% Cr, 0.85% Ni) bar after austenitizing and then quenching in water or oil. Source: Ref 1

Fig. 18

Schematic of a gear that is difficult to harden without the distortion shown. Source: Ref 19

Fig. 20

Design solutions to the quench-cracking problem often encountered in shaft hardening over a cross hole. Source: Ref 19

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Steel Grade and Condition. Although quench cracking is most often due to nonuniform heating and cooling, material problems may also be encountered. Some typical material considerations include (Ref 19):

The compositional tolerances should be checked to assure that the alloy is within specification. Some alloys are particularly problematic. For example, some steel grades must be water quenched when the alloy composition is on the low side of the specification limit. Conversely, if the alloy composition is on the high side, cracking is more common. Steel grades that exhibit this problem include 1040, 1045, 1536, 1541, 1137, 1141, and 1144. As a rule, steels with carbon contents and hardenability greater than 1037 are difficult to water quench (Ref 19). Some steel grades with high manganese are prone to microsegregation of manganese and gross segregation of chromium and are prone to cracking. These include 1340, 1345, 1536, 1541, 4140, and 4150. If possible, it is often a good choice to replace the 4100

Fig. 21

Distortion often encountered when quenching a notch. Source: Ref 20

series with the 8600 series of steels (Ref 20). “Dirty” steels, those containing greater than 0.05% S, such as 1141 and 1144, are more prone to cracking. The reasons include: greater alloy segregation in dirty steels leads to alloy-rich and alloy-lean regions; there are typically more surface seams that act as stress raisers with dirty steels; and steels with higher sulfur levels are often manufactured to coarse-grain practice for improved machinability, which also imparts greater brittleness and propensity for cracking. Decarburization of up to 0.064 mm/ 1.59 mm (0.0025 in./1/16 in.) diameter may be present.

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 quenchant selection. Table 6 summarizes some steel mean carbon content concentration limits for water, brine, or caustic quenching (Ref 22). 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 (Ref 23). It is important to provide a sufficient forging reduction ratio to allow the carbide formation to become fine and uniform (Ref 24). Since 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 that normalized and subcriticalannealed steel is the ideal condition (Ref 18). Subcritical annealing is performed to relieve stresses incurred during normalization without softening or homogenizing the steel. The subcritical annealing process reduces the carbon content and alloy carbide content in the austenite, allowing the production of more lath Table 6 Suggested carbon content limits for water, brine, and caustic quenching Hardening method/shapes

Table 5 Guide to maintain a flat surface Quenchant

Ratio (max) of perimeter/thickness

Water Oil Austemper Gas

30 80 125 150

Carbon, max%

Furnace hardening General use Simple shapes Very simple shapes, e.g., bars

0.30 0.35 0.40

Induction hardening Simple shapes Complex shapes

0.50 0.33

Source: Ref 22


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martensite in the microstructure, which provides higher fracture toughness and higher impact toughness (Ref 23). Machining. Material removal during machining can result in high residual-stress levels and ultimately unacceptable distortion (Ref 18). When excessive machining stresses are imparted, the process may require modification to include a rough machining, then stress relieving, followed by fine machining. Component Support and Loading. Many parts may sag and creep under their own weight when heat treated, which is an important cause of distortion. An example of a component that is susceptible to such distortion is a ring gear. Dimension limits by which ring gears are classified are provided in Fig. 22 (Ref 18). (A general

Fig. 22

Dimensions of a ring gear shape. Shape limitation: length/wall thickness, j1.5; inside diameter/outside diameter (ID/OD), 40.4. Minimum wall thickness (WT) is ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi p 5 defined by: WT i2.25 · module+0.4 · mod · L · OD3 .

Source: Ref 18

Fig. 23

Classification of shapes. Source: Ref 18

dimensional classification of various distortionsensitive shapes is provided in Fig. 23, Ref 18). Proper support when heating is required to minimize out-of-flatness and ovality problems, which may result in long grinding times, excessive stock removal, high scrap losses, and loss of case depth (Ref 18). To achieve adequate distortion control, custom supports or press quenching may be required. Pinion shafts, as defined in Fig. 23, are susceptible to banding along their length if they are improperly loaded into the furnace, as shown in Fig. 24 (Ref 18). When this occurs, the pinion shafts must then be straightened, which will add to the production cost. Surface Condition. Quench cracking may be due to various steel-related problems that are only observable after the quench, but the root cause is not the quenching process itself (Ref 25). Many of these problems have been reviewed earlier and include prior steel structure, stress raisers from prior machining, laps and seams, alloy inclusion defects, grinding cracks, chemical segregation (bonding), and alloy depletion. In this section, three surface condition-related problems that may contribute to poor distortion control and cracking are discussed: tight scale formation, decarburization, and the formation of surface seams or nonmetallic stringers. Tight scale problems are encountered with forgings hardened from direct-fired gas furnaces

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with high-pressure burners (Ref 20, 22). The effect of tight scale on the quenching properties of two steels, 1095 carbon steel and 18-8 stainless steel, is illustrated in Fig. 25 (Ref 23). These cooling curves were obtained by still quenching into fast oil. A scale of not more than 0.08 mm (0.003 in.) increases the rate of cooling of 1095 steel as compared to the rate obtained on a specimen without scale. However,

a heavy scale (0.13 mm or 0.005 in. deep) retards the cooling rate. A very light scale (0.013 mm or 0.0005 in. deep) also increased the cooling rate of the 18-8 steel over that obtained with the specimen without scale. In practice, the formation of tight scale will vary in depth over the surface of the part, resulting in thermal gradients due to differences in cooling rates. This problem may yield soft spots and uncontrolled distortion and is particularly a problem with nickel-containing steels. Surface oxide formation can be minimized by the use of an appropriate protective atmosphere. The second surface-related condition is decarburization, which may lead to increased distortion or cracking (Ref 24). At a given depth within the decarburized layer, the part does not harden as completely as it would at the same point below the surface if there were no decarburization. This leads to nonuniform hardness, which may contribute to increased distortion and cracking because (Ref 20):

Fig. 24

Fig. 25

Typical pinion shaft distortion due to furnace loading. Source: Ref 18

The decarburized surface transforms at a higher temperature than the core (the Ms temperature decreases with carbon content). This will lead to high residual tensile stresses at the decarburized surface or a condition of unbalanced stresses and distortion. Since the surface is decarburized, it will exhibit lower hardenability than the core. This will cause the upper transformation products to form early, nucleating additional undesirable products in the core. The

Centerline cooling curves showing the effect of scale on the cooling curves of steels quenched in fast oil without agitation. (a) 1095 steel. Oil temperature: 50 C (125 F). (b) 18-8 stainless steel. Oil temperature: 25 C (75 F). Test specimens were 13 mm diam by 64 mm long (0.5 by 2.5 in.). Source: Ref 23

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decarburized side will be softer than the side that did not undergo decarburizing, which is harder. The greater amount of martensite leads to distortion.The solution to this problem is to restore carbon into the furnace atmosphere or machine off the decarburized layer. The third surface-related condition that may lead to cracking or material weakening is the formation of surface seams or nonmetallic inclusions, which may occur in hot rolled or cold finished material. The presence of these defects prevent the hot steel from welding to itself during the forging process, for example, creating a stress raiser. To prevent this problem with hot rolled bars, stock should be removed before heat treatment. Recommendations made earlier by Kern are provided in Table 7 (Ref 22). Although not a published standard, Kern has reported that a seam or nonmetallic depth of 0.025 mm/3.3 mm (0.001 in./0.13 in.) diameter maximum is usually acceptable for cold finished bars (Ref 22). If the seam depth is excessive, it is recommended that the bars be magniflux inspected prior to heat treatment. Heating and Atmosphere Control. An important source of steel distortion and cracking is nonuniform heating and not using the appropriate protective atmosphere. For example, if steel is heated in a direct-gas-fired furnace with high moisture content, the load being heated may adsorb hydrogen, leading to hydrogen embrittlement and subsequent cracking that would not normally occur with a dry atmosphere (Ref 19, 26). One source of distortion is when a part is in contact with the furnace hearth during heating, which may produce sufficient nonuniform temperature distribution within the part. This will occur because the portion of the part in contact with the furnace hearth will be heated conductively much faster than the remainder of the

Table 7 Minimum recommended material removal from hot rolled steel products to prevent surface seam and nonmetallic stringer problems during heat treatment Minimum material removal per side(a) Condition

Turned on centers Centerless turned or ground

Nonresulfurized

3% of diameter 2.6%

Resulfurized

3.8% of diameter 3.4%

(a) Based on bars purchased to special straightness, i.e., 3.3 mm in 0.04 m (0.13 in. in 5 ft) maximum. Source: Ref 22

part surface, which is heated primarily by radiation. Thus, as the hotter surface tries to expand, it will be restrained by the cooler steel, leading to a hot upsetting condition and possibly significant distortion even if quenched uniformly (Ref 21). A similar condition exists if the tray of gears is placed near radiant tube heaters or electric heating elements in the furnace wall and the remainder of the gear surface is heated by radiation from the roof of the furnace. Localized overheating is particularly a potential problem for inductively heated parts (Ref 4, 26). Subsequent quenching of the part leads to quench cracks at sharp corners and areas with sudden changes in cross-sectional area (stress raisers). Cracking is due to increases of residual stresses at the stress raisers during the quenching process. The solution to the problem is to increase the heating speed by increasing the power density of the inductor. The temperature difference across the heated zone is decreased by continuous heating or scanning of several pistons together on a single bar (Ref 26). For heat treating problems related to furnace design and operation, it is usually suggested that (Ref 19):

The vestibules of atmosphere-hardening furnaces should be loaded and unloaded with purging. Load transfer for belt and shaker hearth furnaces should only occur with thorough purging to minimize atmosphere contamination. If the load being heated in the furnace is excessively large, either nonuniform heating over the specified heating cycle or nonuniform cooling when quenched will result. In such cases, either the production rate can be increased or smaller loads should be processed.

Retained Austenite. Dimensional changes may occur slowly or quickly and are due to the volume composition of the transformation products formed upon quenching. One of the most important, with respect to residual-stress variation, distortion, and cracking, is the formation and transformation of retained austenite. For example, the data in Table 8 illustrate the slow conversion of retained austenite to martensite, which was still occurring days after the original quenching process for the two steels shown (Ref 15, 16). This is particularly a problem when dimensional control and stability is one of the primary goals of heat treatment. Therefore,

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microstructural determination is an essential component of any distortion control process.

additional general comments regarding quenchant selection include (Ref 4, 20):

Quenching Process Other than component design, the quenching process itself is one of the most frequently encountered problems in heat treating. When designing a quenching process, it is important to consider quenchant selection, quench severity, and quench uniformity. Quenchant Selection and Severity. Quench severity is defined as the “ability of a quenching medium to extract heat from a hot steel workpiece; expressed in terms of the Grossmann number (H)” (Ref 27). A typical range of Grossmann H-values (numbers) for commonly used quench media is provided in Table 9. Figure 26 provides a correlation between the H-value and the ability to harden steel, as indicated by the Jominy distance (J-distance) (Ref 20). Although Table 9 is useful to obtain a relative measure of the quench severity offered by different quench media, it is difficult to apply in practice, because the actual flow rates for “moderate,” “good,” “strong,” and “violent” agitation are unknown. Alternatively, the measurement of actual cooling rates or heat fluxes provided by a specific quenching medium does provide a quantitative meaning to the quench severity provided. Some illustrative values are provided in Table 10 (Ref 28). Typically, the greater the quench severity, the greater the propensity of a given quenching medium to cause increased distortion or cracking. This usually is the result of increased thermal stress, not transformational stresses. Specific recommendations for quench media selection used with various steel alloys is provided by standards such as AMS 2759. Some

Most machined parts made from alloy steels are oil quenched to minimize distortion. Most small parts or finish-ground larger parts are free quenched. Larger gears, typically those over 20 cm (8 in.), are fixture (die) quenched to control distortion. Smaller gears and parts, such as bushings, are usually plug quenched on a splined plug typically constructed from carburized 8620 steel. Although a reduction of quench severity leads to reduced distortion, it may also be accompanied by undesirable microstructures, such as the formation of upper bainite (quenched pearlite) with carburized parts. Quench speed may be reduced by quenching in hot (150 to 205 C, or 300 to 400 F) oil. When hot oil quenching is used for carburized steels, lower bainite, which exhibits properties similar to martensite, is formed.

Table 9 Typical quenching conditions and Grossmann H-values Quenching medium

Poor (slow) oil quench—no agitation Good oil quench—moderate agitation Very good oil quench—good agitation Strong oil quench—violent agitation Poor water quench—no agitation Very good water quench—strong agitation Brine quench—no agitation Brine quench—violent agitation Ideal quench

Grossmann H-value

0.20 0.35 0.50 0.70 1.00 1.50 2.00 5.00 ...

Note: It is possible with high-pressure impingement to achieve H-values greater than 5.00.

Table 8 Dimensional variation in hardened high-carbon steel with time at ambient temperature

Steel type

Change in length after time, Tempering Hard% · 103 temperature, ness, °C HRC 7 days 30 days 90 days 365 days

1.1% C tool steel, 790 C quench

None 120 205 260

66 65 63 61.5

9.0 0.2 0.0 0.0

18.0 0.6 0.2 0.2

27.0 1.1 0.3 0.3

40.0 1.9 0.7 0.3

1% C/Cr, 840 C quench

None 120 205 260

64 65 62 60

1.0 0.3 0.0 0.0

4.2 0.5 0.1 0.1

8.2 0.7 0.1 0.1

11.0 0.6 0.1 0.1

Fig. 26

Quench severity in terms of Grossmann (H) values. J, Jominy distance. Source: Ref 20


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Excellent distortion is typically obtained with austempering, quenching into a medium just above the Ms temperature. The formation of retained austenite is a significant problem with austempering processes. Retained austenite is most pronounced where manganese and nickel are major components. The best steels for austempering are plain carbon and chromium and molybdenum alloy steels (Ref 20). Aqueous polymer quenchants may often be used to replace quench oils, but quench severity is still of primary importance. Gas or air quenching will provide the least distortion and may be used if the steel has sufficient hardenability to provide the desired properties. Low-hardenability steels are quenched in brine or vigorously agitated oil. However, even with a severe quench, undesirable microstructures, such as ferrite, pearlite, or bainite, can form.

Kern and Suess have provided guidelines for hardening steels to achieve optimal microstructural control (Ref 21). To minimize the potential for cracking:

In carbon or alloy steels containing50.3% C maximum, use a water quench. Steels with 0.3 to 0.38% C can be water quenched if they are in the form of simple shapes such as round bars. If the carbon content is 40.38%, an oil quench should be used. (More current references, such as AMS 2759, would permit polymer quenching of some alloys if appropriate quench bath maintenance procedures are used.) Exceptions are carbon steels with low alloy content (maximum of 1% Mn). Carbon steels containing40.95% C and 0.30 to 0.50% Mn

Table 10 Comparison of typical heat-transfer rates Quench medium

Still air Nitrogen (1 bar) Salt bath or fluidized bed Nitrogen (10 bar) Helium (10 bar) Helium (20 bar) Still oil Hydrogen (20 bar) Circulated oil Hydrogen (40 bar) Circulated water

Heat-transfer rate, W m

50–80 100–150 350–500 400–500 550–600 900–1000 1000–1500 1250–1350 1800–2200 2100–2300 3000–3500

2

K

1

can be water quenched if they are in the form of simple shapes and have no drilled or punched holes. Other guidelines of Kern and Suess for hardening steels include (Ref 21):

If the part has widely varying section sizes (ratio of 3 to 1), or if it has holes, keyways, or grooves, water quenching may produce cracking regardless of the carbon content. Designing with generous fillets in these regions may resolve the problem. If the distortion must be as low as possible, oil or salt quenching should be used with appropriate qualification. More recent work has shown that polymer quenching may be used in some cases. If 100% bainite is required, austempering in molten salt should be performed. To assure that no retained austenite remains, a final temper slightly below the austempering temperature is recommended.

Quenchant Uniformity. Quench nonuniformity is one of the greatest contributors to quench cracking. Quench nonuniformity can arise from nonuniform flow fields around the part surface during the quench or nonuniform wetting of the surface (Ref 20, 29–32). Both lead to nonuniform heat transfer during quenching. Nonuniform quenching creates large thermal gradients between the core and the surface of the part. When there is nonuniform cooling within the part between the Ms and Mf, there will be a stretching or elongation in areas where the cooling is slow, which will act as a push stress, leading to push cracking (Ref 33, 34). Another form of cracking is pull cracking, which occurs with nonuniform surface cooling between the austenitizing temperature and Ms. Push cracking and pull cracking are the opposite of each other, although the cracking event takes place for both between the Ms and Mf. Figure 27 provides illustrations of both push and pull cracking (Ref 32). Poor agitation design is a major source of quench nonuniformity. The purpose of the agitation system is not only to take hot fluid away from the surface and to the heat exchanger, but it also provides uniform heat removal over the entire cooling surface of all of the parts throughout the load being quenched. The batch quench system in Fig. 28 illustrates a system where axial (vertical) quenchant flow occurs throughout a load of round bars lying

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horizontally in a basket (Ref 19). In this case, the bottom surfaces of the bars experience greater agitation than the top surfaces. Cracks form on the upper surfaces because of the nonuniform heat loss. Agitation produces greater heat loss at the bottom, creating a large thermal gradient between the top and the bottom surfaces. If a submerged spray manifold is used to facilitate more uniform heat removal, the following design guidelines are recommended:

The total surface of the part should experience uniform quenchant impingement. The largest holes possible (2.3 mm or 0.09 in. minimum) should be used. The manifold face should be at least 13 mm (0.5 in.) from the surface of the parts being quenched. Repeated removal of hot quenchant and vapor should be possible.

Excessive distortion was also obtained with an agitation system illustrated in Fig. 29 when the quenchant flow was either in the same direction relative to the direction of part immersion or in

the opposite direction (Ref 31). The solution to this problem was to minimize the quenchant flow to that required for adequate heat transfer during the quench and to provide agitation by mechanically moving the part up and down in the quenchant. Identifying sources of nonuniform fluid flow during quenching continues to be an important tool for optimizing distortion control and minimizing quench cracking. Nonuniform thermal gradients during quenching are also related to interfacial wetting kinematics, which are of particular interest with vaporizable liquid quenchants, including water, oil, and aqueous polymer solutions (Ref 32). Most liquid vaporizable quenchants exhibit boiling temperatures between 100 and 300 C at atmospheric pressure. When parts are quenched in these fluids, surface wetting is usually timedependent, which influences the cooling process and the achievable hardness. Another major source of nonuniform quenching is foaming and contamination. Contaminants include sludge, carbon, and other insolubles. It includes water in oil, oil in water, and aqueous polymer quenchants. Foaming and contamination lead to soft spotting, increased distortion, and possibly cracking.

Stress Raisers and Their Role in Quench Cracking

Fig. 27

Two forms of quench cracking. Source: Ref 32

Fig. 28

Harmful effects of impeded vertical quenchant flow through the load of a batch quench system. Source:

Ref 19

Not all quench failures occur immediately following the quench; some failures that occur during subsequent use may be due to unacceptably high and/or nonuniform stresses that are imparted during the quenching process and may even be unpredictable. As already discussed, quench cracking occurs due to thermal contraction stresses coupled with

Fig. 29

Effect of quenchant flow direction on distortion. Source: Ref 31

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the volumetric expansion that accompanies the martensitic transformation. It is directly proportional to carbon content and microstructural factors. These cracks can be instantaneous upon quenching, or they may be delayed. Also, some components may be crack-free, whereas seemingly identical components may have cracked. Delayed quench cracks can be the result of additional transformation of retained austenite in steel. This occurs when heavily stressed retained austenite continues to transform to martensite prior to tempering or even after tempering, if there is sufficient retained austenite. However, as mentioned previously, surface damage and inadequate microstructure (decarburizing, banding, inclusions, coarse grain size), among others, may also cause the part to fail (Ref 33–35). Quench cracking typically initiates at the surface, particularly at positions where geometrical changes occur, such as at corners, defects, and inclusions. Quench cracks always begin at the part surface and have characteristics that are easily recognized. First, the fracture generally runs from the surface toward the center of the mass in a relatively straight line, with either a longitudinal or radial orientation unless located by a change in section size. The crack is also likely to open or spread and may exhibit a shear lip. Shear lips are ledges on the side of the specimen that make a 45 angle to the plane of fracture and may be present on the edges of some predominantly brittle fractures to form a “picture frame” around the surface (Ref 35). The fracture surfaces of quench cracks are always intergranular. It is common to find secondary cracking, which forms from and after the main crack, indicating that the component was under high stress. Because the quenching process involves high levels of thermal and transformation stresses, the presence of imperfections in the microstructure can increase the risk of cracking the part. Imperfections such as very small cracks, inclusions, some second-phase particles and defects from prior machining, and laps and seams work as stress raisers. At positions far removed from those defects, the stress is just nominal stress, that is, the load divided by the cross-sectional area. This does not pose a problem if the applied stress is below the elastic limit. However, in the vicinity of small defects or cracks, the situation changes, and the stress is amplified. Because of this, they are called stress raisers and are very important during quenching as well as during service.

When cracked parts are subsequently tempered, the intergranular morphology may form a thick oxide scale from the tempering process. The microstructure adjacent to the crack will not be decarburized unless a specimen with an undetected quench crack is rehardened. In quenched and tempered steels, proof of quench cracking is often obtained by opening the crack and looking (visually) for temper color typical for the temperature used. The following are some case studies showing sources of cracking that are often attributed to the quench but whose root cause is not the quench itself. The quench only exacerbates the problem. There is only one example (case study 1) where the cracking root cause was the quench severity.

Case Studies in Quench Cracking Case Study 1: As-Quenched 4340 Steel (Ref 25, 36). A component (Fig. 30) from AISI 4340 steel cracked during heat treatment. Chemical analysis of the component confirmed a composition compatible with AISI 4340 steel. As shown in Fig. 30, the crack passes straight from the surface to the core. Quench cracks always begin at the part surface and have characteristics that are easily recognized. First, the fracture generally runs from the surface toward

Fig. 30

Macrograph of AISI 4340 quenched and tempered steel illustrating macroetched pure quench crack. Source: Ref 25, 36

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the center of the mass in a relatively straight line, with either a longitudinal or radial orientation unless located by a change in section size (Ref 35). The component steel is considered a highhardenability steel because of alloying elements such as chromium, nickel, and molybdenum. It is recommended that when using high-hardenability steels, quenchants exhibiting lower quench severity and time should be used. This is important because cracking can occur during quenching due to thermal contraction stresses coupled with the volumetric expansion that accompanies the martensitic transformation. As already described, excessive cooling rates (high quench severity) will produce greater thermal stresses in addition to greater transformation stresses. If the total residual stresses in the part exceed the yield strength of the steel, distortion will occur. If the ultimate strength is exceeded, cracking will occur (Ref 4, 5).

With respect to the case being discussed here, the cause of cracking was identified as due to excessively high quench severity. Case Study 2: Cracking of 4140 Block Forging after Quenching and Tempering. Cracking was observed to occur with an AISI 4140 block forging subsequent to quenching and tempering. Chemical analysis was performed on the steel block and compared to the specification range for this alloy (Table 11), which confirmed that the steel was nominally 4140. To verify the presence or absence of inclusions (quantity, morphology, and distribution), a metallographic examination in the unetched condition was performed. In this condition, although the microstructure is not revealed, it is easier to identify inclusions. The steel was examined near the crack, and the results are shown in Fig. 31(a). No evidence of nonmetallic inclusions was found that could be attributed to the observed crack formation.

Table 11 Chemical analysis Chemical composition, wt% Material

AISI 4140 Block

Fig. 31

C

Mn

Si

P

S

Cr

Ni

Mo

0.38–0.43 0.39

0.75–1.00 0.88

0.15–0.35 0.15

0.035 0.013

0.040 0.028

0.80–1.10 0.86

0.25 0.06

0.15–0.25 0.16

Representative view of the surface and crack profile from the block sample. (a) Unetched condition. (b) Etched with 2% nital. Original magnification: 100 ·


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AISI 4140 AISI 4130 Component

C

Mn

Si

P

S

Cr

Ni

Mo

Cu

0.38–0.43 0.28–0.33 0.40

0.75–1.00 0.40–0.60 0.88

0.15–0.35 0.15–0.35 0.25

0.035 0.035 0.025

0.040 0.040 0.031

0.80–1.10 0.80–1.10 1.03

0.25 0.25 0.09

0.15–0.25 0.15–0.25 0.22

... 0.35 0.25

Fig. 32

Representative view of the crack propagating from porosity or voids within the brazed joint. Unetched. Original magnification: 100 ·

The surface was also examined, and no evidence was observed of detrimental surface conditions, such as small cracks or defects due to machining that could contribute to the cracking. This is important because quenching involves high levels of thermal and transformation stresses, and the presence of imperfections in the microstructure can increase the risk of cracking the part. Imperfections act like very small cracks. Inclusions and other surface defects function as stress raisers. In Fig. 31, it is possible to see evidence of secondary cracking. It is common to find secondary cracking that forms around the main crack, indicating that the component was under high stress due to thermal contraction stresses coupled with the volumetric expansion that accompanies the martensitic transformation. Metallographic examination in the etched condition is necessary to verify other microstructural characteristics. For steels, the most common etchant is 2% nital (2 mL HNO3 + 98 mL ethanol, 95%). Figure 31(b) shows the same region of the Fig. 31(a) but in the etched condition. Examination of the crack profiles revealed no evidence of decarburization in the crack, although tempering oxide was observed. The

fracture surfaces of quench cracks are always intergranular, since it is a brittle crack (Ref 37). When cracked parts are subsequently tempered, intergranular morphology is usually observed in the quench crack, which is due to a thick oxide scale from the tempering process. Its presence means that the fracture surface was present before the tempering process. After these analyses, it was possible to conclude that quenching stresses were the main cause of the failure, and it was recommended that a less severe quenching condition be used. Case Study 3: Use of Improper Steel Alloy and Presence of Voids in a Steel Brazed Joint (Ref 25). A reamer fabricated from an AISI 4130 steel shaft was brazed to an AISI 4130 steel body. After quenching and tempering, cracking was observed at the brazed joint, which propagated into the reamer body. The nominal range of chemical compositions for AISI 4130 and 4140 steel are provided in Table 12. Chemical analysis of the component, which was thought to be AISI 4130, is also shown, which confirms that the wrong steel alloy was used. The chemical analysis of the component is consistent with that for AISI 4140. The higher carbon content of AISI 4140 relative to AISI 4130 means greater hardenability and therefore greater probability for cracking and increased distortion to occur. In such cases, quenching should be less severe. From these data, it would appear that the AISI 4130 steel reamer body was most likely exposed to an excessively high cooling rate for this steel alloy if heat treatment parameters were set for 4130 steel. To verify the presence or absence of inclusions (quantity, morphology, and distribution), a metallographic examination in the unetched condition was made. In this case, the examination was also made near the crack, and the results are shown in Fig. 32. Figure 32 did not reveal evidence of nonmetallic inclusions, although it is possible to observe the presence of voids in the brazed joint. These are undesirable and should be avoided, since voids are stress raisers by amplifying the


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stress during quenching and are nucleation sites for crack formation in the brazed joints. This occurs during fast cooling, leading to increased thermal stresses in the component. Figure 32 shows that cracking initiated from porosity or voids within the brazed joint and appears to have propagated into the reamer body from quenching stresses. Secondary cracks are also observed. Examination of the brazed joint in the etched condition, using 2% nital (2 mL HNO3 +98 mL ethanol, 95%), revealed the predominance of tempered martensite and a uniform microstructure (Fig. 33). From these analyses, it can be concluded that the presence of voids (stress raisers) within the brazed joint has nucleated cracks that propagated under quenching stresses. The incorrect steel grade increased the potential for cracking.

Fig. 33

Representative view of the brazed joint between the reamer body (bottom) and reamer shaft (top). Microstructure is tempered martensite. Etched with 2% nital. Original magnification: 100 ·

Case Study 4: Presence of a Seam Defect (Ref 25, 36). An AISI 4140 steel bar cracked after austenitization and quenching in a 20 to 21% aqueous polymer quenchant solution. Chemical analysis confirmed that the steel used was compatible with AISI 4140. Metallographic examination in the unetched condition was performed on the steel near the crack, which is shown in Fig. 34. No evidence of nonmetallic inclusions was observed. However, examination of the defect profile revealed the presence of seam defects. A seam defect is an unbounded fold or lap on the surface of the metal that appears as a crack and is usually the result of a seam that was formed but not closed during the working process (such as rolling, forging, etc.) of the material. These defects typically exhibit the presence of scale and high-temperature oxidation adjacent to the crack, as shown in Fig. 34. Examination of the etched condition (2% nital), shown in Fig. 35, reveled a uniform cross-sectional martensitic microstructure. Examination of the crack profile revealed a seam defect. Cracking appears to have initiated from this defect and propagated from quenching stresses. Evidence of decarburization (lighter regions) and high-temperature oxidation can also be observed within the defect profile. Thus, cracking was caused by seam defects that nucleated crack formation, which then propagated due to quenching stresses. Case Study 5: Presence of Slag Inclusions and a Lap Defect (Ref 25). Longitudinal cracks after quenching and tempering were obtained with an unthreaded AISI 4140 stud bolt (25.4 mm diameter).

Fig. 35

Fig. 34

View of the identified seam defects in a bar sample of AISI 4140. Unetched. Original magnification: 100 ·

Aspect of the defect in the etched condition. In the cross section, it is possible to see a uniform microstructure compounded by martensite. Decarburizing and hightemperature oxidation can be observed. Etched with 2% nital. Original magnification: 100 ·

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AISI 4140 Sample

C

Mn

Si

P

S

Cr

Ni

Mo

Cu

0.38–0.43 0.39

0.75–1.00 0.81

0.15–0.35 0.22

0.035 0.014

0.040 0.027

0.80–1.10 0.92

0.25 0.15

0.15–0.25 0.18

0.35 0.26

Fig. 36

(a) Representative view of large slag-type inclusions observed throughout the sample cross sections. Unetched. Original magnification: 100 ·. (b) View of the crack profile and slag-type inclusions observed adjacent to the cracking. Unetched. Original magnification: 100 ·

Fig. 37

Representative view of the crack surface profile. Unetched. Original magnification: 100·

Chemical analysis of this steel is shown in Table 13, together with the nominal composition range of AISI 4140 steel. These data confirm that the steel is consistent with AISI 4140. Metallography results for the steel in the unetched condition are shown in Fig. 36(a

and b). Large slag-type inclusions were observed throughout the cross section, as shown in Fig. 36(a). Figure 36(b) shows these slag-type inclusions adjacent to the crack. The steel test specimen used for Fig. 37 is also unetched and shows a surface profile of the crack. A surface seam or lap-type defect is evident, and the crack appears to propagate through or from a surface seam. A lap is a surface defect that appears as a seam and is caused by folding over of hot metal, fins, or sharp corners and then rolling or forging them into the surface, although they are not welded close by the hot surfaces involved (Ref 38). Secondary cracks are also observed. Those observations are important, since quenching involves relatively high levels of thermal and transformation stresses, and the presence of imperfections in the microstructure can increase the risk of cracking. Imperfections such as inclusions and other surface defects act as stress raisers. Figure 38 shows the steel in the etched condition (2% nital), which reveals a uniform tempered martensite microstructure with a slag inclusion stress raiser. From those observations, it can be concluded that cracking was caused by quenching stresses acting upon stress-concentration sites of large


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cracking appeared to be intergranular, showing tempering oxide within the profile. No evidence of decarburization was observed. As noted before, when cracked parts are subsequently tempered, the intergranular morphology may form a thick oxide scale from the tempering process. Quenching stresses associated with the nonuniform microstructure, caused, in part, from a slack quench condition or inherent chemical segregation, have contributed to the observed cracking. Slack quenching is related to hardening of steel from the austenitizing temperature at a rate slower than the critical cooling rate, resulting in incomplete transformation and the formation of one or more transformation products in addition to or instead of martensite. Case Study 7: Network Carbides and Coarse Grain Size. A low-alloy 17CrNiMo6 (0.18% C, 0.25% Si, 0.50% Mn, 1.65% Cr, 0.80% Mo, 1.55% Ni) carburized steel gear produced cracks after carburizing, quenching, and tempering. The carburized case was 1.8 to 2.0 mm, and the measured surface hardness was 57 to 61 HRC. After carburizing, the gear was quenched from 840 C using an aqueous polymer quenchant and subsequently double tempered for 5 h at 240 C and 3 h at 260 C. Figure 41(a) illustrates the sectioned gear. Examination of the cracking zone in the etched condition (2% nital) showed that intergranular cracking occurred in the boundary of the coarse-grained structure, as illustrated in Fig. 41(b). Quench cracking typically initiates at the surface, particularly at positions where geometrical changes occur, such as at corners, defects, and inclusions. It always begins at the part surface and has characteristics that are easily recognized. The fracture surfaces of quench cracks almost always occur intergranularly (Ref 37). Quench cracking is considered a complex mechanism of intergranular fracture and can be aggravated by the various mechanisms of grain-boundary weakening (such as segregation of embrittlement elements to the grain boundary) and grain size. However, it also is heavily influenced by volumetric expansion during transformation hardening and the

slag inclusions adjacent to the crack and a small seam or lap defect at the surface. Cracks were also observed propagating from and/or through an apparent surface seam or lap-type defect. Case Study 6: Presence of Chemical Segregation (Ref 25, 36). A press-formed steel flange made from AISI 1035 steel produced cracks after quenching in an aqueous polymer quenchant solution. Chemical composition for this steel grade is provided in Table 14. Comparison with the actual composition of the steel from the flange confirmed that it was the correct AISI 1035 steel grade. Cracking was observed at the press-formed ring location. Examinations (unetched) of the press-formed ring location were performed. Figure 39, shows that in the unetched condition, there is no evidence of surface imperfections or inclusions that could be attributed to the cracking of the steel. However, examinations in the etched condition revealed a microstructure that exhibited chemical segregation in the form of banding, as shown in Fig. 40(a). The microstructure is nonuniform and consists of bainite and tempered martensite, as illustrated in Fig. 40(b). Examination of the outer radius of the pressformed ring revealed evidence of cracking. The

Fig. 38

Etched condition showing tempered martensite microstructure and slag inclusion. Etched with 2% nital. Original magnification: 100·


AISI 1035 Flange

C

Mn

Si

P

S

Cr

Ni

Mo

Cu

0.32–0.38 0.33

0.60–0.90 0.73

... 0.18

0.04 0.012

0.050 0.003

... 0.04

... 0.02

... 50.01

... 50.01


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temperature extremes of quenching. Causes of intergranular brittle fracture include brittle second-phase particles and/or films in grain boundaries (Ref 37).

Fig. 39

Representative view of the cracking associated with the radius of the press-formed ring. Unetched. Original magnification: 100 ·

The microstructure of the carburized case was carefully examined, which also showed the presence of brittle network carbides in the prioraustenitic grain boundaries. The formation of network carbides is an indication that the carbon potential employed was too high for the steel concerned (Ref 39). If, during carburizing, the austenite is supersaturated with carbon, that is, it contains carbon in excess of the eutectoid composition (0.8% C), the carbide will precipitate at the grain boundaries during slow cooling from the carburizing temperature. Under equilibrium cooling conditions, an austenitized steel, having a carbon content above the eutectoid carbon content, will reject the excess carbon as carbide (Fe3C). However, if the same austenite were to be cooled quickly, most of the excess carbon would be retained by the resultant martensite-austenite structure (Ref 39).

Fig. 40

(a) Representative view of the chemical segregation (banding). Etched with 2% nital. Original magnification: 50 ·. (b) Higher magnification of the microstructure showing tempered martensite and bainite. Etched with 2% nital. Original magnification: 400 ·

Fig. 41

(a) Carburized steel gear (17CrNiMo6). (b) Representative view of the cracking zone. Presence of coarse grains and intergranular cracking. Etched with 2% nital

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Carbides, being ceramic compounds, are brittle. The more accelerated the diffusion of carbon, the coarser the carbides become. Oversized carbides do not resist thermal fluctuations, because they are ceramic compounds; therefore, cracking initiates due to thermal shock or fast ramp-up and ramp-down of a furnace. Because carbides in the network configuration possess a brittle nature, once a carbide segment of a network starts cracking, crack propagation is very fast (Ref 40). During carburizing, it is necessary to carefully control the carbon content of the surface case. If this layer becomes hypereutectoid, cementite will be present in the boundaries of the grain, forming a network, as demonstrated in Fig. 42(a and b). In addition to network carbides, nonuniform grain size (ASTM 1 and 2 ) is also observed in Fig. 42(a and b). Nonuniform grain size is a problem, since the hardenability of a carbon steel may increase as much as 50% with an increase in austenite grain size from ASTM 8 (6 to 10) to ASTM 3 (1 to 4). This phenomenon causes a nonuniform martensitic transformation, contributing to increased stress during quenching. The effect becomes more pronounced if the carbon content is increased at the same time (Ref 39). These factors contributed to the cracking observed during quenching. Excessive carbon content (more than 0.8%) in the carburized layer was related to the high carburizing temperature that was used (950 C). Higher temperatures result in greater solubility of the carbon in the austenite phase. Grain growth can also be related to high process temperature and steel chemistry.

Fig. 42

Alloying elements such as aluminum, niobium, vanadium, or titanium function as graingrowth inhibitors, and the steel grade used in this case did not contain these elements (Ref 38, 39). Therefore, it was concluded that the presence of network carbides and nonuniform grain size, coupled with quenching stresses, was responsible for the observed cracking. Case Study 8: Presence of Stringer Inclusions and Chemical Segregation (Ref 25). Pins of AISI 1144 steel (resulfurized steel grade) were through hardened prior to induction hardening of the pin tip, and cracking and soft spots were obtained. The nominal chemical composition of AISI 1144 steel is provided in Table 15 along with the chemical analysis of the component. These results confirm that the steel used for the component is consistent with AISI 1144. Examinations in the unetched condition revealed evidence of many long stringer inclusions, which were oriented in streaks or bands that were parallel (or longitudinal) with respect to the length of the pins. A stringer inclusion occurs when an impurity, either metallic or nonmetallic, is trapped in the ingot and is elongated in the direction of hot working. It appears as a narrow streak that is parallel to the direction of hot working (Ref 38). Stringers were also observed extending to the pin tip surface within the induction-hardened case (Fig. 43). Stringers can, like other inclusions, act as stress raisers. At positions that may be far removed from those defects, the applied stress (applied load/crosssectional area) may be normal and will not pose any problem. However, if the applied stress is

(a) Microstructure of the tooth top showing boundary carbides and coarse grains. (b) Detail of the brittle carbide network showing prior-austenitic grain size and tempered martensite

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AISI 1144 Pin

Fig. 43

C

Mn

Si

P

S

Cr

Ni

Mo

Cu

0.40–0.48 0.44

1.35–1.65 1.50

... 0.23

0.04 0.008

0.24–0.33 0.29

... 0.05

... 0.02

... 0.02

... 50.01

Aspect of the stringer inclusions observed within the pin tip location. Unetched. Original magnifica-

tion: 100·

less than the elastic limit in the vicinity of small defects or cracks, the stress is amplified. Therefore, defects such as stringer inclusions are called stress raisers and are very important during quenching as well as during use. Examination in the etched condition showed that the microstructures consisted of many areas of chemical segregation. Unavoidable chemical segregation of alloying elements occurs during the solidification of an ingot in the steel production process. If this occurs on a grain-sized scale, it is called microsegregation. If chemical segregation occurs on a much larger ingot-sized scale, it is referred to as macrosegregation, and the inhomogeneous steel structure will possess nonuniform properties throughout, particularly in the direction that is transverse to the hot working direction (Ref 38). In the AISI 1144 steel sample being analyzed, the chemical segregation appeared as ferrite bands and was associated with the stringer inclusions, which can be associated with the soft spots that also can contribute to cracking (Fig. 44a and b). It should be noted that stringer inclusions are not uncommon for resulfurized steel grades, because the distribution and shape of these inclusions are often difficult variables to control. The fracture pattern of the pin tips subsequent to induction hardening appears to have propagated from the stringer inclusions and

associated chemical segregation in the presence of quenching stresses. Case Study 9: Decarburization and Oxidized Grain Boundary (Ref 25, 36). A heavy wall tube section of AISI 4140 tube stock produced cracks after quenching and tempering. The component was austenitized at 843 C for 2 h and then quenched into an aqueous polymer solution (25%) and tempered at 565 C for 2 h, then air cooled. Chemical analysis is shown in Table 16 together with the nominal composition of AISI 4140. These data confirm that the proper steel alloy was used. Although it is not shown here, the steel in the unetched condition revealed no evidence of a large number of nonmetallic inclusions. Examination in the etched condition of the crosssectional microstructure shows that it consists of uniform tempered martensite (Fig. 45). However, examination of the surface in the etched condition revealed high-temperature grain-boundary oxidation (Fig. 46). Hightemperature grain-boundary oxidation occurs when grain boundaries starting at the surface of the part are oxidized. Normally, when parts are being heat treated, such as during carburization or austenitization (hardening), if the furnace contains free oxygen from air leakage (ingression) into the furnace or excessive vapor or steam is present, oxygen will diffuse into the surface of the material, resulting in oxidation of the grain boundaries and degradation of the engineering properties at the surface (Ref 40). Hightemperature grain-boundary oxidation also acts as a stress-concentration site for crack initiation. The surface profile in the etched condition (Fig. 47) revealed evidence of partial decarburization and tempering oxide within the cracks. When cracked parts are subsequently tempered, the intergranular morphology may form a thick oxide scale from the tempering process. Decarburization appears when steels are processed by forming, forging, heat treating, or any other thermal treatments where the material temperature may exceed 760 C for some time with no atmospheric protection. If this occurs, the steel may start losing carbon from the heated surfaces, leading to a decarburized surface. This


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Fig. 44

View of the induction-hardened pin tip location. Ferrite bands and inclusions can be observed. Etched with 3% nital. (a) Original magnification: 100 ·. (b) Original magnification: 200 ·


AISI 4140 Sample

Fig. 45

C

Mn

Si

P

S

Cr

Ni

Mo

Cu

0.38–0.43 0.44

0.75–1.00 0.97

0.15–0.35 0.24

0.035 0.012

0.040 0.016

0.80–1.10 1.06

0.25 0.12

0.15–0.25 0.19

0.35 0.18

Cross-sectional microstructure showing uniform tempered martensite. Etched with 2% nital. Original magnification: 400 ·

Fig. 46

Surface profile adjacent to the cracking. Evidence of high-temperature grain-boundary oxidation. Etched with 2% nital. Original magnification: 400·


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Fig. 47

View of the crack profile showing tempering oxide within the crack and decarburization at the surface. Etched with 2% nital. Original magnification: 50 ·

decarburized layer can be partial or full and can degrade the engineering properties of the surface relative to the matrix of the material. Based on this evidence, the identified cracking was attributed to quenching stresses acting on the oxidized surface grain boundaries. REFERENCES

1. Y. Toshioka, Heat Treatment Deformation of Steel Products, Mater. Sci. Technol., Vol 1, 1985, p 883–892 2. S. Mocarski, Carburizing and Its Control— I. Basic Considerations, Ind. Heat., Vol 41 (No. 5), 1974, p 58–70 3. A. Bavaro, Heat Treatments and Deformation, Trait. Therm., Vol 240, 1990, p 37–41 4. F. Legat, Why Does Steel Crack During Quenching, Kovine Zlitine Technol., Vol 32 (No. 3–4), 1998, p 273–276 5. R.W. Bohl, “Difficulties and Imperfections Associated with Heat Treated Steel,” MEI Course 10, Lesson 13, ASM International 6. K.E. Thelning, Steel and Its Heat Treatment, Butterworths, London, U.K., 1985 7. H. Berns, Distortion and Crack Formation by Heat Treatment of Tools, Radex Rundsch., Vol 1, 1989, p 40–57

8. T. Kunitake and S. Susigawa, Sumitomo Search, May 1971, p 16–25 9. R.R. Blackwood and L.M. Jarvis, Ind. Heat., March 1991, p 28–31 10. H. Fujio, T. Aida, and Y. Masumoto, Bull. Jpn. Soc. Mech. Eng., Vol 20, 1977, p 1655– 1662 11. Yu.A. Geller and V.P. Brimene, Steel USSR, July 1971 12. G.E. Totten and M.A.H. Howes, Chapter 5—Distortion of Heat Treated Components, Steel Heat Treatment Handbook, G.E. Totten and M.A.H. Howes, Ed., Marcel Dekker Inc., New York, NY, 1997, p 292 13. C.E. Bates, G.E. Totten, and R.L. Brennan, Quenching of Steel, Heat Treating, Vol 4, ASM Handbook, ASM International, 1991, p 67–120 14. D.A. Canonico, Stress-Relief in Heat Treating of Steel, Heat Treating, Vol 4, ASM Handbook, ASM International, 1991, p 33–34 15. W.T. Cook, A Review of Selected SteelRelated Factors Controlling Distortion in Heat-Treatable Steels, Heat Treat. Met., Vol 26 (No. 2), p 27–36 16. Y. Toshioka, Heat Treatment Deformation of Steel Products, Mater. Sci. Technol., Vol 1, Oct 1985, p 1883–1892 17. A. Thuvander and A. Melander, Calculation of Distortion During Quenching of a Low Carbon Steel, First ASM Heat Treatment and Surface Engineering Conference in Europe (Proc.), Part 2, Trans. Tech. Publications, 1992, p 767–782 18. P.C. Clarke, Close-Tolerance Heat Treatment of Gears, Heat Treat. Met., Vol 25 (No. 3), 1998, p 61–64 19. R.F. Kern, Thinking Through to Successful Heat Treatment, Met. Eng. Q., Vol 11 (No. 1), 1971, p 1–4 20. R. Kern, Distortion and Cracking, II: Distortion from Quenching, Heat Treat., March 1985, p 41–45 21. R. F. Kern and M. E. Suess, Steel Selection: A Guide for Improving Performance and Profits, Wiley-Interscience, New York, 1970, p 15–50 22. R. Kern, Distortion and Cracking, III: How to Control Cracking, Heat Treat., April 1985, p 38–42 23. H.-J. Chen and Z.W. Jiang, Microstructure Improvement and Low-Temperature Quenching of Dies Made of Cr12-Type Steel, Jinshu Rechuli, No. 8, 1992, p 39–41

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24. H.-H. Shao, Analysis of the Causes of Cracking of a 12% Cr Steel Cold Die During Heat Treatment, Jinshu Rechuli, No. 11, 1995, p 43 25. R.R. Blackwood, L.M. Jarvis, D.G. Hoffman, and G.E. Totten, Conditions Leading to Quench Cracking Other Than Severity of Quench, Heat Treating Including the Liu Dai Memorial Symposium, Proc. of the 18th Conf., R.A. Wallis and H.W. Walton, Ed., ASM International, 1998, p 575–585 26. X. Cheng and S. He, Analysis of Quenching Cracks in Machine-Tool Pistons Under Supersonic Frequency Induction Hardening, Heat Treat. Met. (China), No. 4, 1991, p 51–52 27. J.R. Davis, ASM Materials Engineering Dictionary, ASM International, 1992, p 407 28. P.F. Stratton, N. Saxena, and R. Jain, Requirements for Gas Quenching Systems, Heat Treat. Met., Vol 24 (No. 3), 1997, p 60–63 29. H.M. Tensi, G.E. Totten, and G.M. Webster, Proposal to Monitor Agitation of Production Quench Tanks, Heat Treating: Including the 1997 International Induction Heat Treating Symposium—Proc. of the 17th Conf., D.L. Milam, D.A. Poteet, G.D. Pfaffmann, V. Rudnev, A. Muehlbauer, and W.B. Albert, Ed., ASM International, 1997, p 423–441 30. S. Owaku, Quench Distortion of Steel Parts, Netsu Shori (J. Jpn. Soc. Heat Treat.), Vol 32 (No. 4), 1992, p 198–202 31. R.T. Von Bergen, The Effects of Quenchant Media Selection on the Distortion of Engineered Steel Parts, Quenching and

Distortion Control, G.E. Totten, Ed., ASM International, 1992, p 275–282 32. H.M. Tensi, A. Stich, and G.E. Totten, Fundamentals of Quenching, Met. Heat Treat., Mar/April 1995, p 20–28 33. V.D. Kalner and S.A. Yurasov, Internal Oxidation During Carburizing, Met. Sci. Heat Treat. (USSR), Vol 12 (No. 6), June 1970, p 451–454 34. A.A. Polyakov, Quenching Properties of Parts Having Stress Concentrators, Met. Sci. Heat Treat., Vol 37 (No. 7–8), 1995, p 324–325 35. W. Becker, Fatigue Failure, Failure Analysis and Prevention., Vol 11 ASM Handbook, ASM International, 2002, p 700–727 36. L.C.F. Canale, G.E. Totten, R.R. Blackwood, L.M. Jarvis, and D.G. Hoffman, “An Overview of Non-Quench Related Problems Often Attributed to the Quenchant and Quenching Process,” 59th Congresso Anual da ABM-Internacional, July 19–22, 2004 (Saõ Paulo, SP, Brazil) 37. S. Lampman, Intergranular Fracture, Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 2002, p 641–649 38. “Manufacturing Techniques,” http://www. mslab.boun.edu.tr/Heat_treatment.doc (Accessed April 24, 2007) 39 G. Parrish, The Influence of Microstructure and the Properties of Case-Carburized Components, ASM International, 1980, p 236 40. A. Abada, Why Do Heat-Resistant Alloys Fail?, Ind. Heat., Oct 2002, p 55–59

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Steel Failures due to Tempering and Isothermal Heat Treatment Jan Vatavuk, Universidade Presbiteriana Mackenzie L.C.F. Canale, Universidade de Saõ Paulo USP

FAILURE of swords made by early metalsmiths was a complex phenomenon for bladesmiths. The repeated working, heating, and cooling could cause embrittlement, with sword failure occurring in the most critical moments of a battle. Likewise, some of the earliest cannons would break apart after the first shots following manufacture. These problems occurred in the ferrous alloy application until the benefit of tempering became recognized. In the middle of the 18th century, the tempering process (and/or stress relief) received attention as a fundamentally important process in the heat treatment of the ferrous components of tools. Some ironsmith tools were treated by the so-called process of water annealing, whereby steel was tempered in the range of 300 to 600 C. The slow cooling was substituted by water cooling. At the beginning of the 20th century, Krupp developed a great number of patents based on water cooling after tempering of chromiumnickel steel. This phenomenon received attention after the start of WWI, when large amounts of steel were used by the armament industry. In 1917, the term tempering embrittlement was introduced by Dickenson, having been published in papers by Brarley, Hatfield, Philpot, and Grenet. Some investigators, such as Greves and his collaborators, began a set of experimental methods using notched bars to determine the susceptibility of tempering embrittlement. A relationship between the energy absorbed after water cooling and annealing was termed the steel susceptibility ratio. At that time, all the experiments were performed at room temperature, because no one anticipated that temperature may also have an effect on the results. The effect of test temperature received attention in the beginning of 1944, when Jolivet and Vidal introduced experiments at different

temperatures, resulting in a revision of all former data. A very important technological mark was the development of the beneficial molybdenum effect on the embrittlement reduction phenomenon, through work by Greaves and Jones (Ref 1). For some time, embrittlement due to the tempering process has been shown to be an important failure related to heat treatment. In this chapter, the causes and cases associated with problems originated by tempering are reviewed. However, to provide background on this phenomenon, a brief description of the martensite reactions and the steel heat treatment of tempering is given to review the different stages of microstructural transformation.

Martensite Before describing the solid-state reactions resulting from the tempering process in the ferrous matrix, it is important to define the martensitic structure as a function of the alloying elements, especially for the carbon effect. Ferrous martensite is composed of a body-centered tetragonal crystallographic structure, with lattice parameters (c and a) related to the carbon contents of its chemical composition, as shown in the expression (Ref 2): c=a=1+0:0467 · (wt% C)

The lattice ratio for the tetragonal structure is approximately 1.0467, with 1 wt% C in solid solution. As shown in Fig. 1, hardness varies with carbon content, and that effect is strongly related to the distortions caused by the carbon atom in the body-centered tetragonal structure. Martensite is extremely hard (maximum of 800 to 900 HV) and brittle.

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The transformation of martensite from austenite is a nonequilibrium (athermal) process that occurs during rapid cooling from the austenite phase. Unlike the phase transformation from atomic diffusion at equilibrium temperatures, the martensitic transformation occurs when many atoms together undergo a shear displacement. This rapid shear displacement of atoms results in a rapid change in crystal structure during cooling. Thus, the martensitic transformation is referred to as a diffusionless process, because the transformation involves a nonequilibrium (athermal) crystal change from a shear transformation during rapid cooling from austenite.

Martensite is a nonequilibrium structure and thus does not appear on the iron-carbon equilibrium phase diagram. The face-centered cubic lattice shearing of austenite (Bain’s deformation) in a martensitic transformation is illustrated in Fig. 2. The deformation is large and rapid over many atoms in the lattice, and the change in the polycrystalline system is accommodated by lattice deformation. This can occur by slide, mechanical twinning, or even a mixture of both mechanisms according to the steel chemical composition. The crystal change results in an expansion of the polycrystalline system. The effect of temperature on martensite formation is directly related to the transformation temperatures of martensite start (Ms) and martensite finish (Mf). Carbon is the alloy element that has a higher influence on Ms temperature, which is mainly responsible for the martensite morphology of steels. There are several empirical formulas to calculate Ms temperature. Some examples are reported as follows (Ref 2–4): Ms =539 432 (%C) 30:4 (%Mn) 17:7 (%Ni) 12:1 (%Cr) 7:5 (%Mo)

For medium-carbon alloy steels (Ref 4): Ms =520 320 (%C) 50 (%Mn) 30 (%Cr) 20 ½%(Ni+Mo) 5 ½%(Cu+Si)

In Ref 3: Ms =561 474 (%C) 33 (%Mn) 17 (%Ni) 17 (%Cr) 21 (%Mo) (Ref 3)

The higher the transformation temperature, the higher the probability of the plastic deformation mechanism occuring by dislocation slide, although the low temperatures provoke a

Fig. 1

Hardness of martensite as a function of carbon content. Source: Ref 3

Fig. 2

(a) Body-centered tetragonal cell in austenite. (b) Body-centered tetragonal cell before (left) and after (right) the lattice deformation from austenite to martensite. Source: Ref 2, 3

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Steel Failures due to Tempering and Isothermal Heat Treatment / 287

plastic deformation by mechanical twinning. This way, it is possible to establish a martensite morphology with respect to the alloy content, giving special attention to carbon. Figures 3 to 5 show the morphological aspects of martensite as a function of carbon content for steels through optic microscopy. In the case of lath martensite, the deformation mechanism of the lattice is dislocation slip. This kind of martensite is also known as slipped martensite. Morphologically, this martensite presents lath packages, which are separated by low-angle boundaries (Ref 5). In the past, it was thought that the transformation units happened as lath packages, although recently it became clear that each lath is independently formed, and the evidence shows that an austenite film exists, which can be seen in Fig. 6.

Measurements of dislocation density found in martensite are on the order of 0.3 to 0.9 · 1013 cm/cm3 of the crystal. This dislocation density is higher than the maximum that can be obtained by elevating the percentage of cold plastic deformation. Some observations, made by transmission electron microscope, show a very small cellular structure (approximately 0.2 to 0.3 mm) inside the lath (Ref 5). Figure 4 shows the martensite with a high carbon content, observed with an optical microscope. With a high carbon content, the microstructure has twinned martensite or plate martensite. With higher carbon contents, higher volume of retained austenite (Fig. 5) occurs, because a higher carbon content lowers the Ms

Fig. 3

Fig. 5

Lath (low-carbon) martensite in SAE 8620 alloy steel (Fe, 0.2% C, 0.8% Mn, 0.55% Ni, 0.5% Cr, 0.2% Mo) after heat treatment (954 C, or 1750 F, for 1 h, water quench)

Microstructure of quenched 1.3% C steel. Dark needles of plate martensite and white areas of retained austenite (white arrow)

Fig. 6 Fig. 4

High-carbon martensite (0.8% C). Etched with nital

Ref 2

Illustration of the austenite film surrounding martensite laths in a Fe-10Cr-0.2C steel. Source: Adapted from

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temperature. Twinning density can be seen by transmission electron microscopy, because density is high even though twins are very nar˚ ). row (on the order of 10 A The percentage of slip martensite and twinned martensite in carbon steel and tool steel can be experimentally determined, as shown in Fig. 7. As the carbon content increases, the amount of lath martensite decreases. The untransformed austenite increases by the Ms temperature (martensite start temperature) reduction. The increase in the retained austenite volume fraction can reduce the as-quenched hardness mainly in the higher content range. Martensitic transformation causes an increase in volume and size variations, which contributes to the residual tension stresses that develop in the surface after the heat treatment of quenching, when transformation takes place in all of the sample cross sections, and transformation between surfaces and nucleus occurs anachronically. The volume variation measured during the transformation from austenite to martensite in a 1% C steel is approximately 4% (the transformation to pearlite results in a 2.4% expansion) (Ref 2), decreasing as far as the carbon is added in the matrix. This occurs because of the different carbon effect in the austenite related to the

Fig. 7

martensite. In the first, the deformation has a volumetric character, while in the second, it is more directional (Ref 7). This behavior can be seen in Fig. 8. It can be seen in Fig. 8 that the difference in the specific volume between austenite and martensite is approximately 15% from very low carbon content to very high carbon content (2% C). It is also interesting to observe that for low carbon, the change of volume from an asannealed condition to an as-hardened condition is practically nil. On the other hand, increasing carbon content raised that difference. These observations are important during the component process design. As mentioned earlier, the greater the carbon content, the greater the embrittlement of the martensite plates (Ref 8). However, since retained austenite also increases with carbon content, this fraction of retained austenite will bring some toughness to the matrix as well as reduce the volume variation as shown in the curve “A-FC” in Fig. 8. This fact results in a lower load to the cold surface because of the incomplete transformation of the nucleus. It is important to remember that the nucleus presents a low yield strength when retained austenite has a low mechanical resistance, decreasing the possibility of

Effect of carbon content on the lath martensite volume, retained austenite volume fraction, and Ms temperature. Source: Adapted from Ref 6

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developing tension stress in the surface of the component. It can be seen in Fig. 8 that for 2% C, theoretically it could be possible to obtain a complete transformation to martensite, but, this is not the case. However, it is interesting to observe that the retained austenite line matches the austenite line for very high carbon content. This is the case for Hadfield steels (Ref 9), which have a high manganese content (approximately 12%) that guarantees an austenitic microstructure, even though the carbon amount is approximately 1.2%. In this situation, it is possible to quench large components with complex geometry without the risk of developing cracks, even while increasing quenchant severity. The ability to form martensite is described in terms of hardenability, which is related to the presence of other alloy elements besides carbon. For example, molybdenum and manganese increase hardenability, while cobalt lowers the hardenability of steel. A higher hardenability allows martensite formation with a slower cooling rate. This is beneficial for reducing the tensile residual stresses in the component surface.

properties can be changed when the component is held isothermally at a temperature where austenite cannot form. It is important to emphasize that tempered martensite usually does not contain martensite. Instead, it is a structure of fine carbide particles in ferrite, which has formed from martensite during the tempering. This structure has a lower hardness than the martensite, but by proper choice of temperature and time used, the structure developed will be fine to give the desired hardness. Table 1 lists the colors associated with the tempering heats, and Table 2 illustrates the times required to reach furnace temperature during tempering (Ref 10). Effect on Mechanical Properties As noted, martensitic structures are too brittle for most practical applications. However, it is possible to enhance the structure toughness through tempering. The toughness usually comes at the expense of a decrease in yield

Table 1 Colors of tempering heats Temperature(a)

Tempering Tempering is historically associated with the heat treatment of martensite in steels. The resultant microstructure is called tempered martensite. The main purpose for tempering is to develop a usable combination of hardness and toughness. The microstructure and mechanical

Fig. 8

Temperature(b)

°C

°F

Color of oxides

°C

°F

188 199 210 221 232 254 265

370 390 410 430 450 490 510

Faint yellow Light straw Dark straw Brown Purple Dark blue Light blue

238 265 293 321 337 349 376

460 510 560 610 640 660 710

(a) Temperature held for 1 h. (b) Temperature held for 8 min

Specific volume (DV/V) of carbon steels relative to room temperature. Source: Adapted from Ref 7


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strength and hardness, as illustrated in Fig. 9 and 10. Figure 11 shows other modifications of mechanical properties that occur when an oilquenched AISI 4340 steel is tempered at temperatures above 200 C. The interrelationship between tempering temperature, steel chemistry, and hardness can be estimated by the equation (Ref 13): HB=2:84Hh +75(%C) 0:78(%Si) +14:24(%Mn)+14:77(%Cr)+128:22(%Mo) 54:0(%V) 0:55T t +435:66

where HB is the Brinell hardness after hardening and tempering, Hh is the Rockwel (HRC) hardness after hardening, and Tt is the tempering temperature in C. This equation was developed for the following conditions:

pffiffiffiffi Tt =647 [S(60 C+20)=Ht 70:9]1=4 73:45 SHt +(5377561S)(%C)+505S(%V)+219S(%Mo) +75S(%Cr)+66S(%Si)751 [C ]

where Ht is the hardness after hardening and tempering (HRC), S is the degree of hardening, Sj1.0, and the alloying elements are given in weight percent. This formula is valid for a tempering time of 2 h. Tempering Reactions Tempering is a process in which the microstructure approaches equilibrium under the influence of thermal activation. It follows that the tendency to temper depends on how far the starting microstructure deviates from

Hh = 20 to 65 HRC and Tt = 500 to 600 C C = 0.20 to 0.54%, Si = 0.17 to 1.40%, Mn = 0.50 to 1.90%, and Cr = 0.03 to 1.20%

An average relation between the hardness after hardening (Hh) and the hardness after hardening and tempering (Ht) can be found through: Hh =(T t =167 1:2)Ht 17 ½HRC

where Ht is the hardness after hardening and tempering (HRC), and Tt is the tempering temperature ( C). This equation is valid for 490 C5Tt 5610 C and for a tempering time of 1 h. The tempering temperature for a specified hardness after hardening and tempering is also possible to calculate when chemical composition and the degree of hardening are known (Ref 13):

Fig. 9

Effect of tempering on the true stress in a carbon steel. Source: Adapted from Ref 11

Table 2 Approximate heating times for tempering Per inch of diameter or thickness, with furnace maintained steadily at Tmax, and steel having dark or scaled surface Temperature °C

121 149 177 204 260 316 371 427 482

Heating time, min °F

Cubes or spheres(a)

Squares or cylinders(a)

Average flats(a)

Cubes or spheres(b)

Squares or cylinders(b)

Average flats(b)

250 300 350 400 500 600 700 800 900

30 30 30 25 25 25 20 20 20

55 50 50 45 40 40 35 30 30

80 75 70 65 60 55 50 45 40

15 15 15 15 15 15 15 15 15

20 20 20 20 20 20 20 20 20

30 30 30 30 30 30 30 30 30

(a) In hot air oven, without circulation. (b) In circulation air furnace or oil bath (can be used only in lower temperatures)


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equilibrium. Martensite microstructure is the farthest, followed by bainite, ferrite, and cementite (Ref 14). When the martensitic structure is metastable, there is a natural tendency to transform it to a structure with more stability, and those modifications are accelerated by increasing the temperature during the tempering. The modification that occurs during tempering is complex, and the transformations that take place during the treatment conditions necessary to produce the best mechanical properties combination are a result of accumulated knowledge, not just from the academic point of view but also the practical aspect of observation. Most of the time, the structures developed during isothermal heat treatments are influenced by the low content of other elements besides iron and carbon. Tempering stages Solid-state reactions follow a sequence of precipitation that is related to variables such as:

Diffusivity of the involved element Surface energy of interfaces produced by the reactions

Fig. 10

Crystallographic adjustment (coherence stresses) between the precipitated phases and the ferrous matrix Thermodynamic stability of reactions

During tempering, the martensitic structure is submitted to a sequence of reactions, often superimposed and defined as temper stages (Ref 2, 3, 4, 5, 12). Stage 1. In high-carbon steels, the precipitation of excess carbon begins with the formation of a transition carbide, such as e (Fe2.4C). The e-carbide can grow at temperatures as low as 50 C. Martensite is said to be supersaturated with carbon when the concentration exceeds its equilibrium solubility with respect to another phase. However, the equilibrium solubility depends on the phase. The solubility will be larger when the martensite is in equilibrium with a metastable phase such as e-carbide. Approximately 0.25 wt% C is said to remain in solution after the precipitation of e-carbide is completed. Although most textbooks will begin a discussion of tempering with this first stage of tempering, involving the redistribution of carbon and precipitation of transition carbides, cementite can

Effect of carbon content on the hardness of tempered plain steels. Source: Ref 10

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precipitate directly (Ref 14). This is the case for the lath martensite structure, where the dislocation density can be as high as 1012 to 1013/cm2 (Ref 15). Trapped carbon atoms will not precipitate as transition carbides, but cementite is more stable than trapped carbon. This stage begins at room temperature and extends to 250 C. A fine adjustment between the ecarbides and the ferrous structure is attributed to the precipitation-hardening effect of martensite in high-carbon steels tempered between 50 and 100 C. Stage 2. Tempering at higher temperatures, in the range of 200 to 300 C, for 1 h induces the retained austenite to decompose into a mixture of cementite and ferrite. When the austenite is present as a film, the cementite also precipitates as a continuous array of particles that have the appearance of a film (Ref 3, 5, 12, 14). The

Fig. 11

martensite of the steels with less than 0.5% C content has a retained austenite amount lower than 2%, reaching 6% for 0.8% C. There are some indications that austenite decomposes, turning into ferrite and cementite, but presently a consensus does not exist about whether this structure can be correlated to lower bainite, typically from the isothermal decomposition of austenite, in the temperature range of 230 to 300 C. Stage 3. Tempering at even higher temperatures leads to a coarsening of the cementite particles, with those located at the plate boundaries growing at the expense of the intraplate particles. This precipitation is responsible for the embrittlement phenomenon observed at the temperature of 250 to 400 C. It can be avoided by adding silicon, which is an insoluble element in cementite. This allows cementite formation at

Changes in the mechanical properties of AISI 4340 steel with tempering temperature. Source: Ref 12

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temperatures where silicon diffusion occurs, thus slowing the process. Silicon is a substitutional alloy element that accumulates in the ferrous matrix adjacent to the Fe3C carbide in the growth process. It increases the carbon activity in this region, decreasing the carbon flow to the growing particle and consequently reducing its growth velocity (Ref 5, 16, 17). This silicon effect resulted in the development of alloy 300M, which substitutes for 4340 steel in those applications where it is necessary to use tempering temperatures that cause embrittlement of tempered martensite, which is soon defined. This reaction begins to occur at temperatures on the order of 100 C. Cementite can also be observed during quenching when the Ms temperature is elevated, as is the case of steel with a low carbon content, mainly in the martensite formed just below the Ms temperature. This phenomenon is known as self-tempering (Ref 5). Stage 4. In carbon steels, stage 3 marks the end of the tempering process. Spheroidization of Fe3C occurs as cementite coalesces. This phenomenon is sometimes called stage 4 of tempering (Ref 2). The lath boundary maintains stability up to approximately 600 C. Intense rearrangement occurs between the lath and its low-angle boundaries above 600 C. This recovery process is replaced by recrystallization and coarsening (Fig. 12) at temperatures between 600 and 700 C (Fig. 13).

Effect of Temperature and Alloying. The effect of the tempering temperature on steels with increasing carbon contents can be inferred from Fig. 13. During tempering, the continuous decomposition of martensite to ferrite and carbides changes the state of stress because of continuous dimensional changes. At low temperatures (first stage), a volume contraction takes place as a consequence of e-carbide precipitation. In the second stage, with the transformation of retained austenite (approximately 300 C), the volume is increased. In stage 3, the progressive decomposition of martensite leads to a volume decrease. It is important to observe that the austenitization temperature, which determines the amount of carbon dissolved and the amount of retained austenite, has a strong influence on the expected volume changes (Ref 2, 4, 12, 13). Table 3 shows the changes in length for various steels as a function of tempering temperature. Alloyed steels can also have another stage with the precipitation of alloy carbides, including M2C (molybdenum), M7C3, M6C, M23C6 (chromium rich), V4C3, TiC, and so on, where the “M” refers to a combination of metal atoms. However, all of these carbides require longrange diffusion of substitutional atoms. They can only precipitate when the combination of time and temperature is sufficient to allow this diffusion. The alloy carbides grow at the expense of the less stable cementite. If the concentration of strong carbide-forming elements, such as molybdenum, chromium, titanium, vanadium, and niobium, is large, then all of the carbon can be accommodated in the alloy carbide, thereby completely eliminating the cementite. Figure 14 illustrates the effect of alloying elements on hardness as a function of tempering temperature in carbon steels (Ref 5). Increases in hardness with additions of titanium, vanadium, molybdenum, and chromium are related to the alloy carbide precipitation. This phenomenon is common for tool steels and can affect their toughness, as illustrated in Fig. 15.

Embrittlement

Fe-0.17C alloy quenched in water from 900 C and tempered at 650 C for 5 h. Microstructure shows ferrite grains and spheroidized Fe3C

Fig. 12

Hardness decreases with increasing tempering temperature (Fig. 10, 11). Consequently, yield strength and tensile strength decrease as well. On the other hand, elongation and ductility increase. In this general context, a failure related

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to tempering may be attributed to an incorrect choice of temperature (and/or time of tempering), resulting in an incorrect hardness or low toughness. However, most failures are related to embrittlement phenomena. Quenched and tempered steels are susceptible to a number of different types of embrittlement. Some of them are due to structural modifications during tempering, as previously described. However, there are some due to the interaction of the environment with the quenched and tempered microstructures, such as hydrogen embrittlement and liquid metal embrittlement. Examples of the first type of embrittlement are tempered martensite embrittlement and temper embrittlement, which are described as follows. Tempered Martensite Embrittlement. It is well known that tempered martensite embrittlement (TME) is related to tempered martensite of specimens tempered between 250 and 370 C, as shown in Fig. 16. The impact toughness after tempering at this temperature range is lower than that obtained on tempering at temperatures below 250 C. This type of

Fig. 13

brittleness is inherent to some extent in all steels, including carbon grades. For that reason, medium-temperature tempering is, as a rule, not employed in practice, although it can ensure a high yield limit. According to Krauss (Ref 12), TME may or may not be associated with impurity atom segregation to prior-austenitic grain boundaries, but the most common factor, at least for medium-carbon steels, is the phenomenon that takes place due to decomposition of retained austenite to cementite in the interlath Table 3 Length variations related to metallurgical reactions as a function of tempering temperature ranges Stage

Temperature range, °C

1

0–200

2

200–300

3

230–350

4

350–700

Metallurgical reactions

Precipitation of e-carbide; loss of tetragonality Decomposition of retained austenite e-carbides decompose to cementite Precipitation of alloy carbides; grain coarsening

Source: Ref 18

Hardness as a function of tempering temperature for plain carbon steels. Source: Ref 6

Expansion (E) or contraction (C)

C E C E


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region of martensite plates along the prioraustenite grain boundaries (Ref 19). There are basically three different modes of fracture through the tempered martensite of specimens tempered in the range of 260 to 370 C. First is brittle transgranular fracture, which results from the decomposition of retained austenite in the second stage of tempering, as mentioned previously. Films of retained austenite between laths of martensite in quenched medium-carbon steels transform into thin plates of cementite on tempering. The second mode of fracture associated with TME is intergranular. This kind of fracture is quite common and has been related to phosphorus segregation to the austenite grain boundary

Fig. 14

Fig. 15

Hardness and toughness of a tool steel as a function of tempering temperature. Charpy V-test performed at room temperature in the short-transverse direction

Effects of titanium, vanadium, chromium, and molybdenum on tempering hardness behavior. Source: Ref 5

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during austenitization. However, data on the effect of phosphorus on the fracture surfaces in the untempered condition indicate only that the presence of phosphorus at the prior-austenite grain boundary is not sufficient for the development of TME. So, the interaction between phosphorus and cementite is necessary for the intergranular mode of TME (Ref 20, 21). It is important to understand that the fracture is occurring along a prior-austenite grain boundary, which is now a high-angle ferrite grain boundary (Ref 22). Krauss (Ref 12) describes another type of transgranular fracture mode associated with TME, which is observed in 4340-type steels. It is interlath cleavage induced by cracking parallel to the cementite formed from the retained austenite transformation. Transgranular fractures in TME may be related to the interlath carbide thickness; thinner carbides cause interlath fracture, and thicker carbides promote translath cleavage. In some low-carbon steels, embrittlement is associated with peculiar carbide morphology that provides numerous sites for microcrack initiation, growing by microvoid coalescence and then fracture, with little gross plastic deformation. Silicon additions to carbon steels raise the temperature range in which TME occurs (as mentioned earlier), because silicon delays both the conversion of the transition carbide (e) to cementite within the martensite laths as well as cementite coarsening at boundaries at higher tempering temperatures (Ref 16, 17). The effect of silicon content on the impact properties of 0.6%C-0.47%Mn-0.52%Cr-1.77%Ni0.19%V-1.0%-2.5%Si steels can be seen in Fig. 17 (Ref 16).

Fig. 16

Investigations using AISI 4140 steels, have shown that austenitizing temperature has an influence on the TME phenomenon. High austenitizing temperature was found to be associated with more pronounced TME, favoring brittle failure modes, even in specimens showing virtual absence of phosphorus segregation. These investigations found that high austenitizing temperatures increase carbide dissolution in austenite, apparently due to more intensive carbide precipitation and growth during tempering (Ref 23). When TME appears, there is no heat treatment to reverse the effect, other than to reaustenitize and quench the steel, then temper in a temperature range where TME does not occur. TME is also called one-step embrittlement. Temper Embrittlement. The second type of embrittlement from tempering at high temperatures is known as temper embrittlement (TE). Temper embrittlement occurs when tempering in the high-temperature range of 450 to 600 C. It is not a major problem because it may be avoided simply by quenching from the tempering temperature. This embrittlement can be reversible under high-temperature tempering. When steel that has undergone TE is heated to a temperature above 600 C and then cooled very quickly, its impact is restored. Therefore, such brittleness is termed reversible. Temper embrittlement is also called two-step embrittlement, because two tempering treatments or a heating step and a cooling step are sometimes required to induce embrittlement. Figure 18 shows different situations where embrittlement may or may not appear.

Illustration of toughness loss after tempering in the embrittlement range. Source: Ref 17

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Temper embrittlement of this category of steel takes place due to impurity segregation at the grain boundaries and, finally, decohesion of the grain boundary. This leads to intergranular fracture morphology. It is signaled by a material toughness loss. It is pointed out that grainboundary segregation depends on the alloying elements of the steel. Carbon steels with less than 0.5% Mn are not prone to reversible TE. The phenomenon can only appear in alloy steels. Alloying elements may have a different effect on steel after tempering at the steel propensity to TE. Unfortunately, the most widely used alloying elements, such as chromium, nickel, and manganese, promote TE. When taken separately, they produce a weaker effect than in the case of combined alloying. The highest embrittlement effect is observed in chromium-nickel and chromium-manganese steels. A fundamental fact is that alloy steels of very high purity are utterly unsusceptible to TE, which is caused by the presence of various impurities, such as phosphorus, antimony, and arsenic, in commercial steels. Relatively small amounts of these elements, on the order of 0.01% or less, have been related to TE (Ref 24–27). Steels made of pure elements do not become brittle after tempering as can be seen in Fig. 19.

Fig. 17

Molybdenum is one of the main alloying elements in many low-alloy steels and is an effective method of alleviating TE. Small additions of molybdenum (0.2 to 0.3%) can diminish TE, while greater additions enhance the effect. The mechanism of actuation is related to molybdenum segregation (equilibrium and nonequilibrium) during quenching, and the quenched-in vacancies play a role in the temper process (Ref 28, 29). As mentioned earlier, in chromiummolybdenum steel, phosphorus is the major embrittling element. However, the kinetics

Fig. 18

Schematic diagram showing thermal cycles leading to the reversal of temper embrittlement and the deembrittlement heat treatment

Charpy tests at room temperature as a function of tempering temperature for 0.6% C steel with different silicon contents. Source: Adapted from Ref 16

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of phosphorus segregation in this steel are different from that in steels due to molybdenum. Molybdenum has a strong effect in lowering the solubility of phosphorus in iron. This is interpreted as due to a molybdenum-phosphorus scavenging reaction; so, if molybdenum is free in solution, it appears to either prevent phosphorus from segregating to the grain boundaries or to reduce the brittleness potency of phosphorus at the grain boundary. Although molybdenum is an effective element to reduce the susceptibility for TE, the precipitation of molybdenum as carbide must be taken care of. To avoid that, it is observed that vanadium is added in this grade of steel. Vanadium is a strong carbide former compared to molybdenum and chromium. Vanadium initially forms MC-type carbides; this changes the molybdenum-tocarbon and chromium-to-carbon ratios. The increase in molybdenum-to-carbon ratio is favorable for Mo2C-type carbides, and that of the chromium-to-carbon ratio is favorable for Cr7C3 carbides in this grade of steel. These changes in carbide formation sequences basically slow down the precipitation of molybdenum as carbides. When the molybdenum in solid solution in the ferrite matrix is fully removed,

phosphorus is free to segregate, and the material thus becomes embrittled (Ref 28, 29). Failure analysis of high-temperature studs concluded that the failure related to TE could be delayed due to the presence of molybdenum and vanadium. However, reductions of phosphorus contents in the steel and a grain size of approximately 10 mm could reduce the tendency for brittle fracture (Ref 30). Table 4 shows elements that can segregate to the former austenite boundaries that are now ferrite boundaries (Ref 2). This segregation was shown in a conclusive form through Auger electron microscopy performed on intergranular fracture surfaces. This technique allowed the exact determination of atomic specie concentration segregated in con˚ of fined boundaries with approximately 10 A depth. This fraction varies from 0.3 to 2.0 for steels when these matrix elements are lower than 0.1%. Although there are similarities in the effects of the two types of embrittlement from a practical standpoint, TME and TE are separable into two different phenomena because they occur in two different ranges of temperature and also because TME is a much more rapid process than TE. The former develops during the first hour of the normal tempering period, and it is independent of section size and/or cooling rate after tempering. Second, TE needs many hours to develop, and it is an important concern, mainly for heavy sections that are tempered at higher temperatures (out of embrittlement range) and cooled very slowly over a period of many hours through the critical range of embrittlement. Temper embrittlement develops in relatively soft structures because of the high tempering temperature required to develop it (Ref 12). Temper embrittlement is clearly associated with the high-angle ferrite grain boundaries (former austenite grain boundaries). Thus, it is expected that austenite, which has the necessary concentration of segregants to allow the development of TE when aging tempered martensite, Table 4 Chemical elements that can segregate to the grain boundary Group IV B

Fig. 19

Influence of phosphorus and antimony on roomtemperature impact energy as a function of tempering temperature in a Ni-Cr-Mo steel. Arrow shows the laboratory alloy. Source: Ref 19

C Si Ge Sn ...

VB

VI B

N P As Sb Bi

O S Se Te ...


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will develop TE upon aging a bainitic microstructure (Ref 22). In fact, bainitic microstructures, present in many kinds of low- and medium-alloy structural carbon steels, can also develop embrittlement when tempered (TBE). Tempering at low temperature has only a small effect on the cementite size and morphology. Consequently, the low-temperature embrittlement phenomena are not found in conventional bainitic microstructure (Ref 31). However, for higher temperatures, embrittlement can be developed. The tempering temperature range for TBE depends on the chemical composition and microstructure of the steel (Ref 32). Mechanisms that provoke TBE are similar to TME, that is, precipitation of cementite formed by the decomposition of retained austenite film at boundaries of bainitic ferrite laths or even around the martensite-austenite islands during soaking at the tempering temperature (Ref 33, 34). Secondary quenching (formation of martensite during cooling from the tempering temperature) and transformation of retained austenite into martensite because of plastic deformation after tempering are other possible reasons for the occurrence of TBE. The influence of impurity segregation as well as the as-quenched microstructure on the TE phenomenon in low-alloy steel was verified. The susceptibility of lower bainite or martensite to

Fig. 20

embrittlement was investigated. It was shown that martensitic microstructures are more susceptible to intergranular fracture than bainitic microstructures (Ref 35). Mechanical Tests for TE Determination. As previously mentioned, TE is characterized by decreasing the impact resistance by heating and maintenance in the critical temperature where the phenomenon occurs (450 to 600 C) or by slow cooling through this temperature range. The impact resistance can be recovered by heating up the embrittlement temperature range (4600 C), followed by a quick cooling. The TE apparently does not have an influence on hardness, yield strength, and elongation measured in conventional tension tests. The same can be said about fatigue resistance, although this phenomenon produces a drastic reduction in toughness fracture and also increases the transition temperature, mainly for steels with nickel, chromium, and magnesium additions. Experimental methods that are more adequate for TE studies consider the kind of load, the test temperature, and also the velocity of the imposed deformation. Figure 20 shows the test results of AISI 1340 steel, performed in different tempering temperatures. Figure 21 indicates that the embrittlement phenomenon of tempered martensite is just sensitive to determinate test conditions. The room-temperature tensile test, using unnotched proof tests, does not present changes in terms of reduction of section and tension rupture in the

AISI 1340 steel quenched and tempered (at different temperatures). A, unnotched sample at room-temperature test. B, unnotched sample at very low-temperature test. C, notched sample at room temperature. D, Charpy impact test at room temperature. Source: Ref 19

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pg 300


embrittlement temperature range. It is possible to make an analogy with the embrittlement phenomenon that occurs in ferritic stainless steel having chromium above 15% when exposed to temperatures at approximately 475 C. In some cases, this embrittlement is not detected in tension tests, and almost no modifications occur in elongation during the test, but a strong variation in impact resistance occurs, decreasing the value compared to the same unembrittled alloy steel, as seen in Table 5 (Ref 36). In examining Fig. 21, it is possible to once again realize the importance of the kind of test used to verify the embrittlement phenomenon. As seen in Fig. 20, a common method of detecting loss of toughness during tempering is the impact test, which measures the energy absorbed in fracturing a sample in a specified impact loading for a specified specimen size and geometry. Results can be reported as impact energy (for fracture) at a specific temperature or impact energy as a function of the test temperature. It is possible also to give the impact transition temperature, which is the temperature where the impact energy-temperature curve changes from a relatively high value to a relatively low value. This temperature is often taken as the temperature at the inflection point into the impact energy-temperature curve, as shown in Fig. 21 (Ref 22). In Fe-Ni-C steel, the TME associated with the formation of grain-boundary cementite was observed only when the impact test was performed below the critical test temperature

Fig. 21

(approximately 40 C). Impact toughness is controlled by the intergranular fracture below this critical temperature, while it is controlled by transgranular fracture (i.e., the matrix toughness) above this critical temperature. Temperature is an important test parameter to determine embrittlement during tempering (Ref 37). Temper embrittlement is not detected by simple plots of impact energy versus temperature. However, it is detected by more extensive impact testing that measures variation of the impact transition temperature with the tempering temperature. In order to determine the maximum embrittlement temperature and the boundaries of the brittleness to reversible TE, the secondary tempering method is recommended (Ref 38). It is often possible to tell if a steel has failed because of one of these embrittlement problems by examining the fracture surface. A grainboundary fracture mode is characteristic of both types of embrittlement. However, in steels with Table 5 Mechanical properties of ferritic stainless steel at room temperature under two different heat treatment conditions Mechanical properties

Annealed at 800 °C

Annealed at 800 °C and tempered at 450 °C for 4 h

sR, Kgf/mm2 sE, Kgf/mm2 A, % Impact, Kgm/cm2 (J)

58.4 33.0 22.8 12 (94.5)

59.7 40.0 23.6 1.4 (11.0)

Transition of the fracture behavior of two hypothetical steels in two situations. A, tough. B, embrittled. Source: Adapted from Ref 22


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extremely low levels of phosphorus and sulfur, TME can display a cleavage surface. Although traditionally, TME is usually detected by Charpy tests, a series of experiments were carried out on three commercial steels to explore the possibility of characterizing TME by macro- and microhardness tests. Results indicate distinct hardness peaks in two steels and an inflexion in the other at approximately the TME temperature. These experiments are based on the fact that TME is associated with impurity segregation, and microhardness measurements have elucidated such segregation effects (Ref 39). Interaction of the TE phenomenon with Hydrogen Embrittlement. Hydrogen embrittlement (HEM) occurs when high-strength steels absorb an excessive amount of hydrogen in a variety of environments. Hydrogen can be introduced into the material in service or during materials processing, for example, in the presence of hydrocarbons or hydrogen sulfide or during pickling in acids, plating, welding, and heat treatment (Ref 40). When tensile stresses are applied to a hydrogen-embrittled component, it may fail prematurely. Hydrogen embrittlement failures are frequently unexpected and sometimes catastrophic. An externally applied load is not required, because the tensile stresses may be due to residual stresses in the material. The threshold stresses that cause cracking are commonly below the yield stress of the material. High-strength steel, such as quenched and tempered steels or precipitation-hardened steels, are particularly susceptible to HEM. Tensile stresses, susceptible material, and the presence of hydrogen are necessary to cause HEM. Residual stresses or externally applied loads resulting in stresses significantly below yield stresses can cause cracking. Thus, catastrophic failure can occur without significant deformation or obvious deterioration of the component. Very small amounts of hydrogen can cause HEM in high-strength steels. It is believed that hydrogen is likely to be trapped by structural defects due to its extremely low solubility in the iron lattice. The most commonly recognized structural defects that have a strong hydrogen trapping effect are grain boundaries, dislocations, carbides, and microvoids. In the case of carbides, their size is related to the hydrogen trapping effect. For the boronbearing steel tempered to 1050 and 1300 MPa, fine cementites having an effect on hydrogen

trapping were found. Lower susceptibility to HEM was found when this steel was tempered to a lower strength level, due to delayed onset of brittle intergranular fracture (Ref 41). The fracture mechanism produced by this embrittlement is not simple, because it is related to the imposed load as well as the hydrogen amount in the sample. Tests to determine this embrittlement process are frequently static tests, taking many hours or even many days to obtain the results. This kind of test is sometimes called static fatigue. Figure 22 shows the relation between the crack propagation mode for different load levels. Hydrogen embrittlement susceptibility is related to the hardness of the ferrous matrix, and a very small amount of hydrogen can be deleterious in the case of high hardness of the matrix. Under the influence of hydrogen, the resistance to fracture or crack growth of steels is greatly reduced in gaseous hydrogen or hydrogen-containing environments. Thus, some experiments with 4140 steel under different heat

Fig. 22

Illustration of a crack growing at different load levels. Source: Adapted from Ref 19

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treatments were carried out. Experiment results indicated that the acceleration of the crack growth in hydrogen was more pronounced for quenched and tempered conditions when the tempered temperature range coincided with the TME range of the 4140 steel (230 to 370 C). At higher tempered temperature (approximately 550 C), the influence of hydrogen became insignificant. Specimens in the austempered condition have the best performance on the tests, as seen in Fig. 23 (Ref 42). The effect of the microstructure on HEM was investigated for a low-carbon (Mn-Si-Cr) steel. Microstructure formed by bainite and martensite has better behavior than tempered martensite in intermediate-temperature tempering. At the same strength level, the impact energy of the mixed microstructure is 17% higher than only tempered martensite. It is attributed to the fact that bainite-martensite delays the TME onset (Ref 43). Similar work was performed by Lantsman et al. Experiments with cadmium-plated 65S2VA steel springs were carried out. In this process, the steel absorbs hydrogen, which leads to HEM. Results show that the susceptibility of HEM has a strong dependence on the preliminary heat treatment and structure. With an identical hydrogen content, the austempered steel, with lower internal stresses, will fracture

under a higher load than the steel subjected to standard quenching and tempering (Ref 44). Interaction of the TE Phenomenon with Liquid-Metal Embrittlement. The presence of select metallic specimens on the surface of various alloys can provide for a very detrimental reaction under load, known as liquid metal embrittlement (LME), causing brittle fracture by intergranular cracking. LME is the reduction in elongation to failure that can occur when normally ductile metals are stressed while in contact with liquid metals. Failure of components related to LME is less common than failures caused by other processes, such as fatigue, HEM and stress-corrosion cracking, but a significant number of industrial failures related to LME do occur (Ref 45–47). LME should depend on time of contact with liquid metal while the solid is stressed. When this occurs, the solid metal fails instantly, because the flow of liquid metal into the crack tip during crack growth plays a significant role (Ref 48). The liquid metal affects the fracture behavior at the tip of the crack, reducing the critical stress intensity for fracture and altering the micromechanism of fracture at the crack tip (Ref 49). Sources of the aggressive elements vary, including unintentional or accidental exposures, for example, during fabrication or service when there is intimate contact between the structural and embrittling metals and also when there are tensile stresses above a threshold value. Eventually, small amounts (0.1 g) of embrittling metals can produce extensive cracking by LME (Ref 45). Plain carbon and low-alloy steels may be embrittled by exposure to liquid lead, cadmium, brass, aluminum bronze, copper, zinc, lead-tin solders, and lithium (Ref 50, 51). Other metal alloys also have susceptibility to LME. Table 6 lists embrittling environments for some common structural materials (Ref 45). Table 6 Examples of embrittling metal environments for some structural materials Structural materials

Fig. 23 Ref 42

Impact values of 4140 specimens after various types of tempering treatments. Source: Adapted from

High-strength martensitic steel Austenitic stainless steel Titanium alloys Aluminum alloys Copper alloys Zirconium alloys Nickel alloys Magnesium alloys Source: Ref 45

Embrittling environments

Hg, In, Sn, Pb, Cd, Zn, Li, Cu Zn, Cu, Li Hg, Cd, Ag, Au Hg, Ga, In, Sn, Pb, Cd, Zn, Na Hg, Ga, Bi, Zn, Li, Sn, Pb, In Hg, Cd, Cd-Cs, Zn Hg, In, Li, Zn, Ag Na, K, Rb, Cs, Zn

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Industrial environments provide diverse sources of embrittling metals, for example, during application of coatings or during service if temperatures and stress are sufficiently high and there is close contact with the substrate, as mentioned previously. During soldering, brazing, or welding, LME can be developed if the residual or assembly stresses are sufficiently high. Metal lubricants and overheated bearings are other examples that can provoke LME (Ref 45). Literature presents various cases of failures associated with LME, such as failures in gas turbines, presented by D.W. Cameron (Ref 50). In this case, high temperatures inherent in the gas turbine aggravated the phenomenon. Research has shown that the martensitic steel 91 (9% Cr, 1% Mo) is prone to LME by liquid lead when some conditions are fulfilled (Ref 52– 54). Quenched and tempered steels are susceptible to lead embrittlement, and the conditions to develop this phenomenon are:

Presence of either external or internal lead in the steel Tensile loading Temperature between 200 and 480 C

If any one of these conditions is not observed, LME is avoided (Ref 12). There are a great number of possible mechanisms for this embrittlement. In a general way, it occurs in metals when they are in contact with low-melting metals and can happen even

Fig. 24

Worn surface with excessive generation of heat (arrow)

when the stresses occur at lower temperatures than the melting point. LME is not typical, but it is an important failure mechanism, and industrial awareness of potential problems is still limited.

Case Studies Case Study I: Grinding Cracks. The origin of grinding cracks may be related to low efficiency of the cooling system, microstructure and material cleanliness, and also excessive rate of material removal per pass. The heat generated due to these parameters can produce visual characteristics on the worn surface, as seen in the cam shaft surface shown in Fig. 24. This component (AISI 5160) was induction hardened and tempered. The surface cam shaft heat effect can be studied by metallography and microhardness determinations (100 g load) in a cross section containing the worn surface. Figure 25 presents this microstructure with a low-depth white layer (approximately 50 mm) containing the lower indenter marks, which means fresh martensite (indentations 1 to 3). The fresh martensite has different chemical reactivity than tempered martensite or even other austenite decomposition products. Because of this the overheated surface can be identified by special chemical etching methods, for example, cooper sulfide. Table 7 presents the hardness values.

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Fig. 25

Microstructure close to the worn surface. Etched with 2% nital

Table 7 Microhardness values of the worn cam shaft cross section Identation

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24

Depth, mm

Hardness, HV0.1

Average

40 35 35 95 92 90 155 150 160 220 220 225 405 405 405 610 605 625 1000 1000 1000 Base material Base material Base material

782 803 782 433 433 455 520 493 563 592 606 642 642 782 690 858 690 762 724 772 803 256 230 251

... 789 ... ... 440.3 ... ... 525.3 ... ... 613.3 ... ... 704.7 ... ... 770 ... ... 766.3 ... ... 245.3 ...

Going deeper in the case (right side of Fig. 25) close to the white layer, a darker etched zone (indentations 4 to 9) has lower hardness (bigger indentations), which means that the temperature was sufficient to cause tempering and reduce the local hardness. Indentations 13 to 21 (400 to 1000 mm) show higher hardness (Table 7) related to the process tempering operation, where hardness values are close to the martensite asquenched condition. This condition is deleterious because worked surfaces are more prone to grinding cracks. One example can be seen in Fig. 26. The

Fig. 26

Grinding cracks on AISI 5160 steel cam shaft after induction hardening and low tempering (high-hardness tempering). Original magnification: 200·

observed grinding cracks are small and difficult to detect by nondestructive tests as well as optical techniques classified as macrographic techniques. In this case, they could be seen only through microscopy techniques applied directly on the worn surface or using replica methods. These defects must be avoided, because they may increase the precore failure probability. Machining operations must be carefully performed, but surface compressive residual stresses may contribute to avoiding the grinding cracks. Case Study 2: Transgranular and Intergranular Crack Path. Intergranular cracks are frequently related to the TE phenomenon. However, during low-temperature tempering, very close to the embrittlement range, a predominantly intergranular crack path was found in an AISI 5160 commercial steel after a tempering operation at 2000 C (Ref 55).

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In a recent failure analysis of a compact vehicle suspension spring (chemical composition shown in Table 8) with a specified hardness range of 53 to 55 HRC, a fatigue crack nucleated from a surface defect that had grown quickly (radial marks), as seen in Fig. 27. A scanning electron microscopy image of the radial fracture surface (Fig. 28) shows transgranular and intergranular fracture paths. The spring hardness was on the order of 54 HRC, which means that the tempering process was carried out very close to the beginning of the TE range. However, a significant intergranular crack path was found in the fracture surface. This case shows that the fracture can assume an intergranular path even outside of the embrittlement region. The high hardness means high yield stress, which prompts grain-boundary decohesion. Similar behavior was found when analyzing a carbonitrided sample quenched from 880 C and tempered at 180 C. Figures 29 and 30 show a carbonitrided layer and core microstructure, respectively. The hardness of the carbonitrided layer is in the range of 60 to 62 HRC, while the core hardness values are in the range of 30 to 32 HRC. The carbon-enriched surface (approximately

0.9%) shows an intergranular overload fracture path, and the nucleus low-carbon structure (approximately 0.2%) presents a dimpled transgranular crack pattern (Fig. 31a, b). The higher carbon content increases hardness values, which can promote stresses higher than the cohesive grain-boundary strength or even tempering outside of the embrittling temperature range, as mentioned previously. ACKNOWLEDGMENT The authors would like to acknowledge J.C. Vendramim from ISOFLAMA Ind. Com Equip. Ltd., Brazil, for helpful suggestions in writing this chapter.

Table 8 Vehicle suspension spring analysis Chemical composition, wt% C

0.49

Si

Mn

Cr

S

Ni

P

Cu

1.12

0.65

0.59

0.008

0.21

0.017

0.22

Fig. 28

Secondary electron image showing intergranular (horizontal arrows) and transgranular (vertical arrows) crack path

Fig. 27

Surface fracture spring. The dotted arrow shows the fatigue crack nucleus, and the dashed arrow shows the fatigue-to-brittle fracture transition. The solid arrow shows the surface analyzed by scanning electron microscopy. Original magnification: 6·

Fig. 29

Tempered martensite with some retained austenite (approximately 20%). Etched with 2% nital

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Fig. 30

Base material. Low-carbon martensite (horizontal arrow) and upper bainite (vertical arrow). Etched with 2% nital

Fig. 31

Secondary electron image. (a) Intergranular crack path on the carbonitrided case. (b) Transgranular dimpled fracture in the base material

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REFERENCES

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41. M. Wang, E. Akiyama, and K. Tsuzaki, Hydrogen Degradation of a Boron-Bearing Steel with 1050 and 1300 MPa Strength Levels, Scr. Mater., Vol 52, 2005, p 403– 408 42. J.H. Chuang, L.W. Tsay, and C. Chen, Crack Growth Behaviour of Heat-Treated 4140 Steel in Air and Gaseous Hydrogen, Int. J. Fatigue, Vol 20 (No. 7), 1998, p 531– 536 43. D. Liu, B. Bai, H. Fang, W. Zhang, J. Gu, and K. Chang, Effect of Tempering Temperature and Carbide Free Bainite on the Mechanical Characteristics of a High Strength Low Alloy Steel, Mater. Sci. Eng. A, Vol 371, 2004, p 40–44 44. P.Sh. Lantsman, G.G. Vernovskaya, and V.N. Kudryavtsev, Use of Austempering to Reduce Hydrogen Embrittlement of Steel 65S2VA, Metalloved. Term. Obrab. Met., Vol 5, May 1973, p 73–74 (translated from Russian) 45. S.P. Lynch, Metal Induced Embrittlement of Materials, Mater. Charact., Vol 28, 1992, p 279–289 46. P.L.J. Fernandez, R.E. Clegg, and D.R.H. Jones, Failure by Liquid Metal Induced Embrittlement, Eng. Fail. Anal., Vol 1 (No. 1), 1994, p 51–63 47. R.E. Clegg and D.R.H. Jones, Liquid Metal Embrittlement in Failure Analysis, Mater. Sci., Vol 27 (No. 5), 1992, p 453–459 48. K. Ina and H. Koizumi, Penetration of Liquid Metals into Solid Metals and Liquid Metal Embrittlement, Mater. Sci. Eng. A, Vol 387–389, 2004, p 390–394 49. R.E. Clegg, A Fluid Flow Based Model to Predict Liquid Metal Induced Embrittlement Crack Propagation Rate, Eng. Fract. Mech., Vol 68, 2001, p 1777–1790 50. D.W. Cameron, Failures in Large Gas Turbines due to Liquid-Metal Embrittlement, Mater. Charact., Vol 33, 1994, p 37– 43 51. Liquid Metal Embrittlement, Failure Analysis and Prevention, Vol 10, Metals Handbook, 8th ed., American Society for Metals, 1975, p 228–229 52. G. Nicaise, A. Legris, J.B. Vogt, and J. Foct, Embrittlement of the Martensitic Steel 91 Tested in Liquid Lead, J. Nucl. Mater., Vol 296 (No. 1), July 2001, p 256– 264 53. A. Legris, G. Nicase, J.B. Vogt, and J. Foct, Liquid Metal Embrittlement of the

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of a Martensitic Steel by Liquid Lead and Other Liquid Metals, J. Mater. Sci., Vol 40 (No. 9–10), 2005, p 2459–2463 55. A Reguly, T.R. Strohaeker, G. Krauss, and D.K. Matlock, Quench Embrittlement of Hardened 5160 Steel as a Function of Austenitizing Temperature, Metall. Mater. Trans. A, Vol 35, Jan 2004, p 153–162

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Failure Analysis in Tool Steels Rafael Agnelli Mesquita and Celso Antonio Barbosa, Villares Metals

TOOL STEELS are an important class of steels due to their distinct applications and, especially, their specific heat treating issues. Tool steels are used in various industrial applications that require some kind of mold, die, or mechanical device. Tool steels are also a complex class of steels with compositions close to carbon steel or more highly alloyed grades. Tool steels are generally classified according to three main definitions (Ref 1, 2):

selection of heat treatment or tool design. A helpful way to think of this synergistic relation is to think of the factors as simple multiplying factors. If one of them is zero, regardless which one, the whole result would be zero. Of course, the factors are not so distinctly separate.

Due to the large number of grades and applications of tool steels, several possible classifications have arisen. This division is important for the discussion of heat treating. One of the most well-known classifications is made by the American Iron and Steel Institute (AISI), dividing tool steels into several classes according to application, composition, or heat treatment. This classification is shown in Table 1. Another possible classification for tool steels is their division into four groups according to the final application: hot work, cold work, plastic mold, and high-speed tool steels. The advantage of this division is to deal with fewer groups and to group the diverse grades within common aspects of each application, such as sizes, hardness, operating conditions (chocks, wear, or plastic deformation), and surface-finishing requirements. This chapter follows such a division, but the grade nomenclatures used here are primarily from AISI.

They are used in some forming process or forming operation for metal, ceramic, or plastic shaping. Tool steel properties are only attained after heat treating, normally performed by hardening (quenching) followed by tempering. This class of material is produced according to rigorous melting and processing controls. Therefore, even for grades with chemical compositions very close to low-alloy carbon or engineering steels, the production practices of tool steels lead to substantial property improvements.

According to this definition, heat treating is a key issue for tool steels. In fact, heat treating does have a strong effect on tool life, as described in this chapter. In some applications, three major points are equally related to tool performance (supposing that operation is constant):

Design and manufacture finishing Steel composition and its quality Heat treatment applied

Even though this approach shows the importance of heat treating, it may be too simple for determining the variables related to tool performances. That is, all these aspects interact with each other and may influence the final result. For example, if a good design and proper heat treatment are applied to a tool produced with an imporperly selected steel grade, the result can be awful. The same may happen for an incorrect

Classification of Tool Steels

Table 1 AISI classification for tool steels (Ref 1) Group

Water-hardened tool steels Shock-resistant tool steels Oil-hardening tool steels Air-hardening tool steels High-carbon and high-chromium die steels Tool steel for application in plastic molds Cr-, Mo-, or W-alloyed hot work tool steels Tungsten-alloyed high-speed steels Molybdenum-alloyed high-speed steels Adopted from Ref 1

Symbol

W S O A D P H T M


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Cold work tool steels are employed for tooling operations in temperatures below 200 C, typically at room temperature, and for steel forming (Fig. 1 shows some examples). The most common tools of this class are punches and dies for cold drawing and stamping, knives, thread-rolling dies, and coining or cold forging tools. In these situations, mechanical strength and wear resistance are the main performancerelated properties, which are only obtained through high hardness after heat treating, normally at the 60 HRC level, and the dispersion of coarse carbides in the tool steel microstructure. Grades of AISI class D are the most important examples for such tools, but A and O class grades are also employed.

Cold work tooling is also applied to shock operation applications, such as knives for shear cutting thick plates (normally thicker than 13 mm, or 1/2 in.), chisels, and some powder pressing molds or cold forging dies. In these situations, high toughness is also important, even if the wear resistance is reduced. Grades of AISI class S are the most important example in these applications. Toughness is also important in the previously discussed application, preventing several catastrophic failures or, more commonly, adhesive wear failures, in which microchipping is a very important issue. In these situations, the advent of new materials with lower carbon and chromium contents has shown interesting results. Hot work tool steels are used for applications in which process temperature is an important aspect for the working property of a tooling material. A common limit for hot work working temperatures is 600 C, although lower-temperature applications may also be classified as hot (or warm) work. In hot work tooling, the operating temperature interacts with the steel structure, which is very important for materials selection. Hot work tooling may be divided into three major applications: hot forging, especially for steel forgings; extrusion; and die casting (Fig. 2). The two last groups are mainly employed for aluminum alloys and for producing construction or automotive parts, respectively. All applications normally employ AISI grades H

Fig. 1

Examples of cold work tools. Courtesy of Villares Metals

Fig. 2

Examples of hot work dies for (a) press forging and (b) die casting. Courtesy of Villares Metals

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Failure Analysis in Tool Steels / 313

steel as core tools. However, for large forging dies and tool holders, low-alloyed materials, such as AISI grade 6F3 (DIN 1.2714), may be employed. The most important properties for hot work tooling are hot strength and toughness. Hot strength is normally related to the operating conditions, specifically, how long and at what temperature the tools are exposed. This gives rise to an important concept called tempering resistance, which describes the steel resistance to hardness reduction after exposure to high temperatures. Toughness, which is related to crack end-life situations, is also very important in hot work tooling. Toughness requirements are so important that hot work tool steels normally have carbon contents close to 0.40% and hardness below 52 HRC, both values much lower than that of cold work grades. Plastic mold steels have become a very important division of tools due to the increase in plastic material applications (Fig. 3). Mechanical requirements for strength and toughness are less important than in the previous applications. In general-application molds, the moldmanufacturing-related properties are the most important, because the end life of these molds is rarely attained. Therefore, steel machinability, polishability, heat treating response, and weldability are of special interest. Although less common, there are high-demand applications, such as the processing of corrosive polymers or reinforced plastics, where the reason for steel

Fig. 3

selection is different, and corrosion or wear resistance should be emphasized more than manufacturing-related properties. AISI P20 or P20-modified grades (mainly DIN 1.2738) are the most employed steels in the mold-making industry, delivered in the prehardened condition with 32 HRC. Nevertheless, more specialized applications normally require higher-alloyed steels, such as H13, modified martensitic stainless steel (e.g., DIN 1.2083), or even highly alloyed powder metallurgy grades. Currently, the advance of manufacturing technologies, especially those related to high-speed machining technologies, has increased the application of mold steels prehardened to high hardness, such as 40 HRC or higher. This new level improves mold quality, especially the polishing characteristics. High-speed steels are also tool steels, but they have important differences from other tool steels. They are usually used in cutting tools, whereas the other classes are mainly employed in forming tools. Figure 4 shows some examples of high-speed tools. Regarding the manufacturing characteristics, high-speed steel tools are mainly produced by single companies and in smaller sizes but larger quantities than molds or dies. From the metallurgical point of view, highspeed tool steels are definitely part of the tool steel groups. The AISI classes for such grades include the AISI M and T series, which correspond to the main alloy element present in their composition: molybdenum (M) or tungsten (T).

(a) Plastic molds for a drilling machine body, showing a polishing operation, (b) Baby bath plastic injection mold, after final polishing. Courtesy of Villares Metals.

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pg 314


High-speed tool steel properties are normally aimed at high wear resistance, which is the main demand in cutting operations. Wear conditions are complex in high-speed steel tools, with important aspects being high hardness, normally close to 65 HRC, and the homogeneous distribution of undissolved carbides. Also important is the material tempering resistance, in order to retard hardness reduction at the high temperatures developed in cutting processes. Toughness should be of an acceptable level and reduce chipping and cracking, but this property is low when compared to other tool steels due to extremely high hardness and the amount of carbides in high-speed steel microstructures. In the following sections, several failures and their relation to heat treatment are described. For better understanding, the aforementioned division of tool steels is kept in two main groups, cold work and hot work, which are the main classes where failures are observed. High-speed steels and mold steels are also discussed in some specific cases. However, many aspects of a given class (or for a specific failure) are also applied to other classes; therefore, this chapter uses references to previous or future parts of the text and figures. Such a situation is not very convenient for the reader but is important for comparing diverse aspects in a given situation, which is the final task of any failure analysis.

The first high-speed steels were the T type, with T1 being the pioneer grade. During WW II, the supply of tungsten, mainly from Germany or eastern Europe, was interrupted, leading to the development of molybdenum-rich grades. The first developed grade was M1, which is still employed in the United States, but M2 has become much more important. M2 has a combination of molybdenum and tungsten, in almost equal parts, which leads to very useful properties. This grade is used for several applications. Today (2008), it is the most important highspeed composition. Recently, the increase in consumption of molybdenum, tungsten, and vanadium has increased the cost of such alloys, providing another driving force to the development of new grades. In China, due to the large amounts of tungsten ferroalloys in that country, the use of tungsten-rich compositions has gained new attention. Another important element is niobium (formerly known as columbium). Its large reserves in Brazil motivated the development of niobium-modified grades in that country.

Heat Treating Failures of Cold Work Tools

Fig. 4

Chemical Composition and Main Characteristics of Cold Work Tool Steels. Typical chemical compositions for the most common cold work tool steels are presented in Table 2. AISI nomenclature is used in most cases. However, some newly developed grades, not yet

Examples of high-speed tools. Courtesy of Villares Metals

Table 2 Typical chemical compositions of cold work tool steels Composition, wt% AISI

D2 D3 D6 8%Cr-0.8%C; e.g., VF800AT(a) A2 O1 S1 M2 reg. C PM M3 : 2(b)

DIN

1.2379 1.2080 1.2436

1.2363 1.2510 1.2542 1.3343 ~1.3344(b,c)

UNS

C

Si

Mn

Cr

Mo

W

V

Fe

Others

T30402 T30403 ... ...

1.50 2.25 2.10 0.85

0.3 0.3 0.3 1.0

0.3 0.3 0.3 0.3

12.0 12.0 11.5 8.5

1.0 ... ... 2.1

... ... 0.7 ...

0.9 ... 0.2 0.5

bal bal bal bal

... ... ... Nb = 0.15

T30102 T31501 T41901 T11302 ~T11323(b,c)

1.00 0.95 0.45 0.89 1.28

0.3 0.3 1.0 0.4 0.4

0.3 1.3 0.3 0.3 0.3

5.0 0.5 1.4 4.2 4.2

1.0 ... 0.2 5.0 5.0

... 0.5 2.0 6.1 6.3

... 0.1 0.2 1.9 3.0

bal bal bal bal bal

... ... ... ... ...

(a)Trademark of Villares Metals Company, Brazil. VF800AT is not standardized; therefore, the brand name is given. (b) Obtained by powder metallurgy. (c)“~”, similar to but not exactly the same


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standardized, are referred to in the commercial grade. The high carbon level of almost all these grades is evident (except for S1 due to its use in shock-resistant applications). These tool steels have high as-quenched hardness and, when associated with high alloy content, high hardenability. The combination of high carbon content and alloy elements, especially chromium, leads to the formation of large (also called blocky) carbides within the microstructures of cold work tool steels. Along with a high matrix hardness, normally 60 HRC, these carbides are the main reason for the wear resistance of cold work tool steels. A simple mechanism for abrasive wear is shown in Fig. 5, with an abrasive particle grooving away portions (chips) of tool steel material. Hardness increase reduces the particle penetration and the wear. However, steels cannot be much harder than 65 HRC (approximately 900 HV), which is much less than a typical abrasive particle (between 1500 and 2500 HV). Therefore, another mechanism takes place, this being performed by the presence of carbides. They can be as hard as or even harder than abrasive particles, conferring to the material a higher resistance to wear losses, which means high wear resistance. In typical cold tooling, the wear mechanism is more complicated, being a combination of abrasive and adhesive wear. The aforementioned mechanism is still valid, but toughness is also shown to be important for situations where microchipping or microcracking is present. High-speed steels are similar to cold work steels in this sense, with high hardness and dispersion of nondissolved carbides. The difference is that high-speed steels present higher hot resistance promoted by higher alloy content, which leads to intense secondary hardening at

Fig. 5

high temperatures, as discussed in later sections for hot work steels. Although important or even essential for some applications, high wear resistance is normally accompanied by low toughness. High hardness naturally causes a reduction in toughness due to its effect on material fracture toughness. Carbides normally act as crack initiation sites; therefore, the higher the volume fraction of large carbides, the lower the tool steel toughness. The previous discussion can be illustrated by materials properties and microstructures. Typical microstructures of tool steels used in cold work tooling are shown in Fig. 6. Coarse carbides, larger than 10 mm, are present within the microstructures. Such large carbides are typically the M7C3 type, with “M” representing the metal element that is primarily chromium. They are typical for AISI D grades and are present in lower amounts in 8%Cr-0.8%C new materials, as shown for VF800AT steel (the commercial brand name). This is the main advantage of the higher toughness of such grades. A combination of toughness and abrasive wear resistance leads to important improvements in die life for 8% Cr steels, especially in metalforming operations where adhesive wear is the main operating mechanism. Lower-alloyed grades, such as O1, have almost no carbides in the microstructure. This is caused by the lack of chromium content to combine with carbon. Therefore, they are normally used in less demanding applications in terms of wear resistance. Besides the microstructural aspects, a discussion about tempering curves is also important in order to understand cold work steel properties. Figure 7 shows the tempering curves for some of these grades, selected because they enable

Schematic showing wear caused by the movement of a hard particle through microgrooving. The base material has three different conditions. (a) Tool steel with 50 HRC. (b) After hardness increase to 60 HRC. (c) Combination of high hardness and presence of carbides within the microstructure

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a discussion of all materials. For almost all applications, the high hardness of cold working tool steels is normally achieved from low tempering treatments. AISI D and O steels normally have low secondary hardening. In fact, the cold work application is dictated by this situation. Because the materials show a low potential for hardness retention at hot conditions, their application is limited to low temperatures, typically room temperature. In practical terms, the aforementioned curves show that tempering must normally be performed at temperatures near 200 C, if a hardness of approximately 60 HRC is desired. AISI S grades, herein represented by S1 steel, are also tempered in low temperatures, approximately 300 C. The chemical composition of these alloys is specially designed for this condition, with a high silicon level, because this element

Fig. 6

is known to dislocate temper embrittlement to higher temperatures (Ref 4). A distinct behavior is presented by the newcomer 8% Cr steels. Besides lower carbon and chromium, this class also has higher molybdenum contents than AISI D grades. This enhances secondary hardening, which enables tempering at higher than 500 C and obtains hardness as high as 62 HRC. A substantial improvement in toughness and surface treatment behavior is obtained through this alternative tempering, which is discussed subsequently. Some heat treating failures are observed immediately after heat treatment and appear as small or (usually) large catastrophic cracks. However, other failure types related to heat treatment are only observed during tool use, when one notices premature failure or a lowerthan-normal performance.

Microstructures of cold work tool steels. (a) AISI D6, which is similar to D3. (b) AISI D2. (c) An 8% Cr tool steel with brand name VF800AT. (d) AISI O1. Regions are typical for midradius of a 63 mm (21/2 in.) bar after hardening and tempering to 60 HRC. (a–d) Etched with 4% nital for the same amount of time. Original magnification: 100 · . Source: Ref 3

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Fig. 7

Tempering curves for most common tool steels used in cold working. Tempering curves are obtained after hardening small (25 mm or 1 in.) specimens of all materials with the usual hardening temperature: 920 C for S1, 800 C for O1, 940 C for D6 (similar to D3), 1010 C for D2, and 1030 C for the 8% Cr steel called VF800AT.

For simplification issues, heat-treatingrelated failures are divided by topics that represent the main cause. However, in many cases, this division is not possible. That is because several causes can and do act in synergy, amplifying their effects and thus leading to the observed failure. Nevertheless, the division is kept. It is up to the reader to combine the presented information, keeping in mind the possibility for interaction when solving or analyzing a specific troubleshooting case. Design-Related Failures. The previous discussion of cold work tool steel metallurgy and characteristics explains why this class of materials is so prone to fracture and cracking. Except for AISI S grades, all other materials are very brittle. This is due to their intrinsic nature—the combination of high hardness and primary carbides—and also because cold work steels are used predominantly at room temperature, where fracture toughness of steels is naturally reduced (Ref 5). This fact is illustrated in Fig. 8, where the lower toughness of A, D, and O grades in comparison to H or S steels is obvious. Cold work tool steels are thus prone to failure under stress concentrators, also called stress raisers, that are imposed by tool design or machining. Today (2008), modern software is able to calculate stresses and tool working conditions and can help to reduce stresses and especially localized stresses under some regions

Fig. 8

Comparison of longitudinal Charpy V-notched impact toughness for various tool steel specimens taken from 89 mm square stock and tested at working hardness. Source: Ref 6

of tools. However, several tools are still designed based only on previous experience. Design faults may cause failures in heat treatment but also during tool use, leading to short service life. Failures just after heat treatment normally occur in the presence of some of the following features: heavy sections adjacent to light sections, sharp corners, stamp marks, blind holes, and improperly spaced holes (Ref 7). Several of these faults are illustrated in Fig. 9. Large section-size variation caused the failures shown in Fig. 9(a) and 9(b), while the presence of sharp corners or closely spaced holes is shown in Fig. 9(c). Other stress-concentration effects can also cause or even facilitate cracking. One example

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is presented in Fig. 9(d), showing cracking produced by stamp marks. It is not evident in Fig. 9(d), but this failure also had a contribution of poor machine finishing, because the deep tool mark also acts as an important stress concentrator. Figure 9(e) shows a typical sharpcorner crack. In this case, the corner was filled, but there was a nick in the corner where the cracking began. The shape of this fixture is also poor for steel that must be oil quenched. As in the case of Fig. 9(a), thinner outer regions cool more rapidly, forming martensite first, while the more massive central regions cool more slowly. In some cases, it is not possible to eliminate all the stress-concentration effects from a tool

Fig. 9

design. However, they can be minimized if heat treatment and service failure issues are considered prior to tool design. In other words, the designer should foresee possible heat treatment or service problems at the beginning of tool design. As a result, several failures can be avoided, and service life may be enhanced. In this context, some basic advice is given in Fig. 10. Another possibility for solving heat treatment or service cracking is tool steel selection. Instead of using water-hardening grades, oil-hardening ones are preferred in situations sensitive to quench cracking. In some circumstances, it is possible to apply air-hardening grades, such as

Examples of heat treatment cracking caused by design faults in hot work tool steels. (a) Cold work punch, made of a high-speed steel, that cracked because of the large difference in section. Source Ref 1. (b) The same for a D2 die, also assisted by poor machine finishing. Source: Ref 8. (c) O-type steel die cracked through the sharp corners Source: Ref 8. (d) Failure of die caused by stressconcentration effect of deep stamp marks. Source: Ref 1. (e) Fixture made from AISI O1 tool steel that cracked during oil quenching. A nick in the fillet region helped to initiate cracking. Original magnification: 0.75 · . Source Ref 9

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series A or the new grade type 8% Cr. For example, this could be a solution for the failure shown in Fig. 9(e). Even D grades can be air quenched, depending on die section. However, the continuous cooling transformation diagram of the cooling material should be analyzed for the hardening condition and also the possibility for carbide precipitation on the grain boundaries. Before finishing this subject, one further point should be considered. Stress raisers increase failures in cold work tool steel, mainly due to the intrinsically low toughness of such grades. However, almost any tool displays some localized stress. Depending on working conditions, one or more cracks can be initiated and propagate throughout, fracturing or spalling the tool. Therefore, cold work steels are sensitive to overload failure. In many situations, no problems exist in the steel, the design, or even in the heat treatment; the only cause may be excessive stressing of the tool due to its incorrect use. Surface Damage by Grinding or Electrical Discharge Machining. In the previous section, the intrinsic brittleness of cold work tool steels was discussed, as well as the correlation of tool failures to poorly designed tools (regarding stress raisers). This section discusses surface defects introduced in tool manufacturing by grinding or electrical discharge machining (EDM). Nevertheless, these processes can introduce more than

the macroscopic stress raiser effect due to two major factors, described as follows. First, both grinding and EDM cause local heating in the tool surface that, depending on operational conditions, induces local tempering or, far worse, reaustenitizing, quenching, and hardening. The high carbon of these grades, normally more than 1%, promotes high hardness, more so than in other lower-carbon tool steels (Fig. 11). As a result of heating and martensite transformation, small cracks, normally hard to see with the unaided eye, may also be formed, acting as stress raisers during tooling and thus enabling premature cracking. Secondly, after such a metallurgical transformation on the tool surface, the microstructure will be predominantly untempered martensite (also known as fresh martensite). This microstructure is very brittle, especially in high-carbon steels such as the cold work grades. The pre-existing cracks or other cracks formed during tool operation are much more prone to propagate, thus accelerating tool failure. Grinding and EDM cause heating and actually act as a heat treatment applied to tool steel. In the next sections, grinding and EDM are treated separately, with some advice for avoiding problems. Incorrect Grinding. Hardenable steels are more prone to grinding cracks than low-carbon, low-alloy steels. Cold work tool steels and highspeed steels are the most sensitive tool steel grades to such problems due to the high hardness

Fig. 11

Fig. 10

Simple possibilities for avoiding (a) sharp corners and (b) large variation in section. Source: Ref 10

Effect of carbon content on the hardness of different microstructures. Martensite hardness increases rapidly with carbon content. Reaustenitizing and quenching, which can occur in the surface of ground or electrical discharge machined tools, can cause high hardness and brittleness in highcarbon grades such as cold work steels, leading to tool failure. Source: Ref 1

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of as-quenched martensite and its low toughness. Examples of grinding cracks are shown in Fig. 12. The formation of grinding cracks can be explained as follows (Ref 11). Almost all the energy used in grinding is converted into heat, partly through pure friction and partly as a result of deformation of the material. If a correct grinding wheel has been chosen, most of the heat will be removed in the chips, with only a smaller part heating up the workpiece. Incorrect grinding of a hardened tool steel can result in such a high temperature at the ground surface that the tempering temperature of the material is exceeded. This results in a reduction in the hardness of the surface, causing low performance when the tool is used in field applications. However, in addition if the temperature is allowed to rise further, the hardening temperature of the material can be reached, resulting in rehardening. Rehardening during the grinding operation produces a mixture of nontempered and tempered martensite in the surface layer, together with retained austenite, as shown in Fig. 13(a). The affected layer normally shows white under optical microscope examination (after metallographic preparation and acid etching); this denotes the presence of untempered martensite, which is more corrosion resistant than tempered martensite. The diagram in Fig. 13(b) shows the hardness profile through the surface of a cold work tool steel, incorrectly ground in such a way as to produce rehardening. The surface exhibits a high hardness due to the untempered martensite. An overtempered zone occurs just below the surface, where the hardness is lower than the basic hardness of the workpiece. The following hints may provide a solution to grinding problems. Incorrect grinding, resulting in a modified surface layer, often reveals itself through burn marks—discoloration of the ground surface (as indicatedted in Fig. 12b). In order to avoid burning and grinding cracks, it is necessary to keep down the temperature of the ground part, for example, by means of good cooling, and to employ properly dressed grinding wheels that cut the material with sharp cutting edges instead of simply generating heat through friction (Ref 11). The majority of grinding operations leave residual stresses in the ground surface, usually being at a maximum close to the surface. The first and most common effect of such stresses is the occurrence of cracks. Stresses can cause

permanent deformation of the ground part when grinding thin materials. This may be accompanied by retained austenite formation, which

Fig. 12

Examples of grinding cracks. (a) Two views of an S1 tool cutter die cracked and spalled after grinding. Asreceived (left) and after magnetic particle testing (right), accentuating the cracks Source: Ref 9. (b) A D2 die that cracked due to incorrect grinding (arrow indicates grinding marks) Failure was also assisted by closely spaced holes and electrical discharge machining procedures. Generally, grinding cracks are not as easy to see as this. It is usually necessary to examine the part under a microscope or with magnetic powder inspection in order to see the cracks. Source: Ref 8

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enhances deformation and the possibility for cracking. Three avenues are available to reduce grinding assisted failures. First is to control the heat emerging from the grinding operation by use of a proper cooling process. Secondly, the effect of grinding stresses and problems can be reduced by stress-relief tempering after grinding. This also involves the tempering of some regions of untempered martensite, if present. The treatment temperature should be approximately 30 C below the previous tempering temperature to avoid any risk of reducing the hardness of the workpiece. Third, another way of reducing grinding stresses is to tumble or blast the ground parts (Ref 11). Obviously, if the heat damage is too high, that is, cracks, stress relief may not help. Incorrect EDM. Electrical discharge machining is often used in the production of cold

work tools for various reasons. Cold work tools have an intrinsic high wear resistance and are normally difficult-to-cut materials under regular machining processes, such as milling. Finishing the die making with EDM may be an interesting solution, especially for complex-shaped tools. However, new developments in high-speed machining that feature low stock removal and high frequency have been used on dies as hard as 60 HRC. The use of EDM on hardened steels, however, can produce a shallow, rehardened layer of rapidly quenched as-cast structure and untempered martensite at the surface, beneath which is a layer of tempered martensite (Fig. 14). The EDM surface layer is known as the white layer because of its lighter appearance under optical microscope observation of etched samples (this is caused by the higher corrosion resistance of untempered martensite). Normally, the white layer contains microcracks that can grow into serious cracks when the tool is loaded in service (Ref 7). When used on hardened steel, EDM also adds surface stresses to the already established residual stresses; the origins of such stresses are the thermal and phase transformation dimensional variations that occur in EDM surfaces. The temperatures developed in such regions are so high that local melting and resolidification occur, as shown in the upper-left microstructure in Fig. 15(c). In summary, the EDM white layer has four major problems that can enable or accelerate die failures: high hardness, residual stresses,

Fig. 14 Fig. 13

(a) White layer on a tool surface rehardened by an incorrect grinding procedure. (b) Typical hardness profile in regions close to cracks. Source: Ref 11

Electrical discharge machining (EDM) white layer found on a die surface made of AISI D6 (similar to D3) tool steel. Note the white aspect of untempered martensite caused by the EDM process and the presence of small cracks in this layer. Original magnification: 500 · . Courtesy of Villares Metals

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Fig. 15

(a) A2 tool steel blanking die, 63 mm (21/2 in.) in diameter by 13 mm (1/2 in.) thick, that cracked in service because of a brittle zone that had formed during electrical discharge machining (EDM) of the cavity at center. Arrows point to cracks emanating from the cavity. Source: Ref 7. (b) Tool failure due to the same reason, where the 3.2 mm (1/8 in.) holes were produced by wire-EDM. (c) The effect of EDM on surface microstructures and approximate hardness of the tool shown in (b) are presented. Etched with 3% nital. Central image in lower magnification; all other images in the same magnification. Source for (b) and (c): Ref 9

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coarse (as-cast) microstructure, and frequent precracked regions. Some typical examples of die failures assisted by incorrect EDM are shown in Fig. 15. In Fig. 15(a), a 63 mm (21/2 in.) diameter by 13 mm (0.5 in). long blanking die for a small part cracked from the corners of the EDM cavity after the die had produced 20,000 pieces. The die was made of heat treated A2 tool steel. In Fig. 15(b), the surface of an AISI A4 cup plate is shown, with spalling at one of the holes, which were made by EDM. A laboratory investigation of this failure led to the typical appearance of EDM-assisted failures: a coarse, white surface layer that is very high and brittle, due to the presence of an untempered martensitic matrix and net carbides. Below, unquenched martensite is observed, followed by a region of overtempered martensite, after which the normal (core) microstructure is observed. In many situations, the affected regions are not exposed as clearly as in Fig. 15(c). However, the typical white layer is always present on the tool surface, as shown in Fig. 14. The coarser this layer, the higher the probability for tool failure, due to its brittleness and the fact that the EDM white layer likely possesses cracks. Three major practices are recommended for avoiding premature cracking caused by EDM:

Reduce the stock removal when finishing the EDM process (if low-frequency EDM was used for roughing, high-frequency should be used for finishing). This is helpful for minimizing the depth of the rehardened white layer. The white layer should be eliminated or minimized by light grinding or lapping. This procedure is time-consuming but, in many situations, can lead to an impressive extension of tool life, especially when white layers are thick and the tool is crack-sensitive. For relieving stresses in EDM-processed dies and increasing the toughness of the remaining white layer, a new tempering treatment should be performed. Its temperature should be 30 to 50 C below the maximum tempering temperature used in the heat treatment, to avoid hardness loss. Normally, this procedure is easy to apply and therefore is highly recommended.

Although both grinding and EDM can damage the tool surface, EDM problems are much more common in industrial tool failures. This occurs in particular for cold work tool steels and

high-speed steels, where the surface white layer has high brittleness and the base material (i.e., the tool steel) has low resistance to crack propagation. However, EDM-assisted failures are also observed in hot work dies. Also important to mention here is the effect of incorrect EDM in plastic molds. Although the mechanical stressing is normally low, surface finishing (by polishing or texturing) is crucial in this application because plastic injected parts are able to reproduce any problems on the mold surface. Thus, EDM defects may cause serious quality problems to injected parts, impairing the mold application. Failures due to the Heat Treating Procedure of Cold Work Steels. The heat treating procedure can itself deeply change the microstructure and properties of all tool steels, not only the cold work grades. This may occur even if the specified hardness is obtained. This section deals with failures caused by improper heat treating procedure and is divided into the three most common causes in cold work tool steels: the use of incorrect temperatures, the use of excessively short tempering times (or even no tempering at all), and the formation of excessive amounts of retained austenite, caused either by improper hardening or incorrect tempering. Incorrect Hardening or Tempering Temperatures. As for other tool steels, the same class of cold work tool steels may present important differences in the indicated heat treating temperatures. If the temperature is higher or lower than that indicated for a certain grade, mechanical properties may be altered, especially for toughness. Thus, the die performance is also strongly influenced. This section describes this effect in a specific grade—the 8% Cr tool steel, which has been highly employed in tools that traditionally use grades from the D or O series. As discussed previously, the 8%Cr-0.8%C steels have a distinct combination of toughness and wear resistance that makes these grades very suitable for cold work tooling. However, their heat treating temperatures are considerably different from that used in the usual grades. To illustrate this effect, an 8% Cr steel was chosen (commercial name VF800), and various temperatures were used for its heat treatment. The composition of this grade is shown in Table 2. Such conditions were analyzed in the laboratory in terms of microstructure and mechanical properties (measured by a bend test, Ref 12). Four conditions were applied, as

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follows, and summarized in Fig. 16. For all, the hardness was maintained at 60 HRC:

Condition 1: standard-condition VF800AT grade, with hardening temperature approximately of 1030 C and tempering at high temperatures, 540 C (twice, for 2 h) Condition 2: low hardening and low tempering temperatures, 970 and 200 C, respectively. This condition is typical for high-chromium and high-carbon D grades, such as D3 and D6. Condition 3: typical hardening but lower tempering temperature. Although not indicated for 8% Cr steels, this condition is typical for D2 steel, a well-known grade for heat treaters. This is used in D2 for attaining the 60 HRC level, due to the weak secondary hardness of this grade. However, the 8% Cr grades normally have higher alloy content in terms of molybdenum or vanadium, allowing 60 HRC to be obtained after high-temperature tempering. In the case of VF800AT, up to 63 HRC is possible, depending on the hardening condition (Ref 13). Condition 4: hardening temperature higher than normal, using a condition typical for high-speed steels (when treated to 60 HRC)—hardening at 1150 C and tempering at 570 C

The results of mechanical properties and microstructures for all conditions are shown in Fig. 16. A substantial reduction is observed for conditions 2 to 4 compared to the material treated under normal conditions (1). This difference in mechanical properties may be understood based on the relative microstructure for each condition. The first condition has a relatively dark martensitic matrix and dispersion of primary carbides, undissolved during the hardening treatment. This is typical for this material. The dark matrix indicates hightemperature tempering, where stress relief of martensite transformation is well performed; the dispersion of primary carbides is important for wear resistance. In the other conditions, the microstructures show a different aspect. In conditions 2 and 3, tempering at low temperatures is denoted by less intense etching, converting to a lighter matrix. In these cases, hardness is produced by a highly unstable and stressed martensitic structure instead of the secondary hardening of hightemperature tempering (adequate condition).

This reduces the toughness to the observed levels. The last condition, 4, produced the lowest toughness values. This is due to the intense grain growth produced by the high hardening temperatures, which are typical for high-speed steels but not applied for this grade. In this microstructure (Fig. 16e), the coarse martensite plates reflect coarse austenite grains. The 8% Cr cold work steels, as with other cold work steels, have much lower alloy content than high-speed steels. This reduces the pinning effect of carbide precipitation on grain boundaries, thus causing rapid grain growth when high-speed steel hardening temperatures are used. Excessive high hardening temperatures are also common problems in heat treating highspeed steels. Hardening temperatures for these steels are close to the solidus temperature (less than 50 C, 90 F), above which liquid formation starts within the microstructure (in carbide rich areas), leading to expressive embrittlement. Hardening temperature control is thus very important. For example, M2 high-speed steel hardening temperature is about 1200 C (2192 F), but exceeding 1220 C (2228 F) may cause loss of toughness without benefits to hardness, and above 1240 C (2264 F), liquation is likely to occur (Ref 38). Figure 17 shows two cases of an incorrect heat treatment procedure applied to an 8% Cr tool steel, for two punches and a cutting blade, that cracked prematurely. The microstructure observed was close to condition 3 of Fig. 16, but the microstructure was very difficult to observe after regular (nital) etching. A stronger etching condition was applied. Besides tempering at low temperature, there were evidence of hardening overheating (coarse austenite grain sizes). This is therefore a combination of two incorrect situations—conditions 3 and 4. Excessive Retained Austenite Content. Hardening of tool steels involves the transformation of an initial phase, austenite, formed during heating (austenitizing treatment). The following transformations are directly dependent on austenite composition. Martensite formation is of particular interest, because this is the expected phase after quenching. Martensitic transformation is distinct from the usual solidphase transformations, because it does not occur by the diffusion process. The formation of martensite depends on the temperature attained; therefore, two important temperatures are defined: the start and finish of martensite transformation, determined by the Ms and Mf

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Fig. 16

(a) Bend strength and fracture energy (energy necessary to fracture the specimen) obtained in a static bend test. Four-point bend test with specimens of 5 mm (thickness) per 7 mm (width) cross section. Tested material is an 8% Cr cold work steel (brand name VF800AT, Ref 13), heat treated to 60 HRC under four different conditions, 1 to 4 (see text). The legend indicates the hardening (hard.) and tempering (temp.) temperatures, all for 30 min and twice for 2 h, respectively. (b) to (e) Respective microstructures for conditions 1 to 4 after etching with 4% nital for 10 s. All regions refer to the midradius of a 60 mm bar. Source: Ref 12

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temperatures, respectively. If no transformation takes place during quenching, martensite is formed gradually after the Ms temperature is reached, and total transformation of austenite to martensite takes place when the part reaches the Mf temperature. In tool steels, most compositions have Mf values below room temperature; as a result, part of the austenite is not transformed to martensite. This austenite fraction is called retained austenite. It is distributed within the material microstructure; after etching, it is observed as a light matrix crossed by plates or laths of martensite (Fig. 18). As is discussed later, higher hardening temperatures (overheating) lead to larger amounts of retained austenite. In the microstructure, it becomes more evident due to the presence of large martensite plates (resulting from larger austenite grain sizes) crossing the austenite phase within the matrix (Fig. 18b).

Fig. 17

Retained austenite content is directly dependent on the chemical composition, because it determines the Ms and Mf values. The following equation (Ref 14) quantifies this dependence for Ms, since it is also a similar rate for Mf. It is important to observe that all the alloy elements (with the exception of cobalt) reduce the Ms values, especially carbon, thus increasing the amount of retained austenite: Ms ( C)=539 423%C 30:4%Mn 12:1%Cr 17:7%Ni 7:5% Mo (all elements in weight percent) (Eq 1)

Consider again the phase transformation taking place, keeping in mind the effect of alloy composition on Ms values. As explained, austenite will be transformed to martensite because of the rapid cooling of the quenching process.

Examples of failures in an 8%Cr-0.8%C tool steel, caused by an incorrect heat treating procedure. (a) Punches and (b) their microstructures. (c) Cutting blade and (d) the microstructure observed in its failure analysis. The primary cause of failure is low-temperature tempering and high-temperature hardening, a combination of incorrect conditions 3 and 4 shown in Fig. 16. Microstructures etched with Villella’s reagent. Courtesy of Villares Metals


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Therefore, the composition that affects the Ms value is the austenite composition at the moment of martensite transformation, not the alloy

Fig. 18

Retained austenite in two cold work tool steels after hardening and tempering. (a) D2, with 60 HRC. (b) O1, with 54 HRC. For both, the retained austenite content is higher than expected (due to overheating in the hardening treatment). (c) O1 punch from which the microstructure in (b) was obtained. Cracked in service after short life. Retained austenite and carbides are lighter in the microstructure after etching, but they can be differentiated; carbides are either faceted or round, but retained austenite has no delineated area, being within the matrix. In both images, martensite plates that cross retained austenite regions are clearly observed. In (a), the retained austenite content is high enough to reduce the desired hardness from 60 to 54 HRC. Compare to the usual microstructure of O1 and D2 in Fig. 6.

composition. For monophase steels, such as low-alloy, low-carbon grades, the austenite composition is practically the alloy composition. However, in tool steels, especially cold working tool steels and high-speed steels, high amounts of carbon and alloy elements are trapped in the undissolved carbides. The release of these elements is only possible through carbide dissolution, which depends on time and particularly on the austenitizing temperatures involved. The higher the austenitizing temperature (hardening temperature), the larger the amount of carbon and alloy elements that go into solid solution, lowering the Ms temperature (according to Eq 1) and thus increasing the amount of retained austenite. Examples are shown in Fig. 19 for a chromium high-carbon tool steel and for D2. As the hardening temperature increases, the Ms value decreases (Fig. 19a), and the amount of retained austenite increases (Fig. 19a, b). Although the basic mechanism of retained austenite in tool steels has been described, its effect on materials properties has not been explained. In general, retained austenite is undesirable (Ref 17). First, it is softer than martensite, due to the crystallographic nature of austenite (face-centered cubic lattice). However, the effect of the amount of retained austenite on the as-quenched hardness is normally only a few HRCs, unless a strong deviation in heat treating procedure has occurred. Secondly, the retained austenite may cause lower toughness. Due to its unstable nature, austenite at room temperature tends to transform to martensite if adequate thermodynamic conditions are established; one possibility, besides the temperature reduction, is stress. Such conditions may be imposed during tool work that involves elevated localized stress. In this condition, austenite transforms to martensite (untempered, of course), leading to brittleness. Because retained austenite is distributed along all regions of the steel microstructure, normally close to carbides (microsegregated areas), its transformation to martensite may deeply embrittle working regions, causing cracks or fracture failures. Third, this transformation of retained austenite also causes a volume increase, due to the lattice difference between austenite (more compact structure) and martensite. This impairs dimensional stability, which is critical for several tools, especially those that work with small clearances, such as fine blanking dies.

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(a)

(b)

Fig. 19

(a) Influence of austenitizing temperature on martensite transformation of a tool steel containing 1.1% C and 2.8% Cr. Higher austenitizing temperatures lower Ms temperatures and increase the amount of austenite retained at room temperature. Source: Ref 15. (b) Amounts of microconstituents in D2 tool steel as a function of austenitizing temperature. Specimens austenitized for 30 min at temperature and oil quenched. Composition: 1.60% C, 11.95% Cr, 0.33% Mn, 0.32% Si, 0.79% Mo, 0.25% V, 0.18% P and 0.010% S. Source: Adopted from Ref 16

Fig. 20

Hardness and retained austenite as a function of tempering in A2 tool steel. Source: Ref 18

Besides these three undesirable effects of retained austenite, the cause of excess retained austenite should also be considered when analyzing a failure. As discussed and shown in Fig. 19, the increase in retained austenite content in cold working tools is normally caused by exceeding the recommended hardening temperature. However, the retained austenite content may also be reduced when it is converted to martensite (or bainite) after the first tempering.

Transformation of retained austenite depends on the tempering temperatures, as Fig. 20 shows for A2 tool steel. For the purpose of converting retained austenite, tempering must be carried out immediately after quenching to avoid stabilization of retained austenite (Ref 15) and a second tempering must always be applied. At the second tempering temperature, this untempered martensite (also known as fresh martensite) or bainite is tempered, being the final microstructure free from hard brittle phases. For some cold work steels, such as AISI A or 8% Cr steels, tempering can be conducted at temperatures (4500 C) that practically eliminate retained austenite. However, in AISI O or D steels, tempering to 60 HRC or more is normally done at lower temperatures (5350 C), where some amount of retained austenite will exist after tempering; in these cases, the control of retained austenite content should be done by not exceeding the hardening temperatures. Nevertheless, in some cases, it is possible to lose some hardness by using higher tempering temperatures to reduce the amount of retained austenite. This can be done for D2 steel if tempered at approximately 520 C (Fig. 7); hardness would be up to 58 HRC, but the amount of retained austenite decreases from approximately 15% to less than 5% (Ref 18), considerably improving toughness (Ref 19). Application of cryogenic treatments is also a possibility to reduce retained austenite content, especially in cold work tool steels that

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are tempered at low temperatures. In cryogenic treatments, tools are led to very low temperatures, approaching or crossing the final Mf temperature. They are normally performed after a stress relief (at about 150 C) to avoid tool cracking. After the cryogenic treatment, tempering is also necessary to avoid brittleness from the just formed fresh martensite. In summary, the presence of retained austenite in unusually large amounts is an indication that either the hardening or tempering treatment has been inadequately conducted. Other sources of brittleness can emerge in such situations, besides the retained austenite itself. For example, higher hardening temperatures promote coarse grain sizes and also may increase the potential for carbide precipitation on austenite grain boundaries; both effects cause intense embrittlement (condition 4, Fig. 16). Incorrect tempering, with shorter times (see next section) or incorrect temperatures (previous section), changes the stress relief of the martensite structure and the whole strengthening mechanism, thus affecting material toughness as well (condition 2, Fig. 16). Therefore, several examples are observed of industry failures assisted by retained austenite that emerged from incorrect procedures. Figure 21 presents two examples. The first, (Fig. 21a, b) shows an AISI O6 tool that cracked after limited service. Retained austenite is clearly observed in its microstructure. The second tool in Fig. 21 is an AISI S7 die. This grade has lower carbon (~0.50%) and much lower undissolved carbides in comparison to the O- or D-series steels. It should be less prone to retained austenite formation and to the effect of incorrect heat treating conditions. However, this tool was carburized, and a surface pickup of carbon took place, leading to a reduction of Ms and Mf and thus causing the high amount of retained austenite. In both cases, overaustenitizing conditions were employed, enabling the existence of such high retained austenite content and leading to embrittlement. Excessively Short or Absent Tempering. “Tempering, the final heat treatment step applied to tool steels, is defined as the heating of a martensitic or hardened steel to some temperature below A1 temperature (initial temperature of austenite formation); this step produces the final structure and mechanical properties of a hardened steel.” This citation, from Ref 1, briefly explains the importance of tempering treatment. However, in practical situations, this

is not so obvious. After austenitizing and quenching, the steel is hardened to a very high hardness—in many cases, the highest hardness possible to attain for a given steel. After tempering, no significant differences can be observed in hardness measurements, especially for cold work steels, which have a work hardness very close to the as-quenched hardness. This can lead to several problems regarding the embrittlement of a tool caused by poor tempering practice. During tempering, several solid-state phenomena occur simultaneously in the steel microstructure. Depending on the alloy content of a tool steel, the tempering curve presents a different aspect, as shown in Fig. 22 by a division in classes. Class 1 is typical for high-carbon, lowalloy tool steels, class 2 for high-chromium cold work steels, class 3 for high-speed steels with strong secondary hardening, and class 4 for hot work tool steels. Typical cold work steels, from AISI D or O series, will have curves close to classes 1 or 2. As shown and discussed in Fig. 22 and quantitatively in Fig. 7, both AISI O and D grades must be tempered at lower temperatures to attain the highest hardness levels—60 HRC or higher. Tempering is typically conducted at temperatures of approximately 200 C. In practical situations, this is not that simple. The heat flux at such low temperatures is also low, and thus, tempering of dies may take several hours. As a consequence, it is common to find examples of tools that were insufficiently tempered. Two cases are shown in Fig. 23. Avoiding these kind of failures is not technically difficult, but it is time (and money) consuming. A basic rule should be observed: Any operation should not be conducted with the steel in the as-quenched condition. If the tempering time is unknown, a rule of thumb to observe is that a steel is hardly ever overtempered (if the time is exceeded) but can be easily undertempered (if the time is too short). This happens because the tempering parameter, as described in Fig. 22, has the time in log scale. However, the most desirable situation is to have good control of the tool temperature during heat treating. One example is monitoring tool heating and treatment time with thermocouples attached to the part; ideally, all tool regions should be maintained at temperature for approximately 2 h. This discussion of tempering times is important, due to the intense microstructural

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transformations that occur in the initial tempering stages. Carbon rearranges into martensite crystals, and transition carbides are precipitated that are only 2 to 4 nm in size. These transformations considerably enhance material toughness, but hardness decreases only slightly from the initial as-quenched hardness. In conclusion, low-temperature tempering must not be suppressed based on hardness but kept due to the necessary changes in steel microstructure that it promotes.

Fig. 21

Heat Treating Failures of Hot Work Tools Chemical Composition and Main Characteristics of Hot Work Tool Steels. Heat treating of hot work tools is usually more critical than for cold work tools. Hot work tools are normally larger and have higher machining costs, besides being applied to high-demand applications. Although cold work tools may also be critical, such as for drawing dies and industrial cutting blades, they usually do not exceed

(a) AISI O6 graphitic tool steel punch machined from centerless-ground bar stock that cracked prematurely. (b) Microstructural examination revealed an overaustenitized structure consisting of appreciable retained austenite and coarse plate martensite. (c) Failed AISI S7 jewelry striking die showing cracks (arrows) that formed shortly after the die was placed in service; (d) Its microstructural examination revealed that the surface was slightly carburized and the die had been overaustenitized. Note coarse plate martensite and unstable retained austenite. Source: Ref 9

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more than 500 kg. As discussed in the previous section, it is useful to first describe the heat treating conditions of the common grades and then to discuss failure analysis. For this reason, Fig. 24 shows tempering curves for the most common grades, whose chemical compositions are shown in Table 3. By analyzing chemical compositions, one can anticipate several characteristics of hot work tool steels and their differences from cold work steels. First, hot work steels normally have lower carbon, leading to lower as-quenched hardness; also, undissolved carbides will be much lower or even nonexistent in these grades. Thus, abrasive wear resistance, which depends on hardness and coarse primary carbides, is much lower in hot work steels. Final hardness in hot work tool steels is normally determined by desired toughness instead of wear resistance. In most situations, hardness is between 40 and 50 HRC. Other properties, in addition to mechanical strength at room temperature and wear resistance, are important for hot work tool steels. The main metallurgical properties are toughness and strength at elevated temperatures; this last property is dependent on the hardness stability at high temperatures.

Fig. 22

Schematic diagram of hardness versus tempering temperature (assuming constant time at each temperature) or versus a time-temperature tempering parameter for four major types of tempering response in tool steels. Classes 1 to 4 are typical, respectively, for low-alloy tool steels, cold work chromium steels, high-speed steels, and hot work tool steels. Source: Ref 1

Toughness in hot work tools is an essential property for avoiding cracks and fractures that can be common in normal working conditions. Unexpected and very unstable in cold work tools, cracks may be encountered after a given operation time for hot work dies. The most important example is fine crack networks, large in number but small in length, observed on a tool surface after the tooling operation. This kind of crack is known as heat checking and is typically found in die-casting dies (Fig. 25) as well as in forging dies or other hot work tools. The cause is thermal fatigue at the tool surface, caused by repeated temperature fluctuations (heating and cooling) during tool operation. A full explanation of this behavior is found in the literature, for example, Ref 20 to 22. Depending on the intensity, heat checking can lead to tool failure, which is typically the case in die-casting cavities (Fig. 25). In such situations, the casting metal, normally aluminum, enters into the heat checking cracks, making part extractions difficult or impairing part surface quality. In cases of a lack of a toughness, gross cracking may also occur, leading to complete loss of the working tool. High mechanical strength at high temperature is also fundamental in hot working steels. High temperature means temperatures that can affect steel microstructures, normally above 500 C. Depending on time and stress conditions, such high temperatures lead to a decrease in hardness that accelerates several die failure mechanisms, especially hot wear, plastic deformation at working temperatures, and heat checking. The ability of a given grade to maintain hardness at prolonged time at high temperatures is referred to as tempering resistance. A simple (basic) comparison of tempering resistance may be done by analyzing tempering curves. Higher tempering resistance is related to curves dislocated to the right (Fig. 24). As shown schematically in Fig. 26, such hightemperature hardness is only possible by a very important straightening mechanism—secondary hardening, promoted by the precipitation of fine alloy carbides. The stronger the secondary hardening (meaning more intense carbide precipitation), the higher the tempering resistance of hot work tool steels. Such precipitation intensity depends on the amount of alloy elements in solid solution, which is related to the alloy composition and heat treating practice. Figure 27 explains this effect by comparing the addition of different alloy contents to carbon steels; a stronger precipitation hardening is

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obtained by molybdenum, vanadium, or tungsten alloying. Therefore, the desired mechanical properties of hot work tool steels are only attained if a proper heat treatment is applied. For example, adequate secondary hardening will only be present if alloy elements are in solid solution in the as-quenched structure (usually martensite). This is only possible through an adequate austenitizing treatment for hardening, capable of dissolving the alloy elements present in the form of carbides in the initial (annealed) state. In addition, a proper hardening procedure should avoid excessive grain growth and grain

boundary embrittlement by carbide precipitation. Tempering, on the other hand, should eliminate retained austenite and promote adequate precipitation of the alloy carbide. In summary, tool steel properties and the expected performance are only possible after a quality heat treatment. Otherwise, failures may occur, reducing die life and increasing tooling costs. The following sections describe some typical failures of tool steels after heat treatment and the main mechanism that caused failure. The main mechanism is described and divided systematically, but in practical situations, several mechanisms as well as the tool use should be

(a)

(b)

Fig. 23

(a) Tool called a triturating wheel made from AISI D6 (similar to D3) that had poor performance due to a premature crack. Tool diameter of approximately 300 mm (~11.5 in.). The identified cause was the absence of tempering. The hardness measurement was 65 HRC, the usual value for the as-quenched condition of this grade. Courtesy of Villares Metals. (b) D2 dies cracked during finish grinding operation. Cracks emerged due to the as-quenched condition of the microstructure (not tempered). Source: Ref 9

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tool steel. In Fig. 24, this can be observed for high tempering temperatures, but the variation is even higher if the whole tempering curve is observed (see the tempering curve for H13, Fig. 28, for example). Normally, it is possible to attain values between 30 and 58 HRC in the most common hot work steels. However, the steels are not used within this hardness range. Although some variation may exist, hardness levels higher than 50 HRC or lower than 40 HRC are not typical. To improve wear resistance, one can propose an increase in hardness. However, two aspects should be considered before such a decision. First and most obvious is the toughness necessary for a given application. Reduction of

investigated to discover the root cause of a failure. In specialized literature, it is less common to discuss failures of hot work tools than of cold work. This is because of the intrinsic lower brittleness of hot work tool steels. Nevertheless, the subject is very important for these materials, due to the high value of a tooling set for hot working and also because of the large production performance expected from it. Even if no catastrophic failure occurs, the reduction of hot work tool performance can cause serious damage in terms of cost and time. Premature Cracking Caused by Excessive Hardness. A wide range of hardness levels can result from hardening and tempering a hot work

Fig. 24

Tempering curves for the most common hot work tool steels. Tempering curves are obtained after hardening small (25 mm, or 1 in.) specimens of all materials with the usual hardening temperature: 1020 C for H13, TENAX300 (brand name of lowsilicon H11), and VHSUPER (brand name of high-molybdenum, low-silicon modified H11); and 1100 C for H21 (higher temperatures may be used, increasing tempered hardness but reducing toughness).

Table 3 Typical chemical composition of some hot work tool steels Composition, wt% AISI(a)

~H10 H11 H13 Low-Si H11; TENAX300(b) ~6F3 or ~L6 H21

DIN

UNS

C

Si

Mn

Cr

W

V

Fe

Others

1.2365 1.2343 1.2344

T20810 T20811 T20813 ...

0.32 0.37 0.40 0.36

0.30 1.0 1.0 0.3

0.3 0.3 0.3 0.3

2.9 5.0 5.0 5.0

Mo

2.8 1.3 1.3 1.4

... ... ... ...

0.50 0.90 0.45 0.45

bal bal bal bal

... ... ... ...

1.2714 1.2581

... T20821

0.56 0.32

0.3 0.3

0.7 0.3

1.1 3.5

0.5 ...

... 9.0

0.15 0.50

bal bal

Ni = 1.7 ...

(a)“~”, similar to but not exactly the same. (b) Trademark of Villares Metal Company, Brazil. TENAX 300 is not standardized; therefore, the brand name is given.


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Fig. 25

Examples of heat checking cracks on aluminum die-casting dies. Cracks are white because they are filled with aluminum. Courtesy of Villares Metals

Fig. 26

Schematic of hardness after tempering and the effect of secondary hardening in high-alloy steels. Observe that high-temperature hardness is only possible through precipitation hardening caused by alloy carbides (secondary hardening). At low temperatures, hardness is less than martensite due to the presence of retained austenite, which is eliminated after high-temperature tempering.

toughness may accelerate several mechanisms of tool damage, such as heat checking, gross cracking, and wear by chipping, in addition to the risk of catastrophic cracking. The relationship between hardness, toughness, and tempering temperature can be quantitatively evaluated in Fig. 28 for AISI H13. When the tempering temperature is far from the hardness peak, toughness rapidly increases with decreasing hardness. Temperatures of approximately 500 C can be considered to cause temper embrittlement, with toughness being rather low (Ref 24–26). Such temper embrittlement has been studied in relation to silicon content (Ref 27–30). It has been shown, for example, that reducing the silicon content from 1% (usual in H-series grades) to approximately 0.3% causes a strong toughness increase, close to peak hardness but also for higher tempering temperatures. Secondly, the type of wear should be considered. In cold work tooling, wear is directly related to hardness, but in hot work tooling, the

situation is normally quite different. Wear can occur by a combination of abrasion and adhesion at high temperatures. The main reason to avoid wear is the ability to keep high hardness at working conditions (tempering resistance) as well as adequate toughness to avoid chipping after adhesion. One example of such an effect is presented in Fig. 29 for a precision hot forging punch. The traditional material for such an application was H13 steel with high hardness (55 HRC). However, accelerated wear was observed. Analysis of the tool after end-life showed a strong hardness decrease in the working regions, indicating that a higher tempering resistance was necessary instead of higher hardness. Such an alteration was done, substituting H13 with a higher-molybdenum grade, the commercial brand VHSUPER (not standardized), which led to a 50% increase in the tool performance. It is commonly found that excessive hardness assists failures in hot work applications. Figure 30 shows further examples. For all, low tempering temperature, short tempering time, or even specification changes led to an increase in hardness and premature failure due to the resulting low toughness. While excessively low hardness can lead to problems, hardness also should not be too low. The adequate value, as mentioned previously, is normally between 40 and 50 HRC in hot work tools; in some cases for warm forging, it is possible to use tools up to 58 HRC. An interesting effect of hardness and toughness in final performance can be discussed based on Fig. 31. This graph shows the effect of these two properties on heat checking, focused on die-casting applications. The higher the hardness and toughness, the lower the heat checking damage, measured by an appropriate chart. However, as discussed and shown in Fig. 28, hardness and toughness are also associated. Establishing ideal

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hardness thus depends on several conditions of tool design and use and is also based on previous experience. As a starting value, 45 HRC would be recommended, with increases or decreases depending on the results experienced. Inadequate Heat Treating Procedures in Hot Work Tool Steels. Heat treating of hot

Fig. 27

work steels has several important parameters, but incorrect practice does not always produce failures observable just after heat treating. Some examples were shown in Fig. 29 and 30; several tools were improperly heat treated (to higher hardness), but the problem was only observed during tool use.

Effect of vanadium molybdenum, tungsten, and chromium additions on secondary (high-temperature) hardness of mediumcarbon steels. Source: Ref 23

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This is common for hot work tools due to the low as-quenched hardness and the high hardenability of this class of tool steels. For example, the use of improper temperatures may lead to the specified hardness, which can even be homogeneous, but the mechanical properties, such as toughness or tempering resistance, can be deeply affected. A poor practice during quenching can lead to high brittleness with no observed changes in hardness or strength. These two aspects are described in this section. Before continuing, a useful recommendation should be made. Due to the high value of diecasting tools and their high productivity, a strong effort has been made for improvements in this field. An important recommendation for H13 tool steel was written by the North American Die Casting Association (NADCA) (Ref 32). This recommendation provides important

Fig. 29

Fig. 28

Impact toughness and hardness as a function of tempering temperature. Retained austenite content is also shown. Notice the hash-marked area, indicated as a temper embrittlement region, where very low toughness is observed; this region coincides with the peak hardness. Source: Ref 24

Example of the importance of tempering resistance instead of initial working hardness. (a) Hot forging punch showing wear and cracks as the normal failure condition. For maximum wear resistance, initial hardness was established at 56 HRC for H13 tool steel. However, the end-life mechanism was related to hardness reduction in the working (heated) areas, as shown in (b). This grade was substituted by a higher-molybdenum grade, brand name VHSUPER, with higher tempering resistance, as shown in (c) by the longer times necessary for hardness decrease. The substitution lead to 50% longer tool life. Courtesy of Villares Metals

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information for quality assurance of tool steel and also for applied heat treating. It is very helpful for die-casting applications as well as for quality analysis of AISI H-series steels applied to other processes. Incorrect Hardening and/or Tempering Temperatures. To analyze the effect of incorrect temperatures on hot work steel properties, several studies were conducted in the laboratory (Ref 12). Toughness was evaluated by unnotched impact testing, according to the NADCA procedure (Ref 32). Three conditions were simulated. To show the effect of low hardening temperature on tempering resistance, conditions 1 and 2 were evaluated:

Condition 1: standard condition for H13 grade, with hardening temperature at 1020 C and double tempering at 610 C (for 2 h at temperature), leading to 45 HRC Condition 2: Low hardening temperature (at 890 C), simulating a condition for furnaces

Fig. 30

that reach up to 900 C. For this condition, tempering should be reduced to 250 C. Condition 3: higher hardening temperature, increasing from 1020 C to 1150 C. Such temperature is currently used for high-speed steels heat treated to lower hardness. For 45 HRC, tempering was slightly increased, to 640 C, also twice for 2 h.

Figure 32 evaluates the reduction in toughness promoted by inadequate heat treating conditions, as well as the respective microstructures. The low-tempering situation, condition 2, causes a substantial reduction in toughness (40%) as well as a loss in tempering resistance (the hardness decrease was six times higher than expected). Toughness reduction in this condition, is caused by incomplete austenite transformation, causing a heterogeneous microstructure (Fig. 32c), as well as by the low tempering temperature, which does not promote adequate martensite stress relief. On the other hand, the decrease in

Examples of failures caused by excessive hardness. (a) Tool made of DIN 1.2714 tool steel (similar composition to AISI 6F3 and L6) that fractured after a short life. For this tool, the hardness was expected to be approximately 40 HRC, but a sample was analyzed and found to be 50 HRC. Arrows indicate cracking location and cracking initiation site. (b) Microstructure showing light areas, indicating excess retained austenite and untempered martensite, another indication that low tempering temperature was employed and/ or only one tempering treatment. This led to high hardness as well as a brittle microstructure. (c) H13 forging die that cracked prematurely (arrows). Hardness was measured at 52 HRC but expected values were approximately 44 HRC. The excessive hardness was caused by short tempering times and low temperatures. (d) Typical H13 microstructure tempered at low temperatures. Etched with 4% nital. It is lighter than usual (compare to Fig. 32b) due to low-temperature tempering, which causes poor precipitation of alloy carbides and thus enhances corrosion resistance during etching. Courtesy of Villares Metals

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tempering resistance is caused by the insufficient dissolution and reprecipitation of alloy carbides (poor secondary hardening), which is the main strengthening mechanism in hot work tool steels at high temperatures. The most intense embrittlement was produced by condition 3—too high hardening temperature. In this situation, very coarse austenite grains are produced (Fig. 32d), leading to increased grain-boundary embrittlement. When high hardening temperatures are used, precipitation of proeutectoid carbides on grain boundaries intensifies, causing a great reduction in toughness, as observed (90% lower). Some sources studied the use of higher hardening temperature as one way to improve thermal fatigue (Ref 22, 33). In fact, the increase in hardening temperature leads to more dissolution of alloy carbides, rich in vanadium and molybdenum, which increases the content of alloy elements in solid solution and enhances secondary hardening. A simple observation of this is shown in condition 3, where tempering should be increased 30 C to attain the same hardness as the usual hardening condition. In terms of tempering resistance, 30 C (54 F)

Fig. 31

the dislocation in tempering curve indicates a substantial increase in tempering resistance, because temperature effect is exponential to time effect in tempering conditions. This phenomenon also explains some advantages found in specimens austenitized at 1100 C compared to 1020 C (Ref 22). However, modifications in hardening temperatures are rarely possible in practical (industrial) conditions. Increasing the hardening temperature deeply affects the precipitation behavior on grain boundaries during quenching, causing intense embrittlement. Figure 33 shows this effect (note the dashed lines), but a full explanation is given in the following section (especially regarding Fig. 36). Before continuing, it is interesting to show a case of failure caused by excessive hardening temperature. Figure 34 presents such a case—a tool that cracked after low performance. It is easy to see the grain-boundary crack propagation, caused by the coarse grain size as well as by precipitation of carbides on grain boundaries. This is denoted by the preferential and strong etching of austenite grain boundaries. Slow Cooling during Quenching. Reaching the final hardness in tool steels is quite a

Heat checking resistance (lower readings indicate higher resistance) as a function of unnotched impact toughness and hardness of H13 steel. Heat checking is evaluated by the photographs on the left; the rating is calculated by adding the column representing the largest cracks (leading) and the column representing the severity of the cracks (network). See text for discussion of these results. Source: Ref 31

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simple task for quenching. Although there are some exceptions, tool steels normally have high hardenability, with the as-quenched hardness attained even if improper procedures are applied. However, hardness, although necessary, is not sufficient for the high required performance of tool steels, as already shown. In this context, quenching should be considered as a process for promoting the required mechanical properties, not only for attaining a specified hardness. For this reason, consider Fig. 33 again. For both hardening temperatures, hardness higher than 500 HV (~49 HRC) is obtained within a wide range of the continuous cooling tranformation (CCT) curve. Considering the hardness reduction after tempering, it is therefore quite simple to obtain the final hardness with different

(b) condition 1

Fig. 32

(c) condition 2

quenching practices, even using air quenching. However, two further important aspects should be considered. First, the dashed lines in the CCT diagram (Fig. 33) indicate formation of proeutectic carbides. As mentioned in the previous section, if austenitizing temperature increases, more alloy elements go into solid solution by carbide dissolution. In cooling, the process is reversed, and such carbides tend to form again. This happens by precipitation in high-energy areas, the most important being the grain boundaries. The result is a film of carbides between grains, which weakens the interface and promotes failure (Ref 35). Such a phenomenon is marked by two characteristics: a strong etching at austenite grain boundaries (because the interfaces of carbides and steel are regions

(d) condition 3

Laboratory simulation of adequate and inadequate heat treating conditions for AISI H13. The first situation (condition 1) is the recommended heat treatment: hardening at 1020 C, followed by two tempering treatments at high temperature. In this case, 45 HRC was desired, and thus tempering was performed at 610 C. Condition 2 involves a very low hardening temperature, where austenitizing was done at 890 C. To reach 45 HRC, specimens were heat tempered at 250 C. In addition to toughness reduction, the heat treating condition caused reduction of tempering resistance. Condition 3 describes a situation with excessively high hardening temperature (1150 C), with tempering done at 640 C to attain 45 HRC. In (a), the impact toughness is presented, and in (b) to (d), the microstructure relative to each condition is shown (same magnification; etched with 4% nital). Source: Ref 12

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more prone to corrosion) and, in stronger cases, intergranular failure. As shown in Fig. 33, an increase in austenitizing temperature causes dislocation of the dashed lines to the left, indicating stronger precipitation, even when high cooling rates are applied. Secondly, a slow

cooling also affects the previous microstructure, forming bainite instead of martensite. Although the precipitation necessary for secondary hardening is practically not affected (Ref 33), modification of the initial microstructure, from martensite to bainite, also reduces toughness

Continuous cooling transformation diagrams for H13 tool steel austenitized at 1030 C (1885 F) (top) and 1100 C (2010 F) (bottom). Note the dislocation of the dashed line, indicating more pronounced proeutectic carbide precipitation on grain boundaries for the high austenitizing temperature condition. Source: Ref 34

Fig. 33

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(ref 29, 33 and 36). The effect of decreasing the cooling rate is thus embrittlement for both mechanisms.

Fig. 34

Die failure caused by excessive hardening temperature. Two tools were analyzed: one made of VHSUPER steel (commercial brand name) and the other of AISI H13. (a) One of the tools cracked in the position denoted by the arrow, where a sample was cut for analysis. (b) Typical microstructure from H13 tool and (c) from VHSUPER tool with 54 HRC (45 HRC was expected). Note the coarse grain size, approximately ASTM 3 to 4. For these grades, grain size is expected to be approximately ASTM 7 to 10. Courtesy of Villares Metals

Therefore, both bainite and carbide precipitation on grain boundaries must be avoided by preventing slow cooling in quenching. This is important advice for tools sensitive to failures caused or assisted by cracks. One example is presented in Fig. 35 for hot forging dies that failed after short service time. The microstructural analysis showed coarse grain sizes and strong precipitation on grain boundaries (Fig. 35b, c), illustrating the interaction between the two effects. As a final result, strong embrittlement occurs (Fig. 35d) as well as a clearly intergranular fracture (Fig. 35e). Typically, hot work tool steels were oil quenched, but today (2008), vacuum heat treating with pressurized nitrogen quenching has become very popular. In this treatment, cooling rate control is rather critical, since it is related to nitrogen pressure and gas circulation as well. If too strong and heterogeneously applied, cooling may lead to strong distortion or even quenching cracks. On the other hand, grain boundary embrittlement occurs easily if the cooling rate is too slow. A guide for evaluating tool heat treating quality is described by the NADCA (Ref 32), including the use of coupons for destructive testing after heat treating as well as advice for vacuum hardening. Another important issue is the step in which heat treatment should be applied. With the advance in machining technology, the feasibility for machining in higher hardness has increased; machining hot work dies up to 50 HRC is rather common by means of high-speed machining technologies (high cutting speeds with low feed). Consequently, it is common, mainly in forging dies, to machine from prehardened blocks. However, the probability of embrittlement increases as the section size of the tool increases. Figure 36 shows the effect of section size and austenitizing temperature on the toughness of H13 steel. The tendency for toughness loss is evident when larger sizes or higher austenitizing temperatures are used, because they are directly related to the grain-boundary embrittlement effect and are also affected by bainite formation. Even if the quenching process uses a strong cooling medium, large tools are unavoidably sensitive to embrittlement in core regions. Therefore, heavy-section tools with deep engravings should be heat treated only after rough machining to avoid embrittlement of tool working regions. The most important example in this field is die-casting dies. They are usually heat treated only after machining to improve

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(a)

Fig. 35

Example of die failure in a hot forging die caused by coarse grain size and strong precipitation of proeutectoid carbides on austenite grain boundaries. (a) Aspect of the tool. (b) and (c) Microstructure showing the coarse grain size (approximately ASTM 4; expecte d ASTM 8 to 10), marked by preferred etching on carbides present at grain boundaries and the coarse martensite laths. Samples were taken from the tool midradius and analyzed regarding (d) impact toughness in the as-received condition and after new heat treating to the same hardness and (e) fracture of impact-tested specimens (for the initial condition—as-received) by scanning electron microscopy. Note the strong increase in toughness after new heat treating, indicating the deleterious effect of carbide precipitation on grain boundaries, producing intergranular failure in impact specimens. Courtesy of Villares Metals

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toughness, because die resistance to heat checking is directly related to this property (Fig. 33). Even if the quench is applied after machining, it is still important to control the Quench cooling rate (avoiding too slow conditions), due to the possibility of embrittlement of surface regions, which are the working areas of die-casting dies (Ref 32). Failures of Nitrided Tools. Nitriding is commonly used for several tools and dies. For example, hot extrusion dies are typically nitrided for all uses. They normally work for production of parts in aluminum or other nonferrous alloys, mainly for construction applications. Different from other hot working processes, extrusion is continuous and involves constant flow between the conformed alloy, and the tool steel. This enhances the wear condition, which is the usual end-life mechanism. Nitriding considerably improves surface hardness and, consequently, wear resistance, this being the reason for its application in virtually all extrusion dies. The same approach is applied in hot forging tools.

Nitriding is currently used in several tools and dies to improve wear resistance. However, as hardness increases, surface toughness decreases, and for deep tools, especially those prone to cracks, nitriding is prevented. In die casting, nitriding is used in some cases, claiming that the increase in surface hardness tends to increase the initiation of heat checking cracks; however, as cracks cross the nitriding layer, it has no effect at all. Depending on the nitriding layer condition, it can be harmful to crack initiation. Normally, the nitride layer in tool steels is 0.1 to 0.3 mm thick. Hardness is higher than in carbon steels, due to the formation of alloy nitrides. In hot work steels from the H series, chromium is very important for this effect, and the maximum surface hardness approaches 1100 HV. Typical nitrided microstructures present a diffusion layer and a fine white layer on the tool surface (Fig. 37a). In H-series tool steels, the

Fig. 36

Charpy V-notch (CVN) results for different heat treating conditions of H13 tool steel, carried out at room temperature and at 425 C. Specimens cooled at various rates, simulating the core of 150 and 300 mm round bars, as well as an air-cooled 25 mm specimen. Results were tested for different austenitizing temperatures. Toughness reduction is evident at higher austenitizing temperatures and larger sizes, both related to lower cooling rates during quenching. Source: Ref 33

Fig. 37

Tool steel surface after nitriding. (a) White and diffusion layers (b) Coarse nitrides precipitated on grain boundaries. See text for discussion. Courtesy of Villares Metals

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nitriding layer is darker after etching, due to the depletion of chromium content from the steel matrix, thus causing a reduction in corrosion resistance. In carbon or low-alloy steels, the nitride layer appears lighter in the microstructure and can increase corrosion resistance. In hot work tool steels, it is rather common to observe coarse nitrides precipitated on grain boundaries inside the diffusion layer, as shown in Fig. 37(b). Depending on the intensity, this embrittles the tool and forms an exceptional route for crack propagation (Fig. 42). Precipitation of coarse nitrides on grain boundaries can also be accompanied by a thick white layer on the surface. This layer is extremely hard and brittle, because it is uniquely composed of nitrides. It is therefore a common region for crack initiation. The combination of a coarse white layer and grain boundary precipitation of

Fig. 38

coarse nitrides commonly leads to premature crack initiation, with damage to several tools, especially forging tools prone to cracking or for die-casting dies. Some examples of failures assisted by this phenomenon are shown in Fig. 38. It is thus recommended, for these applications and in general, that the nitriding process be controlled in order to avoid both a thick white layer and grain-boundary nitrides. Today (2008), this control is usually performed in computer-controlled gas nitriding and plasma nitriding processes. Excessive Heating Causing Tool Failure. As discussed previously, hardening of hot work tool steel mainly results from precipitation of alloy carbides during the tempering treatment. This phenomenon is also known as the fourth stage of tempering and occurs after the modification of martensite and the formation

Examples of undesirable microstructures encountered on the surface of nitrided tools. For both cases, the core microstructures are correct, indicating proper hardening and tempering procedures. (a) Surface and (b) core microstructure of a nitrided forging tool, showing (in a) the problems of a coarse white layer and nitrides on grain boundaries. (c) and (d) Extracted from a die-casting die failure analysis, also for surface and core respectively. Note the strong precipitation on grain boundaries in (c), whereas core regions are quite well heat treated (to approximately 44 HRC), leading to 300 J of unnotched (NADCA) impact strength. Nevertheless, an unexpected failure occurred (after less than 100 shots), caused primarily by improper tool use but also assisted by the nitriding layer condition.

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of cementite. At low temperatures, typically below 500 C, the thermodynamic driving force (free energy) is insufficient for such carbide precipitation (mainly molybdenum or vanadium carbides), and no hardness peak is observed (Fig. 27). As the temperature increases, carbide formation starts and very small carbides (a few nanometers) precipitate within martensite laths. Because they are fine and large in number, these carbides are proper for obstructing dislocation movement, causing substantial improvement in strength. This occurs only at high temperatures, with precipitation of alloy carbides pertaining to materials used in hot working, in this case, hot work tool steels. For higher temperatures, carbides tend to transform to more stable types or to coarsen, which increases the size and thus reduces the total number of precipitated carbides. This results in loss of the dislocation blocking effect, leading to strength reduction. The total effect is schematically shown in Fig. 26. Temperature and time affect the precipitation behavior. From a thermodynamic point of view, the lowest energy is obtained after alloy carbide precipitation (reducing the free energy of elements in solid solution) and if their sizes are as large as possible (reducing the surface free energy). Consequently, the higher the time or temperature, the easier it is for carbides to become larger, and strength tends to decrease. In tempering curves, such as those shown in Fig. 24 and 28, this can be observed for a fixed time (twice for 2 h) according to temperature. However, a complete view is given by using the tempering parameter, as shown in Fig. 39. This combines the effects of both time and temperature in only one variable, the parameter, proportional to time and to a logarithm of the temperature. Through this mathematical relation, it can be seen that temperature actually has the highest effect; however, if excessively long times are used at appreciably high temperatures, the same effect may occur. For example, 4 h (twice 2 h tempering) at 600 C (1110 F, parameter = 32,700) leads to ~46 HRC for H11, which is equivalent to 15 h at 577 C (1070 F), or 3 days at 550 C (1020 F), or approximately 1 month at 500 C (930 F). That is why below 500 C, heating has practically no effect. However, a second situation should also be considered. Low times at excessively high temperatures can produce rather important effects. This is the situation for several tool applications, where tools are in contact at very high temperatures but for very short times. For

example, consider a forging tool that reaches 750 C (~1380 F) for a half-second during each stroke. After producing 1000 parts, this tool will be exposed to 750 C for a total time of 8.3 min. Such heating is equivalent to a tempering parameter of 35,260 or 2 h at 691 C (1275 F). In H11, such heating is capable of reducing the hardness from 46 to less than 30 HRC. It is in this context that the failures described in this section should be considered. In some situations, tools are correctly heat treated, but, during hot working operation, heating is so high that strength is reduced and failure is accelerated. One example, shown in Fig. 40, is a hot forging punch that usually exhibits low life if compared to the whole set of tools. It is made of DIN 1.2885, a highly temper-resistant hot work steel that has high amounts of molybdenum and cobalt for improvement of this property. Even though the analysis showed a continuous reduction of hardness approaching the punch tip, which contacted the hot inforging part. Besides the hardness values, it is interesting to note the microstructural behavior. Regions far from the hot working areas are typically tempered martensite, becoming darker closer to the punch tip. As previously explained, secondary carbides are not visible

Fig. 39

Tempering curve as a function of the time-temperature parameter for H11 steel containing 0.40% C, 0.92% Si, 5.09% Cr, 1.34% Mo and 0.52% V. For this curve, t = time in hours, and T = temperature in F+460. Source: Ref 37

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under optical microscope; however, their effect on corrosion resistance (and thus on etching behavior) is evident. High tempering temperatures cause an increase in precipitation and carbides, leading to reduction of corrosion resistance due to the depletion of alloy elements in the steel matrix and to the increased sites (interfaces) for corrosion. This explains why the microstructure becomes darker where heating is more intense and hardness is lower (this will always happen if the hardening practice is not changed). In the punch in Fig. 40, heating is so high that secondary carbides have become large enough to be observable by optical microscope. This is typically the annealed state for this alloy, not obtained through transformation annealing but by temper annealing. The effect shown in Fig. 40 occurs in several hot working tools. A further example is shown in Fig. 41 and 42 of a forging die for automotive valves. After some time in production, wear of the forging die produced grooves on the surface of forged products (Fig. 41a), leading to die substitution. The analysis revealed that wear

Fig. 40

was deeply influenced by thermal fatigue cracking (Fig. 41c, d). By comparing heated and not heated areas, two further problems were observed. First, some regions on the die surface had been heated to high temperatures, leading to reaustenitizing, rehardening, and the formation of untempered martensite (Fig. 42a). Second, the tool surface had serious indications of nitriding problems, with coarse nitride formed on grain boundaries and crack propagations through these regions (Fig. 42c). The combination of intense surface heating and nitriding led to a “crazy pattern” of hardness, differing substantially from that encountered in non-heat-affected areas (Fig. 42b). This was caused by the combination of rehardening and nitrogen diffusion to the core regions. The end result, of course, was a deep embrittlement of the die, as observed by its intense cracking. In cases similar to those in Fig. 40 to 42, two possibilities are possible to solve or reduce the problem. The first is to employ one steel grade with a higher tempering resistance, normally a higher-alloyed grade; however, this approach has some limitations. If the situation is only an

Microstructures of hot forging punch made of DIN 1.2885. Etched with 4% nital for the same time in all microstructures

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increase in tool resistance to softening, changing the steel grade may be a solution. For example, this was the situation described in Fig. 29. However, it should be noted that highly alloyed grades can also be more brittle, leading to problems if gross cracking or heat checking are important issues for tool failure. Also, the increased alloy is directly related to steel cost, with an analysis of the cost and benefits of steel substitution being important. The best way is to consider the total tool cost for each produced part, which includes the increase in steel cost but also its higher performance as well as the reduction of setup times for tool substitution.

Fig. 41

However, in some situations, heating is so intense during tool application that it is impossible to solve the problem only by modifying the tool steel employed; thus, there is a second possibility concerning cooling applied during hot work tooling. Coolants and lubricants are both important in hot work tooling, in many cases being the same product. In cases where appreciable regions of tools are reaching more than 700 C or even reaustenitizing, the best practice is to work on the lubricating/cooling involved in the process, before considering changes to the tool steel used. It would be helpful to contact the lubricant manufacturer and obtain updated

(a) Forging tools for production of automotive engine valves; the analyzed die was painted in blue. (b) Forged valve and, in detail, the grooves produced when using worn dies. (c) Three analyzed regions, marked “A”, “B” and “C”. “A” is the main wear region, but “B” is where a large amount of thermal fatigue cracking is observed. “C” is a region not affected by process heat and used as a reference. (d) Thermal cracks of region “B”, under scanning electron microscopy. Courtesy of Villares Metals

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(b)

Fig. 42

(a) Microstructure of working regions of valve forging die in Fig. 41 (region A). Note the intense cracking and thick nitrided layer (double the expected). Also note the surface white layer, which corresponds to brittle untempered martensite, obtained by rehardening of the tool surface. (b) Hardness profile of this region (A) and another region not affected by heat (C) in Fig. 41(c). (c) Two pictures of region C, not affected by heat, that have a mechanical crack within the nitriding layer. Note the presence of a coarse white layer and nitrides on the grain boundaries. The micrographs in (c) refer to the same region, but the one on the right is slightly underfocus for better observation of nitrides on the grain boundaries. Courtesy of Villares Metals

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information on the possibilities for a specific application. Besides the type of lubricant, a very important issue is the way the lubricant is applied. In modern processes, robots may be used for homogeneous application of lubricants. It is also important to evaluate the possible changes in tool design to avoid heat during hot working. The answer is complex and should consider tool life, process productivity, and the necessary investments, always aiming for the minimum cost per produced part.

4.

5. 6.

Conclusion All theory and data discussed in this chapter may be summarized by the following points:

Heat treating of tools and dies may be considered crucial for their performance. Several examples show that success cases as well as premature failures are often related to heat treating quality. Hardness measurements alone are usually not the best indication of proper quality in too steel heat treating. Other properties such as strength, toughness, and wear resistance, as well as the microstructural features, are better indicators as to whether a heat treatment was done correctly or not. Assessing these properties is usually difficult in tools and dies because it would be a destructive (and expensive) test. Therefore, guaranteeing the correct procedure is the best way to assure that the heat treatment of a tool was performed correctly and thus enables adequate tool performance. In order to apply such correct conditions, the best approach is to follow the specifications for each grade which is normally provided by the tool manufacturer. For several specific points such as design, tool use, and surface treatments, the examples given with this chapter may be helpful.

7. 8. 9. 10.

11.

12.

13. 14.

15. REFERENCES

1. G. Roberts, G. Krauss, and R. Kennedy, Tool Steels., 5th ed. ASM International, 1998 2. Tool Steels, Steel Products Manual, Iron and Steel Society, April. 1988 3. R.A. Mesquita and C.A. Barbosa, Avaliaçaõ das Propriedades de Desgaste e Tenacidade em Aços para trabalho a Frio (Evaluation of Wear and Toughness Properties of Cold Working Tool Steels). Tecnol. Metal.

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Mater., Vol 3 (No. 2), Oct/Dec 2005, p 12–18 (Abstract in English; Article in Portuguese) C.J. Altstetter, M. Cohen, and L. Averbach, Effect of Silicon on the Tempering of AISI 43XX Steels, Trans. ASM, Vol 55, 1962, p 287 G. Dieter, Mechanical Metallurgy, 3rd ed., McGraw-Hill Book Co., 1986 R.M. Hemphill and D.E. Wert, Impact and Fracture Toughness Testing of Common Grades of Tool Steels, Tool Materials for Molds and Dies, G. Krauss and H. Nordberg, Ed., Colorado School of Mines Press, 1987, p 66–69 Failure Analysis and Prevention, Vol 10, Metals Handbook, 8th ed., American Society for Metals, 1975, p 500–507 Met. Eng. Q., Feb 1973, p 31–41 J.R. Davis, Ed., Tool Materials, ASM Specialty Handbook, ASM International, 1995 “Heat Treatment of Tool Steel,” Technical brochure, Uddeholm Tooling, http://www. uddeholm.com/files/heattreatment-english.pdf (accessed Dec 2006) “Grinding of Tool Steel,” Technical brochure, Uddeholm Tooling, http://www. uddeholm.com/files/grinding-english.pdf (accessed Dec 2006) R.A. Mesquita, D.R. Leiva and C.A. Barbosa, Estudos de Tratamento Te´rmico nos Aços Ferramenta VH13ISO E VF800AT (Heat Treating Studies of VH13 and VF800AT Tool Steels), Proceedings of Third Encontro da Cadeia de Ferramentas, Moldes e Matrizes, ABM, 2005, p 30–40, (Abstract in English; article in Portuguese) “Cold Work Tool Steel, VF800-AT ” Datasheet, Villares Metals Company, http://www. villaresmetals.com.br/ (accessed Dec 2006) K.W. Andrews, Empirical Formulae for the Calculation of Some Transformation Temperatures, J. Iron Steel Inst., Vol 203, 1965, p 721–727 W.J. Harris, Jr. and M. Cohen, Stabilization of the Austenitic-Martensite Transformation. Trans AIME, Vol 80, 1949, pp. 447– 470. B.L. Averbach, S.A. Kulin, and M. Cohen, The Effect of Plastic Deformation on Solid Reactions, Part II: The Effect of Applied Stress on the Martensite Reaction, Cold Working of Metals, American Society for Metals, 1949 Heat Treating, Vol 4, ASM Handbook, ASM International, 1991

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18. “Cold Work Tool Steels”, Datasheets, Uddeholm Tooling, Hagfors, Sweden 19. A. Mendanha, H. Goldenstein and C.E. Pinedo, The Role of Microstructure on the Toughness Behaviour of AISI D2 Cold Work Tool Steel, Proceedings of the Seventh International Tooling Conference—Tooling Materials and Their Applications from Research to Market—‘Tool 06, Vol 2, (Politecnico de Torino, Torino, Italy), 2006, p 813–819 20. M.A.H. Howes, Heat Checking in Die Casting Dies, Die Cast. Eng., March-April 1969, p 12–16 21. T.C. Benedyk, D.J. Moracz, and J.F. Wallace, Thermal Fatigue Behaviour of Die Materials for Aluminum Die Casting, Paper 111, Proceedings of Sixth SDCE International Die Casting Congress, 1970 (Cleveland, OH), The Society of Die Casting Engineers, Inc., 1970, p 1–20 22. J. Sjo¨stro¨m and J. Bergstro¨m, Thermal Fatigue Testing of Chromium Martensitic Hot-Work Tool Steel after Different Austenitizing Treatments. J. Mater. Process. Technol, Vol 153–154, 2004, p 1089– 1096 23. W. Crafts and J.L. Lamont, Secondary Hardening of Tempered Martensitic Alloy Steel, Vol 180, TMS-AIME, 1949, p 741 24. “Hot Work Tool Steels”, Datasheets, Uddeholm Tooling, Hagfors, Sweden 25. J.R.T. Branco and G. Krauss, Toughness of H11/H13 Hot Work Tool Steel, New Materials Processes Experiences for Tooling. H. Berns, M. Hofmann, L.-A. Norstro¨m; K. Rasche, and A.-M Schiner, Ed., Materials Search (Interlaken, Switzerland), 1992, p 121–134 26. H. Berns, Strength and Toughness of Hot Working Tool Steels, G. Krauss and H. Nordberg, Ed. Tool Materials for Molds and Dies: Application and Performance, (Ilinois), The Colorado School of Mines Press, 1987, p 45–65 27. W.M. Garrison, Jr., Influence of Silicon on Strength and Toughness of 5wt-%Cr Secondary Hardening Steel, Mater. Sci. Technol., Vol 3, April 1987, p 256–259 28. D. Delagnes, P. Lamesle, M.H. Mathon, N. Mebarki and C. Levaillant, Influence of Silicon Content on the Precipitation of Secondary Carbides and Fatigue Properties of a 5% Cr Tempered Martensitic Steel,

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Mater. Sci. Eng. A, Vol 394, 2005, p 435– 444 M. Umino, T. Sera, K. Kondo, Y. Okada and H. Tubakino, Effect of Silicon Content on Tempered Hardness, High Temperature Strength and Toughness of Hot Working Tool Steels, Tetsu-to-Hagane (J. Iron Steel Inst. Jpn.), Vol 89, (No. 6), June 2003, p 673–679 R.A. Mesquita and C.A. Barbosa, Effect of Silicon and Phosphorous on the Toughness of G11 Hot Work Tool Steel, Proceedings of 61st Brazilian Metallurgical and Materials Science Congress, ABM, 2006 (in Portuguese) L. Eliasson and O. Sandberg, Effect of Different Parameter on Heat-Checking Properties of Hot-Work Tool Steels, New Materials and Processes for Tooling, H. Berns, H. Nordberg and H.-J. Fleischer, Ed. (Bochum, Germany), Verlag Schurmann and Klagges KG, 1989, p 1–7 “Premium Quality H13 Steel Acceptance Criteria for Pressure Die Casting”, 207/ 2003, North American Die Casting Association (NADCA) M.L. Schmidt, Effect of Austenitizing Temperature on Laboratory Treated and Large Section Sizes of H-13 Tool Steel, Tool Materials for Molds and Dies., G. Krauss and H. Nordberg, Ed. (Illinois) Colorado School of Mines Press Center, 1987, p 118–164 K.-E. Thelming, Steel and Its Heat Treating, 2nd ed., Butterworths, London, 1984 C.L. Briant and S.K. Banerji, Intergranular Failure in Steel: The Role of Grain Boundary Composition, Int. Met. Rev., (No. 4), 1978, p 164–199 T. Okuno, Effect of Microstructure on the Toughness of Hot Work Tool Steels, AISI H13, H10 and H19, ISIJ, Vol 27, (No. 1), 1987 J.C. Hamaker, Jr., Die Steel Useful for Ultra High-Strength Structural Requirements, Met. Prog., Dec 1956, p 93 R.A. Mesquita, C.S. Gonçalves, and C.A. Barbosa, Effect of Hardening Conditions on the Mechanical Properties of High Speed Steels, Proceedings of the European Conference on Heat Treatment 2008 — Innovation in Heat Treatment for Industrial Competitiveness, Verona, Italy, 2008, CDROM

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Case Studies of Steel Component Failures in Aerospace Applications Scott MacKenzie, Houghton International, Inc.

STEEL has limited use in modern aircraft, but it is used in landing gear systems, arrestment systems for naval aircraft, and for support equipment. Because of the high strength levels required (1790 to 2070 MPa, or 260 to 300 ksi, yield) to maintain justification for the use of steel, these very high-strength steels are prone to environmental effects (stress-corrosion cracking and hydrogen embrittlement), flaws created during manufacturing (laps, seams, machining gouges, and grinding), and heat treatment (distortion, decarburization, and quench cracking). Care is needed to protect and inspect these highvalue and safety-critical steel components. The following case histories illustrate typical failures experienced by these high-strength steels. The case histories in this chapter illustrate a variety of failure mechanisms. The causative reasons vary from manufacturing to operational to environmental. Because of the many different types of root causes, the failure engineer must be aware of the manufacturing process, the assembly process, as well as the environment to which component is exposed to effectively determine the primary cause of failure.

Fig. 1

Failure Analysis of a Catapult Holdback Bar This investigation analyzed the failure of a repeatable-release holdback bar. This bar consisted of a failed strain bar and a failed T-head (rod end connector). These components failed during characterization fatigue testing. The strain bar and the T-head (rod end connector) had been subjected to spectrum loading. This fatigue spectrum consisted of two 365 MPa (53 ksi) tensile loads followed by a single 469 MPa (68 ksi) tensile load application. The strain bar failed at 4875 cycles. The T-head failed at 3235 cycles; the desired lifetime for the strain bar and the T-head was 4500 cycles. The failed AISI 4340 steel strain bar and the failed cadmium-plated AISI 4330V steel T-head were submitted for analysis. Figure 1 shows the as-received failed strain bar. The fracture occurred at the aft radius of the circumferential retainer ring groove located near a slot in the strain bar. Figure 2 shows the as-received failed T-head and the location

As-received strain bar from fatigue testing of a hold-back bar assembly

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of the fracture. Fracture occurred in the run-out area of internal threads at a distance of approximately 50 mm (2 in.) from the forward end. All of the parts were subjected to a magnetic particle inspection. In the strain bar, two cracks were found approximately 180 apart in the retainer ring groove location shown in Fig. 3. The

Fig. 2

Fig. 3

crack lengths were 22 and 25 mm (0.85 and 0.99 in.). No cracks were detected in the T-head. The radius at the primary fracture of the strain bar was determined to be 1.3 mm (0.050 in.) At the secondary cracks, the radius was 1.14 mm (0.045 in.). At present, there is no dimensional requirement for these radii.

As-received T-head connector from fatigue testing of hold-back bar assembly

Magnetic particle inspection indications found in the retainer ring groove on the strain bar. (a) Crack indication 1 showing a crack length of 25 mm (0.99 in.) between arrows. (b) Crack indication 2 showing a crack length of 22 mm (0.85 in.) between arrows. Original magnification: 1.4 ·

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Case Studies of Steel Component Failures in Aerospace Applications / 353

The fracture surfaces were examined visually and at up to 50 · magnification using a stereomicroscope. Figure 4 shows macrographs of the fracture surfaces of both the strain bar and the T-head. The strain bar primary fracture surface exhibited multiple fatigue origins around the outside periphery of the fracture. The T-head fracture exhibited fatigue origins primarily near internal thread roots. A scanning electron microscope (SEM) was used to examine the fracture surfaces and to document the mode of failure. Figure 5 shows the fracture topographies of the primary fracture in the strain bar. The secondary fractures in the strain bar are shown in Fig. 6. Fatigue striations emanated from fatigue origins located on the outside periphery of the fractures. Fracture ridges also gave evidence of fatigue origin locations on the outside periphery of the fracture. Figure 7 shows topographic features of the fracture surface of the T-head failures. Fatigue striations emanated from multiple origins located on the inside peripheries of the fractures. Hardness measurements were made on the strain bar and T-heads to verify heat treat conditions. The hardness of the strain bar was determined to be 41.6 HRC, and this met the hardness requirement of 40 to 43 HRC. The hardness of the T-head was determined to be 48.8 HRC. This hardness value exceeded the requirement of 46 to 48 HRC. Metallographic sections were prepared through the strain bar and a typical T-head at the locations shown in Fig. 4. The specimens were prepared using standard metallographic

Fig. 4

techniques. Figure 8 shows the microstructures of the strain bar and a typical T-head. The inclusion contents of the metallographic specimens appeared to be relatively high, as is observed in air-melted steels. An energydispersive spectroscopy analysis identified manganese sulfide inclusions as being present. Manganese sulfide inclusions are typically present in air-melted steels and can severely affect fracture toughness (Ref 1). Inclusions can act as initiation sites for fatigue failures; however, the failure of these parts could not be associated with the presence of inclusions. Banding was also present in both the strain bar and T-head metallographic specimens. Microhardness measurements made on the light and dark bands in the strain bar and T-head yielded hardness readings of 44.1 HRC (light), 42.6 (light), 49.8 (dark), and 49.1 (dark), respectively. This microstructure is typical of an AISI 4340 component that has been improperly normalized (Ref 2). The light etching areas are generally untempered or lightly tempered martensite, while the darker etching regions are martensite that has been tempered more thoroughly. The lighter etching regions contain higher hardenability than the darker etching regions. This is typically due to segregation of chromium, generally as the result of inadequate normalizing prior to heat treatment. This type of nonuniform structure is not optimal. The harder regions tend to have lower fracture toughness and have higher notch sensitivity than the darker, more tempered regions. Because of this, it was thought that the premature fracture occurred

Macrographs of primary fracture surfaces. (a) Strain bar, with maximum depth of fatigue at arrow (4.24 mm, or 0.167 in.), looking forward. Original magnification: 2.4·. (b) T-head, with maximum depth of fatigue at arrow (2.90 mm, or 0.114 in.), looking aft. Original magnification: 2.2·

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because of improper heat treatment. The presence of a higher level of inclusions could also contribute to notch sensitivity reduction.

Cracking in a Main Landing Gear Attach Pin Multiple pins exhibited cracks in the flanges during magnetic particle inspection as part of the manufacturing process. Some parts were returned to the vendor for rework, while others were reworked in-house. The pins were machined from 300M steel. The pins were then heat treated to a strength range of 1930 to 2070 MPa (280 to 300 ksi). The pin was ground and chromium plated. Figure 9 shows the main landing gear (MLG) lever attach pin as received for examination. The magnetic particle inspection indications in the flange are shown in Fig. 10.

Fig. 5

After the chromium plating had been removed, the flange area was subjected to a temper-etch operation of 1% nitric acid in methanol. Numerous areas on the flange etched dark, indicating that these regions had been overtempered (Ref 3). A micrograph of the temper-etched inspected flange is shown in Fig. 11. A typical crack was opened, and the fracture surface was examined visually and with a stereomicroscope. The fracture surface was discolored (Fig. 12). The discoloration appeared to be similar to a temper color in steel (straw or tan). After the chrome plating was ground to final dimensions, the attach pin had received a stress-relief operation that could have produced this discoloration. An SEM was used to examine the fracture surface. Figure 13 shows the SEM fractographs of the surface. The fracture topology was primarily intergranular.

SEM fractographs documenting typical topographic features of the strain bar primary fracture surface. (a) Overall view of strain bar origins (micron bar is 1.5 mm long). (b) Typical origins (arrows) on outside edge of strain bar fracture (micron bar is 1500 mm long). (c) Typical fatigue striations on strain bar fracture (micron bar is 10 mm). (d) Typical worn area on strain bar fracture (micron bar is 10 mm)

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A hydrogen analysis was conducted on the MLG lever attach pin before and after baking for 24 h at 190 C (375 F). The results of the analysis indicated that before baking, the lever attach pin contained 1.6 ppm of hydrogen, and after baking, the attach pin contained 1.4 ppm of hydrogen. These amounts of hydrogen are considered to be low and will not produce a hydrogen embrittlement failure (Ref 4). Also, fractures caused by hydrogen embrittlement are generally not discolored. Metallographic sections were prepared through the magnetic particle indications using standard metallographic techniques. Shown in Fig. 14 are areas where the chromium plating had been reduced by possible grinding. Figure 15 shows micrographs of a typical crack. This crack extended through the chromium plating and into the base metal to a depth of approximately 0.25 mm (0.010 in.). There was no visual indication of untempered or overtempered martensite.

Fig. 6

Based on the results of this investigation, it was concluded that the MLG lever attach pin contained numerous probable grinding cracks in the flange.

MLG Linear Actuating Rod and Cylinder The part was removed from the aircraft and was reported to have accumulated 50 flight hours. The failed part was from an early production lot manufactured with electroless nickel on the inner diameter surface rather than the drawing requirement of electrolytic nickel. The manufacturing sequences for the MLG linear actuating piston rod cylinder were:

A 300M steel bar was machined to the required outside diameter and was then bored, honed, and nickel plated. The bar was swaged in an approximate 76 mm (3 in.) length on the end opposite

SEM fractographs of the fracture surfaces of cracks 1 and 2 in the strain bar. (a) Fracture surface of crack 1, showing typical fatigue zones at arrows (750 mm). (b) Typical fatigue striations in crack 1 fatigue zones (30 mm). (c) Fracture surface at crack 2 of the strain bar, showing typical fatigue zones at arrows (830 mm). (d) Typical fatigue striations in crack 2 fatigue zones (1.5 mm)

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Fig. 7

SEM fractographs documenting typical topographic features of the T-head. (a) Typical portion of fracture surface, showing typical origins at arrows (500 mm). (b) Typical fatigue striations on fracture surface of T-head (1 mm)

Fig. 8

Micrographs documenting the microstructure of the failed strain bar. (a) Section 1-1 from Fig. 4 showing a banded structure (670 mm). (b) Location A, retainer ring groove (300 mm). (c) Section 2-2 from Fig. 4 showing banded structure (670 mm). (d) Location B at origin showing banded structure (300 mm)

from the mono-ball end. The bar was supposedly swaged at 650 C (1200 F) and then stress relieved.

The bar was then finish machined and heat treated to a tensile strength range of 1930 to 2070 MPa (280 to 300 ksi). The part was

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then magnetic particle inspected. Electrodeposited nickel was specified for plating on the inside diameter surface. This plating is relatively soft, ductile, and has a relatively high melting point. However, electroless nickel, which has high phosphorus content, was substituted without authorization. This plating is hard and brittle and starts to melt at the eutectic temperature of 880 C (1616 F). Figures 16 and 17 show the MLG linear actuating piston rod cylinder components as received for examination. Two circumferential fractures were located approximately 165 and 215 mm (61/2 and 81/2 in.) from the mono-ball end of the piston rod cylinder. These two fractures were joined by a longitudinal crack approximately 50 mm (2 in.) long. The longitudinal fracture was opened, and the fracture surfaces were examined visually and at up to 50 · magnification using a stereomicroscope. Figure 18 shows macrographs of the fracture surfaces. The circumferential fractures originated and terminated on the longitudinal crack. The longitudinal crack exhibited a single,

Fig. 9

Macrograph documenting the appearance of the as-received main landing gear lever attach pin

intergranular-appearing origin on the inside diameter surface of the piston rod cylinder. The origin was located at approximately midlength of the longitudinal crack, 200 mm (8 in.) from the mono-ball end. An arrest mark was present on the fracture surface at approximately 85% of the wall thickness.

Fig. 10

Magnetic particle inspection indications on the flange of the main landing gear lever attach pin

Fig. 11

Overtempered (darkened areas) in the flange revealed by temper etching (2.5 mm)

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An SEM was used to examine the fracture surface and to document the mode of failure. Figure 19 shows typical SEM fractographs made of the fracture surface. The fracture topography at the origin was intergranular. The dimensions of the origin were 0.5 mm wide by 0.19 mm deep (0.02 in. by 0.0075 in.). Away from the origin was a static rupture area, which was characterized by dimples. Approximately 85% across the thickness, an arrest mark was present. As shown in Fig. 19, the fracture topography in the arrest was intergranular. Past this zone to the edge of the specimen, the fracture topography consisted of dimples (static rupture). Hardness measurements were made on the piston rod cylinder. The hardness was 54 HRC, and this met the hardness requirement of 53 to 55 HRC.

Fig. 12

Discoloration of the fracture surface (330 mm)

Fig. 13

Topographic features of the fracture surface as observed by the SEM. (a) Overall view of the fracture surface (111 mm). (b) Intergranular fracture observed on the fracture surface (10 mm)

Fig. 14

Metallographic examination of cracks evident in the flange of the main landing gear lever attach pin, showing loss of chromium at cracks (20 mm)

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Fig. 15

Fig. 16

Micrographs showing the morphology of the cracks. (a) Overall view (50 mm). (b) Closeup of crack (20 mm)

As-received main landing gear linear actuating piston rod cylinder

A hydrogen analysis was conducted on the piston rod cylinder. The hydrogen content was determined to be 2.3 ppm. This amount of hydrogen is considered sufficient to cause hydrogen embrittlement (Ref 5). A metallographic specimen, section 1-1, was removed at the origin location shown in Fig. 19. The specimen was prepared using standard metallographic techniques. The microstructure at the origin is shown in Fig. 20. The microstructure of the steel, which consisted of tempered martensite, was normal. The plating at the origin consisted of three distinct layers. The layers were subjected to an energy-dispersive x-ray analysis. The outermost two layers consisted essentially of nickel and phosphorus, whereas the inner layer consisted of nickel and iron. The outermost two layers contained cracks, and the inner layer did not. Also evident in the micrograph is plating that had covered a small part of the fracture surface. This gave evidence that a crack was present prior to plating or prior to the heating of the plating, which occurred during swaging. A single intergranular origin was associated with the fracture surface, and this origin was relatively clean. These factors suggest a delayed mode of failure, as occurs in a hydrogen

embrittlement failure. The bulk hydrogen was analyzed to be 2.3 ppm; this concentration would most likely be higher at the crack tip. The intergranular origin was separated from the intergranular arrest zone by a ductile rupture region, indicating that the crack had propagated in a ductile manner before arresting. Hydrogen most likely diffused to and concentrated at the crack tip, which caused the intergranular arrest region to occur. The crack then propagated to failure by ductile rupture. Based on the results of this investigation, it is concluded that the MLG linear actuating piston rod cylinder failure was most likely due to hydrogen embrittlement.

Failure Analysis of AISI 420 Stainless Steel Roll Pin Several failures of AISI stainless steel roll pins were reported. This pin is a standard part and is used to hold pin components together by expanding after compression. The pin is manufactured from AISI 420 stainless steel and is heat treated to 46 to 55 HRC. Figure 21 shows the as-received failed roll pin. The fracture extended the length of the pin. An SEM was used to examine the fracture surface of the service failure and a laboratorycreated overload failure from the same lot of material. Figure 22 shows the results of the SEM examination. In this figure, the topology of the fractures was intergranular, with some dimples present. This failure mechanism indicated that the failure was due to overload (static rupture).

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Fig. 17

Appearance of longitudinal crack in the main landing gear linear actuating piston rod cylinder

Fig. 18

Fracture surfaces. (a) Circumferential crack 1 (4.3 mm). (b) Circumferential crack 2 (4.3 mm). (c) Longitudinal crack (5 mm)

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Hardness measurements were made on the roll pin to verify the heat treated condition. The hardness of the roll pin was measured to be 49 HRC, which met the hardness requirement of 46 to 55 HRC. A metallographic specimen was prepared through the fracture surface. The specimen was prepared using standard metallographic techniques. The microstructure (Fig. 23) consisted

Fig. 19

of tempered martensite, with carbides outlining the prior-austenitic grain boundaries. This is not a typical microstructure for quenched and tempered AISI 420 stainless steel (Ref 2). Generally, because of their high hardenability, martensitic stainless steels can be quenched in either oil or air. Oil quenching ensures maximum ductility and corrosion resistance. Air

SEM fractographs documenting the appearance of the fracture surface. (a) Origin location (670 mm). (b) View of box A showing fracture origin. (c) Location A (5 mm). (d) Location B (2 mm). (e) Location C (5 mm). (f) Location D (2 mm)

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quenching may cause decreases in ductility and corrosion resistance. If slow cooled or an inadequate quench through the critical range of 870 to 540 C (1600 to 1000 F), this steel will precipitate carbides at the grain boundaries (Ref 2). Because of the presence of the intergranular carbides, it is likely that this part experienced an inadequate quench. Based on the results of this investigation, it was thought that the failure mechanism was overload, with limited ductility caused by improper quenching during heat treatment.

Failure Analysis of a Main Landing Gear Lever

Fig. 20

Fig. 21

Microstructure at fracture origin (12.5 mm)

As-received 420 roll pin. (a) Visual view of as-received roll pin (2.5 mm). (b) SEM view of as-received roll pin (200 mm)

The MLG lever was removed from service after a hard carrier landing. Multiple cracks developed during removal of the ion vapordeposited (IVD) aluminum with sodium hydroxide to analyze for residual stresses by x-ray diffraction. The lever was machined from 300M steel forging into a hollow configuration, machined, then heat treated to a tensile strength level of 1930 to 2070 MPa (280 to 300 ksi). The lever was then IVD coated on the outside. Figure 24 shows the failed MLG lever as received for examination. The location of a primary crack and a series of secondary cracks, which were between the up-latch and oleo lugs, is also shown in Fig. 24. The primary crack was approximately 23 mm (0.9 in.) long and is shown in Fig. 24. The primary crack was opened, and the fracture surfaces were examined visually and at up to 50 · magnification using a stereomicroscope. Figure 25 shows a macrograph of the fracture surface. Also shown in Fig. 25 are the appearances of origins (some of which were discolored), which were located on the outside surface of the lever. The origin areas had a faceted appearance, which indicate a delayed mode of failure, that is, stress-corrosion cracking or hydrogen embrittlement (Ref 6). An SEM was used to examine the fracture surface. Figure 26 shows SEM fractographs documenting the topographic features of the fracture origins. The fracture topography was intergranular, which is indicative of a delayed mode of failure. The wall thickness at the origin location was measured to be 5.8 mm (0.229 in.). The drawing

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Fig. 22

SEM examination of the failed roll pin and laboratory-produced overload fractures. (a) Location A of the service failure (20 mm). (b) Location A of the service failure showing intergranular fracture with some dimples (5 mm). (c) Laboratoryproduced overload failure showing intergranular fracture (20 mm). (d) Laboratory-produced failure showing intergranular fracture with dimples (5 mm). Compare to (b)

Fig. 23

Microstructure of failed roll pin. Microstructure consists of tempered martensite with carbides decorating the prior-austenite grain boundaries (10 mm)

requirements for the maximum and minimum wall thicknesses were 5.33 and 4.45 mm (0.210 and 0.175 in.), respectively. Therefore, the drawing requirement was exceeded by 0.48 mm (0.019 in.). Hardness measurements were made on the lever to verify heat treatment. The hardness of the lever was 54.2 HRC, and this met the hardness requirement of 53 to 55 HRC. Hydrogen analyses conducted on the lever yielded values of 3.0, 4.6, 3.4, and 3.5 ppm. The average of the four values was 3.63 ppm, which is considered high enough to produce a hydrogen embrittlement failure. A metallographic specimen was removed through the fracture origin area. The specimen was prepared using standard metallographic techniques. Figure 27 shows the microstructure of the specimen. The microstructure of the lever was tempered martensite, and this is normal for a quenched and tempered high-strength, lowalloy steel. Based on the results, it is thought that the cracks formed in the MLG lever as a result of

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Fig. 24

Appearance of the main landing gear lever showing the location of the primary and secondary cracks. (a) Overall view (33 mm). (b) Location of primary and secondary cracks at site of ion vapor deposit (IVD) removal (10 mm)

the application of sodium hydroxide to remove the IVD coating in addition to the residual stresses present as a result of the hard carrier landing. It is likely that accelerated stresscorrosion cracking occurred because of the high residual stresses and electrolyte.

Failure Analysis of an Inboard Flap Hinge Bolt An inboard flap hinge bolt was found to be failed after 286.5 flight hours. The bolt had been machined from 4330V-modified steel bar and

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Fig. 25

Appearance of primary crack removed from part. (a) Region of primary crack where ion vapordeposited coating had been removed (13 mm). (b) Primary crack showing branching (2 mm). (c) Fracture surface of crack after opening (3 mm)

heat treated to the 1515 to 1655 MPa (220 to 240 ksi) tensile strength range. A nut with the bolt was machined from 4140 steel bar heat treated to the 1240 to 1380 MPa (180 to 200 ksi tensile strength range. Figure 28 shows the inboard flap hinge bolt and nut as received for examination. The bolt was approximately 102 mm (4 in.) long. The

bolt consisted essentially of the shank, which was approximately 23 mm (0.9 in.) in diameter and 64 mm (2.5 in.) long, and a threaded portion, which was 23 mm (0.9 in.) in length with thread type 0.5000-20 UNJF-3A THD in accordance with MIL-S-8879. The unthreaded portion of the bolt was chromium plated, while the threaded portion was aluminum IVD coated. The fracture had occurred at the transition of the shank to the threaded portion of the bolt. The fracture surface was examined visually and at up to 30 · magnification using a stereomicroscope. Figure 29 shows a macrograph of the fracture surface. Fracture ridges emanated from an origin that was located on the outer diameter surface of the bolt at the location shown in Fig. 29. The origin exhibited a reflective, intergranular appearance. Apparent corrosion products were also observed on the fracture surface. Residual stresses were measured on the surface of the bolt at locations 1 and 2, shown in Fig. 28. The residual stresses were measured using an x-ray residual-stress analyzer. Prior to measuring the residual stresses, the chromium plating (location 1) and the aluminum IVD coating (location 2) were removed. The residual stress measured at location 1 was 790+ 100 MPa (–115+15 ksi), which indicated that this area was properly shot peened as required according to the engineering drawing. The residual stress measured at location 2 was 140+ 100 MPa ( 20+15 ksi). This indicated that this area was not shot peened. The engineering drawing does not require this location to be shot peened. An SEM was used to examine the fracture surface and to document the topographic features. Figure 30 shows SEM fractographs taken of the fracture surface. The fracture topography at the origin was intergranular. The intergranular topography extended for approximately 3/4 of the way across the fracture surface, at which point dimples indicative of ductile rupture were present. A laboratory-produced failure yielded a structure that consisted only of dimples. The corrosion products on the service failure fracture surface were subjected to an energydispersive x-ray analysis. Calcium, potassium, magnesium, and chlorine were identified as being present. Rockwell hardness measurements were made on the inboard flap hinge bolt to verify heat treatment condition. The hardness of the bolt averaged 46 HRC. This met the hardness

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Fig. 26

Typical fracture features. (a) Overall view (500 mm). (b) Intergranular fracture apparent at fracture origin (7 mm)

Fig. 27

Metallographic examination of primary crack. Fracture path followed prior-austenite grain boundaries. No precipitates in grain boundaries (50 mm) were observed.

requirement of 46 to 48 HRC for 4330Vmodified steel heat treated to the 1515 to 1655 MPa (220 to 240 ksi) tensile strength range. The hardness of the nut was 45 HRC, which was above the hardness requirement of 40 to 43 HRC. The high hardness value did not appear to be detrimental.

A hydrogen analysis conducted on the hinge bolt yielded a value of 1.5 ppm, which is considered low to cause hydrogen embrittlement. The diameter of the bolt at the fracture was measured to be 10.6 mm (0.419 in.), meeting the drawing requirement of 10.7+0.25 mm (0.422+0.010 in.). The radius at the transition where the failure occurred was 1.5 mm (0.060 in.), meeting the drawing requirement of 1.6+0.25 mm (0.063+0.010 in.). A spectrographic analysis verified that the bolt was 4330V-modified steel. A tensile test was conducted on a production bolt. The bolt failed in the threaded area. The failing load was determined to be 17,600 kg (38,820 lb). No requirement was available. A metallographic section was removed through the origin location shown in Fig. 29. The specimen was prepared using standard metallographic techniques. Figure 31 shows micrographs taken of the fracture origin. The tempered martensitic microstructure was normal. No untempered martensite was present. Corrosion products were observed on the fracture surface at the origin (Fig. 32). However, there was no evidence of pitting or other corrosion processes that would have produced the corrosion products. No aluminum IVD was present on the surface at the origin, because the IVD had been removed prior to measuring residual stresses. Based on the results of this investigation, it is concluded that the inboard flap hinge bolt failed in a delayed mode of failure, which was most likely stress-corrosion cracking. No material anomalies were observed that would account for the failure of the inboard flap hinge bolt.

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Fig. 28

Fig. 29

As-received flap hinge bolt and nut

Fracture surface of the failed flap hinge bolt

Failure Analysis of a Nose Landing Gear Piston Axle This investigation determined the mode and initiating cause of failure of a nose landing gear (NLG) piston axle. Failure of the piston axle occurred on the aircraft during taxiing after the aircraft had accumulated approximately 60 h in service. When failure occurred, the nose wheel completely separated from the piston strut. The NLG piston assembly was fabricated from 300M steel and heat treated to a 1930 to 2070 MPa (280 to 300 ksi) ultimate strength range. The axle was then shot peened and low-embrittlement cadmium plated. White paint was applied to the inside and outside surfaces of the axle, with

exception of the two land surfaces for the wheel bearing cans. The failed NLG piston assembly was removed from the wheel and submitted for examination. Figure 33 shows the as-received piston assembly and the failed axle. A visual examination of the failed axle revealed that the fracture surface followed a circumferential path and contained a large discolored region (Fig. 34). Except for the discolored region, the fracture was typically gray colored and contained chevron markings, as shown in Fig. 35, indicating the failure originated at the discolored region. It was determined that the discolored region, from which the failure appeared to originate, was located at the bottom of the axle, at approximately the 6:30 o’clock position. The sketch in Fig. 36 shows the approximate failure location. Failure of the axle initiated approximately 145 to 149 mm (5.70 to 5.85 in.) from the threaded outboard end. An optical examination with a stereomicroscope at up to 30 · magnification confirmed that the failure of the axle originated at the discolored region. Chevron marks were evident that indicated the crack propagated away from the ends of the discolored region, terminating at an area diametrically across from it, as shown in Fig. 35(b). The discolored region extended along the circumference of the axle for a cord distance of 22.15 mm (0.872 in.) and a maximum depth from the inside diameter surface of 3.38 mm (0.133 in.). Actual wall thickness at that location was 3.81 mm (0.150 in.). The plane of the discolored region was approximately 25 from being perpendicular to the longitudinal axis of the axle, along machining marks. These dimensions were shown in Fig. 33

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Fig. 30

SEM examination of the flap hinge bolt. (a) Overall view of fracture surface (1.25 mm). (b) Intergranular fracture evident at origin (10 mm)

Fig. 31

Micrographs of the fracture origin. (a) Section 1-1, Fig. 29. Fracture is intergranular at prior-austenite grain boundaries (50 mm). (b) Location A showing evidence of corrosion product at fracture facets (25 mm)

and 36. When examined, the discolored region showed a color variance ranging from a gold color to reddish brown. Also, small patches of blue were evident. This observed color variation is a characteristic typically observed on steel surfaces that have been exposed to elevated temperature, such as during heat treatment (Ref 1). The machining marks indicate a surface much rougher than is normal. The hardness measurements were made at several locations around the circumference of the axle adjacent to the fracture surface. The hardness measured 53 to 55 HRC, which conformed to the drawing requirement of 53 to 55 HRC for 300M steel heat treated to a 1930 to 2070 MPa (280 to 300 ksi) condition. A section containing the discolored region was removed from the axle and examined with

the SEM. The locations where SEM examinations were performed are shown in Fig. 35(a). Prior to SEM examination, the discolored region was cleaned with acetone and replication in an attempt to remove the scale from the fracture surface. After several attempts, only a small amount of the scale could be eliminated, indicating the scale was firmly attached to the fracture surface. The SEM fractographs adjacent to the inside diameter surface (area 1) were primarily intergranular and appeared to be covered with scale. At areas beyond the inside diameter surface, the SEM fractographs revealed a mixed intergranular with transgranular features and patches of scale. These are shown in Fig. 37. Along the periphery of the discolored region, fatigue striations could be observed at a higher magnification, as shown in Fig. 38. The depth of

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the fatigue growth region was approximately 0.05 mm (0.002 in.). Rapid fracture extended beyond the fatigue region. The SEM fractograph in Fig. 39 shows dimpled features that characterize a ductile mode of rapid fracture. When examining the discolored region with SEM at low magnification, there was no evidence of shear lip along the inner diameter surface, which

Fig. 32

Micrograph showing the microstructure at the fracture origin. Microstructure consists of quenched and tempered martensite (25 mm). IVD, ion vapor deposited

Fig. 34

Discolored region of the fracture

indicates that the fracture did not initiate subsurface, as would be the case for delayed failure resulting from hydrogen embrittlement. In an attempt to determine the composition of the observed scale or to identify any contaminant that may be associated with the fracture surface, an energy-dispersive x-ray (EDX) analysis was performed on the discolored region. The EDX analysis revealed no other element except the ones common to the base metal composition.

Fig. 33

As-received nose landing gear piston assembly and the failed axle

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Fig. 35

Fracture surface of the failed axle. (a) Black arrows show locations for SEM examination. (b) White arrows show fracture direction and location of metallographic sections.

Fig. 36

Axle wall thickness, discolored region dimensions, and fracture location

To examine the microstructure of the discolored region, a section was removed (section A-A, Fig. 35b) for metallographic examination. The microstructure was tempered martensite, which is typical for 300M steel heat treated to the 1655 to 2070 MPa (280 to 300 ksi) ultimate strength range (Fig. 40). A secondary crack was observed starting at the inner diameter surface (Fig. 40b), at a location showing large surface irregularities. In an attempt to investigate this

secondary crack, the crack was opened and the fracture surface examined with a light microscope and the SEM. The fracture depth was relatively small (0.25 mm, or 0.01 in.), which made it difficult to determine with the light microscope if a high-temperature scale was present. The SEM examination revealed intergranular and ductile transgranular features (Fig. 41). These features suggest a secondary heat treat crack.

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Fig. 37

SEM fractographs showing brittle intergranular structure in discolored region of the fracture surface. (a) Intergranular fracture partially covered with scale at the area adjacent to the inner diameter surface (10 mm). (b) Fracture surface away from the inner diameter surface showing intergranular and transgranular features (10 mm)

Fig. 38

SEM fractographs showing fatigue growth at regions close to the outer diameter surface. (a) At boundary of discolored region (5 mm). (b) Outside the boundary of the discolored region (5 mm)

In order to verify the material composition, one section of the piston axle was chemically analyzed, using the atomic absorption spectroscopy method. The results of the analysis are shown in Table 1. These values met the requirements of AMS 6419 for 300M steel. In summation, the results of this investigation indicated that the failure of the NLG piston axle was introduced from a pre-existing defect. This defect was present on the axle prior to the final assembly of the part to the aircraft. The defect had a brittle intergranular fracture surface feature and a discoloration characteristic of a thermal or quench crack. These phenomena demonstrated that the crack occurred on the piston axle during or prior to the heat treatment

Fig. 39

Rapid fracture at location outside the discolored region (5 mm)

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process. Because of the machining marks and the initiation of cracking at these asperities, it is possible that cracking occurred during the machining operation, in a fashion similar to the formation of grinding cracks. Based on the results of this examination, it is concluded that the failure of the NLG piston axle was due to the pre-existing crack on the axle. This crack was created prior to or during heat treatment of the part.

Multiple-Leg Aircraft-Handling Sling A new multiple-leg aircraft-handling sling that had just received proof testing was found to have a cracked clevis. The failed adapter was

machined from 4330V-modified steel that had been heat treated to a 1240 to 1380 MPa (180 to 200 ksi) tensile strength range. Figure 42 shows the adapter part of the multiple-leg aircraft-handling sling as received for examination. Also shown in Fig. 42 is the location and appearance of the crack that was at a clevis. The crack penetrated completely through the wall of the clevis. The fracture surface was opened, and the fracture surface was examined visually and at up to 50 · magnification using a stereomicroscope. Figure 43 shows a macrograph of the fracture surface. Fracture ridges indicated that origins were present on the outside surface of the clevis. The fracture surface appeared faceted, which is characteristic of an intergranular, delayed mode of failure.

Fig. 40

Microstructure and secondary cracking at the discolored region. (a) Normal tempered martensite typical of 300M (25 mm). (b) Secondary cracking apparent on inner diameter surface (100 mm)

Fig. 41

SEM fractographs of opened secondary crack. (a) Origin of secondary crack (33 mm). (b) Intergranular fracture apparent (10 mm)

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An SEM was used to examine the fracture surface and the side of the clevis at the origin. Figure 44 shows SEM photographs documenting topographic features of the fracture surface and corrosion pitting on the side of the clevis at the fracture origin. The pits were coated with primer, which indicated they were present prior to painting. Also documented in Fig. 44 are intergranular topography and corrosion products at the fracture origins. These features are all characteristic of a stress-corrosion failure (Ref 4). A laboratory overload fracture produced in the clevis was examined in the SEM. Figure 45 shows an SEM fractograph documenting dimples that are characteristic of an overload mode of failure. Hardness measurements were made on the adapter to verify the heat treat condition. The hardness of the adapter was 41.4 HRC, and this met the hardness requirement of 40 to 43 HRC. The hydrogen content of the adapter was determined to be 1 ppm, which is considered

Table 1 Chemical analysis of nose landing gear piston axle Chemical composition, % Si

Mn

Ni

Cr

Mo

V

Fe

1.42

0.72

1.79

0.68

0.49

0.008

bal

Fig. 42

As-received failed multiple-leg aircraft-handling sling

low in relation to producing a hydrogen embrittlement failure. A metallographic specimen was prepared through a fracture origin. The specimen was prepared using standard metallographic techniques. Figure 46 shows the microstructure, which was normal (tempered martensite). Also shown in Fig. 46 is a corrosion pit that was coated with paint, indicating the pit was present prior to painting. Based on the results of this investigation, it was concluded that the adapter failed due to stress-corrosion cracking. Corrosion pits at the fracture origin were present prior to painting.

Failure Analysis of an Aircraft Hoist Sling during Static Test An aircraft hoist sling was successfully tested to an ultimate load of 136,000 kg (300,000 lb). However, shortly after relieving the load, the weld on the aft right-hand fitting failed. The fitting was fabricated from welded 17-4PH stainless steel plates. Figure 47 shows the as-received failed portion of the test fixture. The failure occurred in a consumable electrode weld that traversed around the length of the fracture, which was approximately 69 cm (27 in.). A dark, discolored area was present in the fractured weld over a length of approximately 122 mm (4.8 in.) (location C). Lack of fusion was a general condition of the weld. Figure 48 shows a typical area where lack of fusion was present in the weld.

Fig. 43

Appearance of the fracture surface of the multipleleg aircraft-handling sling

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Fig. 44

SEM fractographs of the service fracture. (a) Overall view. (b) Intergranular topography at origin. Original magnification: 1500 ·. (c) Pits on side of clevis at origin. Original magnification: 400·. (d) Pit and corrosion products at origin. Original magnification: 1000 ·

Figure 49 documents the direction of fracture, as evidenced by the convergence of river patterns, at a typical location. The side of the fracture containing location C had been placed in tension due to bending. Portions of the fracture at locations A, B, and C shown in Fig. 47 were excised and examined on an SEM. Figure 50 shows SEM fractographs documenting topographic features of these portions of the fracture surface. The areas at A and B showed evidence of overload, that is, dimples. At location C in the dark, discolored area, the fracture topography was different than at A and B. The topography at location C was indicative of a heated, oxidized surface. An energy-dispersive spectrographic (EDS) analysis conducted on the fractures at A and B basically detected iron, chromium, nickel, and copper. Nickel was not detected at location C. The EDS spectra are shown in Fig. 51. The base metal also showed the presence of iron, chromium, nickel, and copper, which are present in 17-4PH stainless steel.

Fig. 45

Dimpled rupture indicating overload failure in a laboratory-produced failure. Original magnification:

5000 ·

In conclusion, a different consumable welding electrode was used at location C (no nickel present) than at locations A and B, where nickel was present. It is recommended that 17-4PH filler metal be used when welding 17-4PH steel.

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This was apparently done at locations A and B but not at C, where no nickel was present. Nickel gives ductility to the weld and helps to prevent cracking during cooling from the welding temperature (Ref 7). It appears that at location C, where no nickel was present, a crack was formed, and the fracture surfaces at C became discolored and pebbled in appearance at an elevated temperature. Lack of fusion due to insufficient heating was a general condition in the weld, and this also contributed to the failure. In summary, the weld, when placed under tension, failed due to a crack and lack of fusion.

Figure 52 shows the internal gear as received for examination. Also shown in Fig. 52 is the appearance of typical cracks, which were located on the inside of the internal gear on areas next to the gear spline. The cracks were located completely around the circumference on both sides of the gear. A crack was opened, and the fracture surface was examined visually and at up to 30 · magnification using a stereomicroscope. Figure 53 shows macrographs of the fracture surface. The fracture surface was discolored

Failure Analysis of an Internal Spur Gear This investigation analyzed cracks that were present in an internal spur output gear. The internal gear is part of the planetary gear system for a canopy. The gear was made from 4340 steel heat treated to the 1790 to 1930 MPa (260 to 280 ksi) tensile strength range. The part was rejected after magnetic particle inspection due to multiple crack indications along the inside surfaces next to the gear spline. The typical manufacturing sequence is forge, machine, heat treat, grind, and plate.

Fig. 46

Microstructure consisting of tempered martensite at fracture origin. Original magnification: 200 ·. Corrosion pit coated with paint is evident in the micrograph.

Fig. 47

As-received portion of the failed aft right-hand fitting for the hoist sling

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with shades of black and brown, as can be produced by heating the steel. The color brown corresponds to the temper color, which is produced by heating steel to approximately 238 C (460 F) in air and cooling to room temperature. Chromium plating was observed on the outer

Fig. 48

diameter surface of the gear. There were no shear lips evident around the periphery of the fracture surface. Fracture ridges emanated from origins located on the inner diameter edge, and the fracture surface exhibited a faceted appearance indicative of a brittle intergranular failure.

Typical lack of fusion area

Fig. 50 Fig. 49

Direction of fracture propagation

SEM fractographs documenting the topographic features of the failed weldment. (a) Location A showing overload features. (b) Location B showing overload features. (c) Location C showing weld defect features. Original magnification: 1200 ·

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An SEM was used to examine the fracture surface and to document the mode of failure. Figure 54 shows SEM fractographs made at a typical fracture origin. The fracture topography was intergranular, which is indicative of a brittle mode of failure (Ref 3). No defects were observed that could be associated with the cause of the fracture. A metallographic specimen was removed at the location shown in Fig. 52. The specimen was prepared using standard metallographic techniques. Figure 55 shows the microstructure (tempered martensite) and typical cracks that were present. The cracks were intergranular, and an oxide was present in the cracks. The

Fig. 51

Energy-dispersive spectrographs of failed weld and base material. (a) Location A weld metal. (b) Location B weld metal. (c) Location C weld metal

Fig. 52

As-received internal spur gear showing location of cracks

Fig. 53 surface

Macrographs of (a) magnetic particle indications of cracks and (b) opened crack revealing the fracture

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Main Landing Gear Axle

Fig. 54

SEM fractographs showing the topographic features of a typical fracture origin. (a) Fracture surface. Original magnification: 20·. (b) Location A showing intergranular fracture. Original magnification: 3000 ·

morphology of the cracks is typical of grinding cracks (Ref 8). A hydrogen analysis conducted on the internal gear yielded a value of 1.85 ppm, which is not considered to be high enough to cause hydrogen embrittlement. Hardness measurements were made on the internal gear to verify the heat treat condition. The hardness of the gear near the observed cracking was 44 HRC. This hardness value was below the drawing hardness requirement of 50 to 53 HRC. The bulk hardness of the gear was 51 HRC. The lower-than-expected hardness at the crack initiation sites indicates that there was some event that caused excessive local heating of the microstructure. Because of the nature of the part and typical manufacturing sequence, it is likely that the grinding cracks occurred as a result of aggressive grinding. Based on the results of this investigation, it is concluded that the internal spur gear failed due to aggressive grinding.

During taxi, the MLG axle separated. The axle had been installed for only 90 days, with a total flight time of 62 h. The axle was fabricated from 300M high-strength, low-alloy steel. Shown in Fig. 56 is the as-received failed MLG axle. Also shown in Fig. 56 is the location of the fracture, which occurred adjacent to a chromium- and cadmium-plated area. Also shown in Fig. 56 is a worn (discolored) area of chromium plating. The fracture surface was examined visually and at up to 50 · magnification using a stereomicroscope. Figure 57 is a macrograph of the fracture surface. Shown in Fig. 58 is an SEM photograph of the wear that had removed the protective plating adjacent to the fracture surface. Multiple origins were located around the periphery of the axle. Faceted areas characteristic of an intergranular, delayed mode of failure had progressed through approximately 80% of the thickness. Fatigue propagated from the ends of the intergranular areas, with several areas of the wall being totally penetrated by the fatigue. The relatively large intergranular areas and the irregular shape of the intergranular area are indicative of stress corrosion. An SEM was used to examine the fracture surface and to document the mode of failure. Shown in Fig. 59 are SEM fractographs documenting topographic features, that is, intergranular topography and fatigue striations, on the fracture surface. From the ends of the intergranular area, fatigue propagated through the wall thickness. Also documented in Fig. 59 are the lack of plating at a fracture origin and corrosion products, which were present at the intergranular fracture origins. The lack of plating and the presence of corrosion products are indicative of a stress-corrosion failure. Also shown in Fig. 59 is an area where the chromium was intact. The hardness of the axle was determined to be 53.9 HRC. This met the hardness requirement of 53 to 55 HRC. A chemical analysis by atomic absorption verified that the MLG axle was fabricated from 300M steel. A metallographic specimen was prepared through a typical fracture origin on the axle. The specimen was prepared using standard metallographic techniques. As determined from the metallographic specimen, the thickness of the chromium plating was approximately 0.06 mm

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Fig. 55

Micrographs showing the appearance of the cracks and microstructure. (a) Surface cracking evident. Original magnification: 50·. (b) Surface cracking and MnS inclusions. Original magnification: 100 ·. (c) Extended surface crack. Original magnification: 100·. (d) Intergranular crack along prior-austenite grain boundaries. Original magnification: 200·

Fig. 56

As-received failed axle

(0.0024 in.). Shown in Fig. 60 is the microstructure of the axle at a typical fracture origin. The chromium plating ended approximately 1.6 mm (0.062 in.) from the origin. The hardness of the light, discolored area at the

chromium runout was 58.4 HRC, which exceeds the hardness requirement of 53 to 55 HRC. Past the chromium runout, the part was supposed to be cadmium plated. However, past the chromium runout to the origin, no plating was present. The wear documented in Fig. 60 had removed the protective cadmium plating from this location. This allowed corrosion to occur. Shown in Fig. 61 is a secondary branch crack characteristic of stress corrosion. This micrograph also shows the chromium plating that had been penetrated. Documented in Fig. 62 is the appearance of the crack (unetched) after further polishing. Note the worn (missing) chromium. A hydrogen analysis conducted on the axle yielded hydrogen content of 4.2 ppm, which is relatively high and considered high enough to cause hydrogen embrittlement (Ref 4). This hydrogen is in the form of atomic hydrogen and can be generated in corrosion reactions and then absorbed by steel.

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Based on the results observed, it was determined that the MLG axle failed due to stresscorrosion cracking. The protective cadmium plating had been worn away, allowing corrosion to occur at locations where the cadmium was no longer present.

Nondestructive Testing and Failure Analysis of Fin Attach Bolts after Full-Scale Fatigue Testing During the inspection, after the first lifetime of fatigue testing, one aft fin attach bolt was found to have broken. Five additional bolts were found to have crack indications. The location of the bolts and the nomenclature used for identification of these bolts are shown in Fig. 63. The failed bolt (bolt A) and the additional five bolts, identified as B, C, G, H, and I, were submitted to the laboratory to determine if ultrasonic testing would be a suitable field inspection technique for this application and to examine the failed bolt and cracked bolts to determine the failure modes. The specification requirements for these bolts are for a cadmium-plated, forged hex head metric bolt that has a close tolerance shank. The material can be any of four material specifications: BS S147, S148, S149, or S158. The compositions of these steels are shown in Table 2. These steels are heat treated to

Fig. 57

Macrograph of fracture surface of the failed axle. Typical fracture origins are shown at arrows.

a minimum tensile strength of 1100 MPa (160 ksi) minimum. The threads and the headto-shank fillet radius are rolled after heat treatment. Machining of the fillet radius is not permitted. The specified dimensions are shown in Fig. 64 and Table 3. The failed bolt and the additional bolts were examined using nondestructive testing, visual examination, scanning electron microscopy, metallography, and analytical chemistry. The hardness of the parts was also measured. Nondestructive testing was performed on the submitted bolts to determine if ultrasonic inspection is a suitable field inspection method. Magnetic particle inspection was used as a confirmation of the ultrasonic inspection method. Ultrasonic reference standards were fabricated from AN-8 and AN-6 pan stock bolts. These bolts were machined so that the head thickness and lengths were similar to the test articles. Notches were cut into the bolts at the head radius and at the midpoint length of the grip

Fig. 58

Overall view at origin. (a) Worn cadmium plating at origin. (b) SEM photograph of fracture surface

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Fig. 59

SEM fractographs documenting the fracture features found at the origin. (a) Fatigue striations emanating from the origin (200 mm). (b) Intergranular fracture at origin (location A, 50 mm). (c) Intergranular fracture and corrosion products found at the origin (13 mm). (d) Intergranular fracture found at interface between chromium plating and steel (50 mm)

Fig. 60

Lack of plating at fracture origin. (a) Overall view of microstructure through fracture origin (500 mm). (b) Lack of plating at fracture origin (50 mm)

on the shank. The dimensions of the reference standards and notches are noted in Fig. 65 and Table 4. While every attempt was made to make

the manufactured notches at the head-to-shank fillet radii and the shank identical, it was not always possible because of manufacturing

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Fig. 61

Secondary branch cracking observed. (a) Overall etched microstructure. (b) Location E, showing cracking on prior-austentite grain boundaries. (c) Location F, showing voids in chromium plating and associated cracking

Fig. 62

tolerances. The ultrasonic inspection was conducted with an ultrasonic transducer (20 MHz; 3.18 mm, or 0.125 in., in diameter). The delay line was removed for this application. The instrument settings are noted in Table 5. The sensitivity of the instrument was adjusted so that an 80% signal response was obtained from the notch under the bolt head of the reference standard (4.17 mm, or 0.164 in., long; 0.686 mm, or 0.027 in., deep) and a notch on the shank of the machined reference standard (4.57 mm, or 0.180 in., long; 4.57 mm, or 0.18 in., deep). Signal responses for the reference standards are shown in Fig. 66. Each of the submitted bolts was inspected with an additional 6 dB of gain to increase the sensitivity of the inspection. Signal responses for each of the bolts are shown in Fig. 67. The ultrasonic inspection showed that a crack indication was present at the shank-tohead fillet radius of bolt C. No crack indications were found in the other bolts. The submitted bolts were inspected using standard magnetic particle inspection techniques. The results indicate that bolts C, G, and I had circumferential crack indications at the shank-to-head fillet radius. Bolt I had a circumferential crack indication around the shank of the bolt. The magnetic particle indications are shown in Fig. 68. After nondestructive testing, bolts C, G, and B were inspected visually. The as-received bolts are shown in Fig. 69. After inspection, the cracks

Unetched micrograph showing branch cracking and void in chromium plating. (a) Overall view of secondary crack (500 mm). (b) Location G showing void in chromium plating and associated cracking (50 mm). (c) Location H showing cracking along prior-austenite grain boundaries (50 mm)

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indicated by magnetic particle inspection were exposed. This was accomplished by centering the threaded portion of the bolts in the chuck of a lathe and drilling a small pilot hole. A larger drill bit was then used to drill the center of the bolt, leaving a wall thickness of approximately 0.76 mm (0.030 in.). Using a hacksaw, a small saw cut was made opposite to the crack. The head of the bolt was held in a vise, and the shank

Fig. 63

Schematic of bolt location and nomenclature used

of the drill bit was inserted into the hole of the bolt to support the bolt shank. Using the drill bit, the bolt was bent away from the crack indication, keeping the crack faces in tension. The exposed crack was then analyzed. Bolt B showed no crack indications, using ultrasonic inspection or magnetic particle inspection. This bolt showed significant circumferential scoring around the periphery of the

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pg 384


Bolt C showed a highly polished shank surface, with a circumferential scoring line in the same approximate location as bolt B. Some rounding of the head corners was evident. No identifying mark or part numbers were stamped on the head of the bolt. After exposing the crack,

shank, at approximately the midpoint of the shank. Minor deformation at the corners of the bolt head was detected. Because no crack indication was seen by nondestructive testing (NDT), no additional work was performed on this bolt.

Table 2 Allowable chemistries of the submitted bolts Allowable composition range of British Specifications, wt% BS S147

BS S148

BS S149

BS S158

Element

Max

Min

Max

Min

Max

Min

Max

Min

C Si Mn P S Cr Mo Ni

0.38 0.20 0.75 ... ... 0.40 0.20 0.60

0.43 0.35 1.00 0.025 0.020 0.60 0.30 0.70

0.36 0.15 0.60 ... ... 0.50 ... 1.10

0.41 0.35 0.90 0.025 0.025 0.80 ... 1.50

0.38 0.20 0.65 ... ... 0.70 0.20 1.65

0.43 0.35 0.85 0.025 0.020 0.90 0.30 2.0

0.22 0.15 0.50 ... ... 0.90 0.15 ...

0.29 0.35 0.80 0.020 0.015 1.20 0.25 0.30

Fig. 64

Schematic of bolt dimensions

Table 3 Allowable bolt dimensions Thread Size code

03 04 05 06 08 10 12 14 16

C

D

Diameter

Pitch

A tolerance

B min

Max

Min

Max

Min

F Min

G

H +0.0 0.3

R +0.0 0.2

J

M3 M4 M5 M6 M8 M10 M12 M14 M16

0.5 0.7 0.8 1.0 1.25 1.5 1.5 1.5 1.5

5.5 7.0 8.0 10.0 13.0 17.0 19.0 22.0 24.0

4.88 6.38 7.38 9.28 12.28 16.08 18.17 21.17 23.17

0.4 0.5 0.5 0.5 0.5 0.6 0.6 0.6 0.6

0.2 0.2 0.2 0.2 0.2 0.3 0.3 0.3 0.3

2.990 3.990 4.990 5.990 7.987 9.987 11.984 13.984 15.984

2.965 3.965 4.965 5.965 7.962 9.962 11.959 13.959 15.959

6.08 7.74 8.87 10.95 14.26 18.90 21.10 24.49 26.75

6.0 7.6 8.5 9.5 12.4 14.8 15.5 17.5 19.5

2.0 2.5 3.0 3.5 4.5 5.0 6.0 7.0 8.0

0.4 0.4 0.5 0.7 0.7 0.8 0.9 1.0 1.1

... 1.0 1.2 1.6 2.0 2.0 2.0 2.0 2.0

Note: All dimensions in millimeters


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the resulting fracture surface was examined (Fig. 70). The fracture surface of crack C had the characteristic ratchet marks or shear ridges that indicate multiple initiation sites along the outer edge of the bolt. The shape and appearance of the fracture surface was indicative of fatigue. Bolt G showed circumferential scoring (Fig. 69), and some rounding of the head corners was evident. After the crack was exposed, the resulting fracture surface was examined

Table 4 Dimensions of reference standards used for ultrasonic testing Specimen identification

HL1 HL2 HL3 HL4 HS1 HS2 HS3 HS4 HS5 SL1 SL2 SL3 SL4 SS1 SS2 SS3 SS4 SS5

A

B

mm

in.

mm

in.

12.4 10.1 7.70 5.99 9.47 7.82 6.50 4.16 2.46 12.3 11.2 8.13 5.82 8.97 7.65 5.16 4.57 2.82

0.487 0.399 0.303 0.236 0.373 0.308 0.256 0.164 0.097 0.485 0.440 0.320 0.229 0.353 0.301 0.203 0.180 0.111

5.66 2.39 1.37 0.81 3.96 1.98 1.17 0.69 0.15 5.87 3.40 1.52 0.81 3.18 2.24 0.76 0.46 0.18

0.223 0.094 0.054 0.032 0.156 0.078 0.046 0.027 0.006 0.231 0.134 0.060 0.032 0.125 0.088 0.030 0.018 0.007

Table 5 Ultrasonic inspection techniques equipment settings Tune Reject Gain Frequency Video filter Damping Delay Range Velocity Repetition rate

Fig. 65

Sketch of reference standards used during ultrasonic inspection

Fig. 66

Ultrasonic signal responses from the manufactured reference standards

0 Off 53 dB 25 MHz FW Minimum 0 8.3 1 3


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Fig. 67

Ultrasonic signal responses for each bolt inspected

(Fig. 71). The fracture surface of crack G had the characteristic ratchet marks or shear ridges that indicate multiple initiation sites along the outer edge of the bolt. The shape and appearance of the exposed crack was indicative of fatigue. Bolt H (Fig. 72) showed no crack indications, using ultrasonic inspection or magnetic particle

inspection. This bolt showed significant circumferential scoring around the periphery of the shank, at approximately the midpoint of the shank. More general wear was seen on the shank than bolt B. Because no crack indication was seen by NDT, no additional work was performed on this bolt.

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Fig. 68

Ultraviolet light photographs of magnetic particle indications for each bolt inspected. NDT, nondestructive testing

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Bolt I (Fig. 72) showed scoring along the circumference of the shank. In this bolt, crack indications were found by NDT in both the shank and the head. The exposed shank crack is shown in Fig. 73. The shank fracture surface of crack I had the characteristic ratchet marks or shear ridges that indicates multiple initiation sites along the outer edge of the bolt. The shape and appearance of the exposed crack was indicative of fatigue. The crack indication in the head-to-shank fillet radius of bolt I was opened in a similar manner as the other bolts, and the fracture surface is shown in Fig. 74. The headto-shank fillet radius fracture surface had the characteristic ratchet marks or shear ridges that

Fig. 69

As-received condition of bolts B, C, and G

indicate multiple initiation sites along the outer edge of the shank fillet radius (Ref 9). The shape and appearance of the exposed crack was indicative of fatigue.

Fig. 70

Fracture surface of bolt C with fracture schematic

Fig. 71

Fracture surface of bolt G with fracture schematic

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The as-received condition of bolt A is shown in Fig. 75. A photograph of the fracture surface of the failed bolt is shown in Fig. 76. The

Fig. 72

Fig. 73

Fig. 74

Fracture surface of the exposed crack in the head fillet radius of bolt I

Fig. 75

As-received condition of the failed bolt

Fig. 76

Fracture surface of the failed bolt

As-received condition of bolts H and I

Fracture surface of the exposed crack in the shank of bolt I

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fracture surface shows ratchet marks that are characteristic of fatigue origins. Multiple fatigue origins, located to one side of the bolt, were

observed. A large fatigue zone was apparent. The final fracture zone was badly damaged. The bolt showed evidence of machining marks on the sides of the shank at the site of crack initiation (Fig. 77). Opposite the initiation site, on the side of the shank, the surface appeared to be polished,

Fig. 79

SEM micrograph of the initiation site of the failed bolt, showing cracking beginning at the localized surface wear (1 mm)

Fig. 77

Fig. 78

Machining marks and polishing evident on the shank of the failed bolt

SEM micrograph of representative multiple fatigue origins associated with machining marks on exposed cracks (100 mm)

Fig. 80

SEM micrograph showing initiation of cracking on the shank of bolt I at localized surface wear (400 mm)

Fig. 81

SEM micrograph of representative fatigue striations found on the bolt fracture surfaces (2 mm)

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with more evidence of machining marks (Fig. 77). The fracture surfaces on bolts A, C, G, and I were examined using the SEM. The fracture surfaces of the exposed cracks in the bolts were characterized by multiple fatigue origins, initiating along machining marks (Fig. 78). In the fractured bolt, initiation of cracking began at a region of localized wear and scoring on the shank of the bolt (Fig. 79). This was also evident

Fig. 82

in the initiation of cracking in the shank on bolt I (Fig. 80). Striations, characteristic of fatigue, were seen on the fracture surface of bolt A (Fig. 81). Machining marks found in the headto-shank fillet radius were evidence that the head-to-shank fillet radius was machined (Fig. 82). This is unacceptable according to the design specification. A metallographic specimen was prepared using standard techniques. This specimen was used to examine the microstructure and microhardness of bolt A. A photographic montage of the metallographic specimen is shown in Fig. 83. The tempered martensite microstructure of bolt A was typical for quenched and tempered low-carbon steel (Fig. 84). Bolt A failed in the shank; in addition, a crack approximately 0.15 mm (0.006 in.) long was found in the headto-fillet radius (Fig. 85). Neither magnetic particle nor ultrasonic NDT detected this crack. The failed bolt was submitted to a microhardness survey to determine the extent of decarburization. The microhardness survey was taken with a Knoop indenter, with a 500 g load. Indentations were taken at 0.0318 mm

SEM micrograph of machining marks at the headto-shank fillet radius of bolt I (1 mm)

Fig. 84

Micrograph of representative quenched and tempered martensite found in the submitted bolts. Etched with 2% nital (25 mm)

Fig. 83

Microstructure at bolt A fracture through section 1-1. The overall microstructure is shown as well as the location of additional micrographs. There is no evidence of a cold-formed head. Etched with 2% nital (667 mm)

Fig. 85 (25 mm)

Micrograph of the crack found in the head-to-shank fillet radius of the failed bolt. Etched with 2% nital

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(0.00125 in.) intervals until the hardness readings began to stabilize. The point at which the hardness stabilized was called the core hardness. For the failed bolt, the core hardness was 404 KHN. In accordance with the specification, the point at which the hardness is within 20 KHN of the core hardness is defined as the depth of decarburization. For the failed bolt, this was at 0.064 mm (0.0025 in.) deep, meeting the specification requirements. The hardness indentions are shown in Fig. 86. The results of the microhardness survey indicated that the surface hardness (Knoop), converted to Rockwell C, is 26 HRC. The Knoop core hardness, converted to Rockwell C, is 40 HRC. While the partial decarburization of the bolt satisfies process specification requirements, the steep gradient of the hardness profile and the surface softness measured indicated that there would be degradation in the potential fatigue life of the bolts.

Determination of the composition of the failed bolt was performed by arc-spark spectrometry. The carbon content of the steel was determined by Leco carbon analyzer. Results are shown in Table 6. The chemical results indicate that this bolt does not meet the composition requirements of Table 2 for any of the specified bolt alloys. The hardness of each bolt was measured using standard procedures. The calibration of the hardness tester was verified, using two calibration test blocks (54.1+1.0 HRC and 34.6+1.0 HRC). The measured hardnesses on the test blocks were found to be within the range of the hardness test block calibration. A section of the bolt was taken by cutting perpendicular to the long axis of the bolt, in the threaded region. The hardness of each bolt was determined by taking a hardness measurement on the end of the cut section. Results are shown in Table 7. Based on this investigation, it was concluded that:

Fig. 86

Micrograph of Knoop hardness (500 g load) survey of failed bolt (100 mm)

Table 6 Measured chemistry of the failed bolt A Element

C Si Mn P S Cr Mo Ni

0.42 0.32 0.82 0.010 0.014 0.49 0.15 0.45

Table 7 Measured hardness of submitted bolts Bolt

C B H G I Failed bolt A

Percent, wt%

Hardness, Rockwell C

43.4 36.4 41.6 40.4 39.2 36.9

Ultrasonic inspection does not have the sensitivity to detect cracks of the size and geometry seen in the bolts. The fractured bolt failed by fatigue, initiating at localized wear on the bolt shank. The soft surface hardness from partial decarburization aggravated the fatigue failure. The exposed cracks in bolts C, G, and I initiated at machining marks at the head-toshank fillet radius, and cracking propagated by fatigue. Evidence of machining at this radius was found. Bolt I had a second initiation site on the shank. The composition of the failed bolt did not meet the chemical requirements of S105 bolts. However, the chemistry of the bolts is not thought to have contributed to the failure of bolt A. The bolts were machined instead of cold formed. This results in lower fatigue strength and proof load.

REFERENCES

1. G.E. Dieter, Mechanical Metallurgy, McGrawHill, Inc., 1986 2. Metallography and Microstructures, Vol 9, Metals Handbook, 9th ed., American Society for Metals, 1985 3. D.A. Ryder et al., General Practice in Failure Analysis, Failure Analysis and Prevention,

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Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986 4. B.E. Wilde, Stress-Corrosion Cracking, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986 5. W.J. Jensen, Failures of Mechanical Fasteners, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986 6. Threaded Steel Fasteners, Failure Analysis and Prevention, Vol 11, Metals

Handbook, 9th ed., American Society for Metals, 1986 7. A.G. Glover et al., Failures of Weldments, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986 8. E. Alban, Failures of Gears, Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986 9. R.J. Parrington, Fractography of Metals and Plastics, Pract. Fail. Anal. Vol 2 (No. 5), 2002, p16

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Failure Analysis of Powder Metal Steel Components S. Ashok and Sundar Sriram, Sundram Fasteners Ltd.

POWDER METALLURGY (PM) technology provides a cost-effective method of producing near-net shape products, especially when a large number of the same or similar products are required. While the initial powder used is expensive compared to wrought steels, the manufacturing process can avoid machining steps and provide more uniform composition control. Total cost of producing PM parts can be less expensive than conventional metalshaping (casting, forging) in many cases. It was first adopted by the auto industry. Since this initial inception, other industries, such as the refrigerator, lawnmower, and hand tools industries, have also extensively adopted PM parts (Ref 1). Manufacture of powder metal parts involves pressing of metal powder in a die, sintering the green compact so that metallurgical bonding takes place, sizing or coining for densification and dimensional correction, followed by finishing operations. The finishing operations include machining, case and through hardening, steam treatment, and so on (Ref 1). This chapter reviews failure aspects of structural ferrous PM parts, which form the bulk of the PM industry. Focus is on conventional PM technology of parts in the density range of 6 to 7.2 g/cc. This constitutes 90% of the parts produced by PM technology. The PM part undergoes multiple heat treatments, some of which are unique to PM while others are similar to that used by conventional manufacturing processes, such as forging and casting. These include:

Powder annealing Sintering Case hardening Carbonitriding Case carburizing Through hardening Induction hardening Steam treatment

This chapter briefly introduces the processing steps involved in PM (Fig. 1). These heat treatments and the PM process steps are essential to understanding failure analysis of PM parts. The methods used for analyzing the failures are then discussed. Methods of failure analysis itself have recently become more systematic in terms of prevention and ensuring more robust process and high-quality products. Some case studies are given that illustrate different failures and the methods of prevention of these failures.

Powder Metallurgy Process Powder Production. Solid-state reduction of iron ore and atomization of molten iron are the most common methods of producing iron powders for structural parts. The powder thus produced should be favorable for compaction such that:

It fills the die cavity repeatably, which is measured in terms of the apparent density and flow rate of the powder. It can be compressed to the desired density with the lowest pressure, which is measured in terms of compressibility. It has sufficient strength at that density so that the green part can be handled, which is measured in terms of its green strength.

The filling characteristic of a powder is measured by its apparent density and its flow rate. The compressibility of a powder is a function of the particle morphology, processing history, and purity. Water-atomized iron powders can be compressed to higher densities compared to sponge iron powder. The higher the impurity content (typically oxides), the poorer the compressibility. As a process, all iron powders are generally annealed in a reducing atmosphere. Special high-compressibility

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Fig. 1

The powder metallurgy process

powders are produced by double-annealing operations. Blending. Alloying elements are typically mixed with the iron powder by blending in a double-cone blender. Elemental addition of alloying elements, that is, graphite, copper, or nickel, has its drawback, because these elements tend to segregate in the product as the alloying element content increases. A recent development in this regard is the introduction of diffusion-bonded iron powder and binder-treated premixes. These processes are proprietary. For example, the Distaloy process is diffusion bonding of nickel, molybdenum, and copper to iron powders developed by Hoganas AB, Sweden. The Ancorbond and Starmix methods are binder treated premixs to reduce graphitelubricant segregation. In the case of diffusionbonded iron powders, the alloying elements,

typically copper, nickel, and molybdenum, are mixed with iron, and the mix is annealed at a temperature of approximately 850 to 920 C. As a result, bonding of the alloying elements to the iron particle takes place, and this prevents segregation of alloying elements in the products. The same principle is used to prevent carbon-lubricant segregation. In this case, the binder, which is added during powder mixing, is sprayed on the iron-carbon mix. This binder ensures that each iron particle has a coating of graphite and thus minimizes segregation. Case studies in this chapter illustrate this mechanism. Compaction. The most widely used method of compaction is axial pressing of loose powder. Initially, the die cavity is filled with the loose powder from a hopper by gravity; the density is typically the apparent density of the iron powder, which is 2.4 to 3 g/cc. When compaction commences, that is, the punches move toward each other, densification takes place by particle rearrangement and plastic deformation. Thin sections of individual particles bend or break, cold welding and interlocking of neighboring grains takes place, and the voids are filled by the material, which becomes squeezed into them. The density of 5.5 g/cc is the lowest density of iron powder parts fit for practical handling in the green stage. On attaining this density, a rapid increase in pressure is observed when the part is compacted to densities between 6 and 7.2 g/cc. Compacting in the higher-density range gradually increases the rate of plastic flow of the metal. This results in the following:

Increased work hardening of powder particles Increased friction against the die walls

Structural ferrous powder metal parts typically have sections of varying thickness. With respect to heat treatment, it is very important that all sections in the parts have uniform density throughout. Low sectional densities result in through hardening and wide variation in the case depth within the PM part. Various methods of compaction have been developed to produce parts with minimal difference in density. These include:

Double-action compaction Spring floating die compaction Compaction by withdrawal process Multilevel compaction using multiple moving punches in special multiplaten presses

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Failure Analysis of Powder Metal Steel Components / 397

Sintering. During sintering of ferrous PM parts, the following reactions occur:

Metallurgical bonds are developed from the mechanical interlocks between the powder metal particles in the compact. Metal oxides in the powder compact are reduced by reaction with the carbon from the blended graphite or with the combustible constituents of the atmosphere. Desired final level of combined carbon in the microstructure is obtained by diffusion of carbon from graphite in the powder mix.

Densification of PM compacts can be achieved in sintering. The degree of densification depends on the sintering parameters and the alloys involved. Generally, higher sintering temperatures and longer sintering times promote densification of ferrous sintered parts. For conventional sintering, which generally is conducted at 1120 C, mesh belt conveyor furnaces are widely used. These furnaces can be used up to temperatures of 1150 C, primarily due to the temperature limitation of the belt and metallic muffle material. For stainless steels, proper sintering is critical for proper corrosion resistance (Ref 2).

Case Hardening Powder metallurgy steel parts can be case hardened by several processes, although various available processes are not equally suited to every application. A clear case/core relationship can be obtained only with parts having a density of at least 7.2 g/cc. Carburizing. Powder metallurgy parts with relatively low combined carbon contents of up to 0.20 wt% can be carburized by conventional pack or gas methods. Liquid carburizing is not recommended because of the difficulty of washing the parts free of salt. Gas carburizing is more practical for PM parts than pack carburizing. For this process to be successful, however, density as well as the precise composition of the parts should be known. Low-density parts should not be subjected to gas carburizing, because the carburizing gases penetrate the voids. Consequently, a distinct case is not achieved compared with the case developed on wrought parts under the same conditions. Instead, the carbon penetration in PM parts is generally deeper and relatively nonuniform. The extent of this condition varies

with density. In parts that have been repressed and resintered, this condition may be tolerable or even negligible, but for parts of lesser density, the depth of carbon penetration may be so great that these sections of quenched parts will be brittle. Another reason that conventional gas carburizing enjoys only limited use is because it does not increase hardenability. Thus, plain carbon grades usually must be quenched in an aqueous medium. This may result in cracking, especially if carbon penetration is excessive. As with wrought parts, the depth of carburized case of the PM parts depends on time and temperature. Carbonitriding is a modified form of carburizing. The principal process modification consists of introducing ammonia into the gas carburizing atmosphere, which results in the addition of nitrogen to the carburized case as it is produced. Nascent nitrogen forms at the workpiece surface by dissociation of the ammonia in the furnace atmosphere. Nitrogen diffuses into the steel surfaces simultaneously with the carbon, where the austenite stability is greatly enhanced by nitrogen in solution. This way, the necessary quenching rate to form martensite is reduced, and a martensitic microstructure is obtained without expensive alloying elements. Typically, carbonitriding of PM parts is carried out at 790 to 880 C for a duration of 30 to 60 min. Carbonitriding is widely used for case hardening of PM parts made of ferrous powders. Densities of the sintered compacts vary from approximately 6.8 to 7.9 g/cc. Parts may be infiltrated with copper prior to carbonitriding. Carbonitriding is extremely effective for case hardening high-density (7.2 g/cm3) parts made from sintered iron compacts (Fig. 2). Additionally, it is reasonably effective for case hardening parts of lower density. Equipment and Techniques. Procedures for carbonitriding PM parts are essentially the same as those used for similar wrought parts. Control of temperature and time is generally more critical than for wrought parts because of porosity. Lower temperatures are avoided to minimize the potential danger of explosion, and higher temperatures are avoided because case depth control is more difficult. The processing cycle, including composition of the atmosphere, is critical. The ammonia content (usually 1 to 5% of carrier gas by volume) increases hardenability and affects dimensional stability. Because dimensional changes in heat

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Fig. 2

Effect of density on the case depth as measured through a hardness traverse from the surface. Courtesy of P. Beiss, University of Aachen, Germany

treating are often crucial to the economic justification of producing parts by PM, gas composition, temperature, and quenching medium must be closely controlled. When processing to a new specification, establishment of processing parameters is usually made on small production quantities, thereby requiring sacrifice of only a few parts to arrive at optimal conditions. Such adjustments are then recorded so they can be used when processing the next lot of similar parts. Tempering of carbonitrided parts requires special consideration, largely because the quenching oil they contain will partially evaporate and pollute the environment. Toughness of PM parts in the hardened condition (either by oil quenching or through sinter hardening) is significantly improved on tempering. Induction Hardening. Induction heating is a method of heating electrically conductive materials by the application of a varying magnetic field whose lines of force are intersected by the workpiece. In this process, the varying magnetic field induces an electric potential, which in turn results in generation of electric current depending on the geometry, the frequency, and the electrical characteristics of the workpiece. The induced current, termed eddy current, generates heat that makes it amenable for use in many different heating applications, of which the induction hardening of steels and cast irons is one of the most predominant. Induction hardening of PM parts has several differences compared to hardening wrought steels and cast irons. Chemical composition, microstructural heterogeneity, and low density

due to the porosity are the causes of the different induction-hardening responses of PM parts. The results of hardening are more sensitive to chemical composition and the prior microstructure when compared with alternative processes. The electrical resistivity, thermal conductivity, and magnetic permeability strongly depend on the porosity of the PM part. Low density negatively affects the hardenability of powder metal parts (Table 1). It is recommended that the part selected for induction hardening has a density of at least 7 g/cc. Carbon, copper, nickel, and molybdenum are the most commonly used alloying elements in PM parts. The stresses due to high carbon content or large pores aggravate cracking. Segregation of alloying elements, foreign inclusions, or large pores can serve as stress raisers, making the powder metal part susceptible to cracking. It is quite common for PM parts to absorb oil. Thus, intensive ventilation is required, and one must ensure that the reused quenchant provides the required quench severity. Water-based polymer fluids are the most common quenchants used for induction hardening of PM parts. Steam Treatment. Sintered parts are subjected to steam treatment, wherein controlled oxidaton of the ferrous part is carried out in an atmosphere of superheated steam. As a result, a layer of Fe3O4 forms on the surface as well as in the pores. This oxide is hard and has excellent wear resistance (hardness of 450 HV0.05 kg), increases density, and improves corrosion resistance. The iron oxide also serves to seal the pores, which results in parts that are impermeable to gases. This characteristic of

Table 1 Effect of low density on properties and induction hardening process parameters Property

Thermal conductivity

Electrical resistivity Magnetic permeability

Structural homogeneity

Change

Effect on induction hardening parameters

Decreases with Inefficient conduction decrease in heat transfer density Larger temperature gradients(a) Increases with Larger current decrease in penetration depth density Decreases with Larger current decrease penetration depth in density Lower coil electrical efficiency Decreases with Wide scatter in decrease in apparent hardness density Wide scatter in case depth

(a) The penetration of the quenchant into open porosity overcompensates the effect of lower conductivity. As a result, low-density parts cool faster.

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steam-treated PM products has enabled its widespread use in pistons and valve plates for compressors. The sealing of pores also helps the PM part to be amenable to corrosion-preventive coatings (galvanic as well as barrier coatings) and, in some cases, to gas and salt bath nitriding. The formation of Fe3O4 is exothermic, and thus, a further increase in product temperature by 10 to 20 C is observed. Steam treatment is typically carried out at 500 to 550 C for a time of 30 to 300 min. The upper limit of 550 C is chosen so that the products do not exceed 570 C, which in turn would result in FeO formation (Fig. 3).

Failure Analysis Techniques Failure analysis techniques can be broadly classified as preventive methods or corrective methods. Preventive methods of failure analysis are typically carried out during product or process development. These refer to the standard techniques for failure mode and effects analysis (FMEA), which are specified in the various certifications, such as TS16949 and ISO

9000:2000. The engineer anticipates the failures of the product on the basis of known failure modes and designs the process and product parameters accordingly. However, in practice, failures are observed in spite of carrying out the measures that are recommended by an FMEA. This can be due to the fact that assumptions which go into carrying out the FMEA may change over a period of time, or the basis of the assumptions is incorrect, but more commonly, the failures occur due to the breakdown in one or more so-called 4M parameters (man, machine, method, material). Thus, there is a need for corrective methods of failure analysis, which form the bulk of this chapter. However, the techniques used in the FMEA are useful for analyzing failures, and a few are very relevant to heat treated products. In the case of heat treated powder metal parts, the failure modes and their causes are limited in number and can be quantified easily. The common failure modes of heat treated powder metal parts are:

Wear Fracture Dimensional instability because of plastic deformation Dimensional instability because of phase transformation Corrosion

These can be related to deficiencies in mechanical or metallurgical properties. Properties that play a role in failure of heat treated PM steel include:

Fig. 3

Fe-O-H2O diagram indicating the zones in which various oxides of iron are stable. Courtesy of P. Beiss, University of Aachen, Germany

Metallurgical properties: microstructure; case depth; coating adhesion, hardness, and thickness for wear-resistant coatings; salt spray life in the case of coatings for corrosion prevention; density; and chemical composition Mechanical properties: tensile strength, elongation, impact strength, fatigue strength, fracture toughness, and hardness

Deficiencies in properties that can result in a failure are caused by a combination of one or more of the 4M parameters. For example, low apparent hardness after carbonitriding may be associated with carbonitriding or due to low density in compaction. Thus, in analyzing the failure, one needs to consider the fact that the cause of the failure may be associated with more than one stage of component production.

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Failure Analysis Tools. The objective of the failure analysis is to relate the failure to a product property attributed to one or more 4M parameters of processing. The steps that are typically carried out in the failure analysis of heat treated powder metal parts are shown in Fig. 4. Examination of the product involves one or more of the following techniques:

Visual or any other method of nondestructive examination Fractography or examination under a lowmagnification stereomicroscope Hardness and related mechanical property measurements Metallography under an optical or scanning electron microscope Chemical analysis

The failure can occur because of a nonconforming property in the product even though the product meets the specifications. The cause of the former is related to the manufacturing process, and the latter is related to the design. However, both failures involve determining the 4M conditions that cause the failure and correcting the same. In the case of failure of a heat treated powder metal part, the failure mode as well as the causes of failure are limited. As mentioned previously, a principal activity is an operation-wide 4M analysis to determine the

Fig. 4

Steps in a failure analysis

cause of the failure. Two tools are invaluable aids in this regard. These are process maps and a cause-and-effect matrix. Process Maps. In charting a process map, the entire manufacturing process is considered. The implicit assumption is that the cause of failure can be related to not only the final operation but also to prior operations. Every operation in the manufacturing process is regarded as a process with inputs and outputs. The inputs consist of the 4M conditions and the output of the product requirements. All possible 4M conditions that can result in a nonconforming product are listed as inputs, with the nonconforming properties of the products as output. Table 2 is an example of a detailed process map for the PM process. This exercise is carried out for all stages in the manufacturing process. Once the process map is charted, the potential causes at different manufacturing steps are identified. Cause-and-Effect Matrix. The process map is an effective tool that eliminates the possibility of ignoring an operation that can result in failure of the product. However, information presented by the process map requires sorting for easy and meaningful interpretation. The cause-and-effect (CE) matrix ensures this. In the CE matrix, the 4M conditions that can cause the failure are listed in the y-axis and the defective properties

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Table 2 Detailed powder metallurgy process map Inputs

Compaction Single platen press No pressway No double action High apparent density Low apparent density Segregation of alloys Segregation of carbon and lubricant Sintering Low cooling rate High cooling rate Low temperature High temperature Low soaking time High soaking time Oxidizing atmosphere Decarburizing atmosphere Carburizing atmosphere High rapid burnoff temperature Low rapid burnoff temperature Low density Carbon-lubricant segregation Alloy element segregation Low sectional density Poor sinter hardenability Carburizing/carbonitriding Low carbonitriding temperature High carbonitriding temperature Low soaking time

Outputs

... Low density Density difference within compact exceeding 0.2 g/cc Carbon and lube segregation Alloy segregation Dimensions not conforming to specification ...

... Low hardness Decarburization Oxidation Excess growth Shrinkage Carburized microstructure Low mechanical properties Blistered product Pinholes Low sintered density ... ... ... ... ... ... ... High surface hardness

on the x-axis. These are usually transferred from the process map. A coordinate in this matrix links the potential cause, that is, the 4M conditions, to the failure, that is, the defective product characteristic. Once all the coordinates are filled, all the potential causes of a failure are listed in a user-friendly format. An example of a process map and CE matrix is illustrated in Fig. 5.

Inputs

Induction hardening Too high frequency Too low frequency High heating rate Low heating rate High heating time Delayed quench Severe quench High temperature Low temperature High carbon Low density Steam treatment Continous mesh belt furnace Batch furnace Low temperature Insufficient time Insufficient steam Oil in pores

Case Study 1: Wear after Sinter Hardening Sinter-hardened bushes were developed for an application that required high wear and impact fatigue resistance. The bushes were sinter hardened to a martensitic-bainitic microstructure using a prealloyed Fe-Cr-Mo powder. Accelerated rig testing for 24 h yielded no wear,

... ... Cracked product High case depth Low case depth Low surface hardness Localized zones of melting ... ... ... ... ... Low hardness Leaky product (pores not well sealed) Low oxide layer Poor surface appearance Loose rust Red rust

but the bushes were observed to fail in the field. The parameters that can cause the failure on listed in the CE matrix for wear of sinterhardened bushes, as follows: Operation

Powder chemistry Blending Compaction Sintering

Case Studies of PM Steel Failures

Outputs

Carburizing/carbonitriding (continued) High time of cabon ... potential attainment High soaking time Low surface hardness Oxidizing atmosphere Decarburization Decarburizing atmosphere Oxidation Carburizing atmosphere Excess growth High nitrogen potential Shrinkage Masking of products ... Delayed quench Carbide network in case Hot oil quench Nonmartensitic transformation product in case High-viscosity quench oil ... Delayed tempering Low mechanical properties Poor hardenability High retained austenite Low density Low case depth Carbon segregation High case depth Alloy segregation High core hardness Low sectional density Low core hardness

4M parameter

Observations

Low carbon

Yes

Segregation of carbon and alloy elements Low density Low sintering temperature Low soaking time Oxidizing atmosphere Decarburizing atmosphere High rapid burnoff temperature Bainite in microstructure Overloading of parts

No No No No No No No Yes No

Figures 6(a) and (b) show the failed bush. Figure 7 shows the microstructure of the failed bush, which consists of martensite and upper and


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pg 402

Blending

X X

X X X

X X

Poor mechanical properties Nonconforming dimensions

Blistered product

Pinholes

Poor surface apearance

Loose rust

Red rust

Thick oxide layer

Thin oxide layer

Melt zones

Cracked product

Decarburization

Oxidation

Free ferrite in core

Carbide network in case Nonmartensitic transformation product

Low core hardness High retained austenite

High core hardness

High case depth

Low case depth

Process inputs Excess time Insufficient time Elemental alloy additions High carbon High lubricant

High apparent hardness

Process step

Low apparent hardness


X X X X

X

Compaction Single platen press No pressway No double action High apparent density Low apparent density Segregation of alloys Segregation of carbon and lubricant

Sintering

Low cooling rate High cooling rate Low temperature High temperature Low soaking time High soaking t ime Oxidizing atmosphere Decarburizing atmosphere Carburizing atmosphere High rapid burnoff temperature Low rapid burnoff temperature

Low d ensity Carbon-lubricant segregation Alloy element segregation Low sectional d ensity Poor s inter hardenability

Carburizing/ carbonitriding

Fig. 5

X X X X

X X X X

X

X X X

X

X X X X

X X X

X

X

X

X

X X

X X X X

X

X

X X

X X

X

X X

X X

Low carbonitriding X temperature High carbonitriding temperature X Low soaking time High time of carbon potential attainment High soaking t ime X Oxidizing atmosphere Decarburizing atmosphere X Carburizing atmosphere X High nitrogen potential X Masking of products X Delayed quench X Hot oil quench X High-viscosity quench oil Delayed t empering

X

X X X

X

X

X

X

X

X X X

X

X X X X X

X XX

X X

X

X X

Cause-and-effect matrix derived from powder metallurgy process map

X X

X


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Carburizing/ carbonitriding

Induction hardening

Steam t reatment

Too high frequency Too low frequency High heating r ate Low heating rate High heating time Delayed quench Severe quench High t emperature Low t emperature High carbon Low density Continous mesh belt furnace Batch furnace High t emperature Low temperature Insufficient time Insufficient steam Oil in pores

X X

X

X

Poor mechanical properties Nonconforming dimensions

Blistered product

Pinholes

Poor surface apearance

Loose rust

Red rust

Thick oxide layer

Thin oxide layer

Melt zones

Cracked product

Decarburization

Oxidation

Free ferrite in core

Carbide network in case Nonmartensitic transformation product

Low core hardness High retained austenite

High core hardness

High case depth

Low case depth

Process inputs Poor hardenability Low density Carbon segregation Alloy segregation Low sectional density

High apparent hardness

P ro c e s s s te p

Low apparent hardness


X

X X X X

X

X X X X X X X

X

X

X

X X X X

X

X X X

X

X

X X X X X

X X X

X X X

X X X

X

Fig. 5 (continued)

Fig. 6

Sintered bush. (a) Outside diameter wear. (b) Outside diameter crack

lower bainite. The failure was due to the softer bainite being worn out, resulting in eventual accelerated fatigue failure of the bush. Bainite is an essential component for toughness. A fully martensitic microstructure obtained by increasing the cooling rate and the carbon content has poor toughness. Thus, the decision was made to use a different material that can give the required hardness and toughness.

Corrective Measures. An improved chemistry was derived, where the base iron powder was prealloyed with nickel and molybdenum, and the carbon content was increased to 0.9%. This resulted in a predominantly martensitic microstructure. The presence of nickel ensured the toughness requirements, and the increased carbon ensured the wear requirement. Figure 8 shows the new microstructure, wherein martensite with some retained austenite is observed.


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Fig. 7

Micrographs of the failed bush. (a) Martensite and lower bainite. (b) Upper bainite

Results. The improved bushes were assembled in chains and tested in a rig as well as in the field. Figure 9 shows the comparative evaluation of the PM and wrought chains after the change in material of the PM bush. It is observed that in addition to withstanding the tests, the chains with PM bushes have less elongation compared to chains with wrought bushes. Case Study 2: Pinholes after Sintering Valve seats have a chemistry that is highly alloyed and a carbon content greater than 1%. Because most of the elements are admixed, segregations of these elements are not uncommon. The parameters that can cause pinholes in valve seats one: Operation

Blending

Compaction Sintering Machining

4M parameter

Segregation of carbon and lubricant Copper agglomeration No inspection for agglomerates in powder Not applicable Low sintering temperature Low soaking time Pullout of nonmetallic inclusions in machining

Fig. 8

Martensite and some retained austenite in the modified chemistry

Observations

Yes No Yes ... No No No

In all of these above cases, the pinholes that are observed after sintering or machining are in the regions where the agglomerates were present, prior to sintering or machining. The copper melts and diffuses into the iron matrix, carbon diffuses in the matrix, and the lubricant burns off, leaving the pinhole in each case. In the case of nonmetallic inclusions, they are pulled out during machining, thus leaving a pinhole.

Scanning electron microscopy/electron dispersive x-ray analysis was carried out to determine the cause. Figures 10 to 13 reveal high carbon content near the pinhole. Thus, the likely cause of the pinholes was concluded to be carbon-lubricant agglomeration during blending. Corrective Measures. Binder-treated premixes that minimize carbon-lubricant segregation were recommended for the products. An inspection procedure was evolved to check the premix for the presence of agglomerates and coarse particles, wherein the premix was sieved and the +150 mm fraction was inspected for agglomerates under a stereomicroscope. This ensured that a premix with agglomerates is not issued for compacting the products.

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Fig. 9

Comparison of sintered versus regular bush in 12B model chain

Fig. 10

Fig. 11 Pinhole in the valve seat at a magnification of 40 : 1

Results. The rejections due to pinholes decreased significantly as a result of using binder-treated premix.

SEM micrograph of the pinhole of Fig.10 indicating a suspected graphite particle

The parameters that can cause blistered sintered products are: Operation

Powder chemistry Blending

Case Study 3: Blistered Sintered Products Synchronizer keys used in an automobile gear box were produced with Fe-3%Ni-0.5%C. Elemental nickel and carbon were mixed with the iron powder. Blistering of the products after sintering was observed in ~5% of the products.

Compaction Sintering

4M parameter

Observations

High nickel content Segregation of nickel powder Elemental nickel additions Not applicable Low rapid burnoff temperature H2 in hot zone

Yes No Yes ... No Yes

It was observed that delubing the keys was not completed in rapid burnoff, and the lubricant


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Fig. 12

Region of suspected graphite agglomeration in the pinhole (Fig. 11) at higher magnification

Fig. 13

High carbon in electron-dispersive x-ray analysis, confirming the likely cause of the pinhole (Fig. 10–12) as carbon-lubricant segregation in blending

was carried into the hot zone, where the atmosphere consists of N2-10%H2. In the presence of H2, nickel acts as a catalyst favoring the cracking of the lubricant. This results in the lubricant exploding, and consequently, blistered products. Corrective Measures. Two possible solutions to the problem were considered:

Ensure complete removal of the lubricant prior to contact with H2 Reduce the catalytic activity of nickel

Due to the limitations of sintering in a continous mesh belt furnace, the former could not be

completely ensured. Thus, it was resolved that reduction in catalytic activity of nickel could be an effective solution. Sulfur effectively poisons nickel and limits its catalytic properties, so 0.03% S was added to the blend. Results. The addition of sulfur to the blend effectively reduced the rejections from 5% to nil. Case Study 4: Dimensional Instability during Shrink Fitting Exhaust valve seats with tool steel powder as a major constituent were shrink-fitted in a

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cylinder head of a sport utility vehicle. It was observed that the valve seats were expanding during shrink-fitting instead of contracting. This resulted in interference between the valve seat and the cylinder head, causing rejection of the cylinder heads. The parameters that can cause dimensional instability during shrinkfitting are: Operation

Powder chemistry Blending Compaction Sintering Tempering

4M parameter

Observations

Not applicable Not applicable Not applicable Not applicable High retained austenite

... ... ... ... Yes

A significant amount of retained austenite was observed even after tempering. This retained austenite was transforming to martensite during shrink-fitting, where the products were cooled to 60 C, which resulted in a volume expansion. Corrective Measures. A complete transformation of retained austenite can be effected by:

Multiple tempering Cooling to below the martensite finish temperature (in this case, a deep cryogenic treatment) Combination of tempering and deep cryogenic treatment

Trials revealed that a combination of subzero treatment and tempering gave the best results with respect to dimensional stability as well as metallurgical properties. Results. The valve seats subjected to a combination of subzero treatment and tempering did not expand during shrink-fitting. Further dimensional measurements taken at 60 C were in line with the theoretically calculated shrinkage in contrast to the earlier expansion.

parameters that can cause wear of a free-graphite bush are: Operation


Oil impregnation

4M parameter

Observations

Not applicable Not applicable High density High sintering temperature High sintering time Cementite network Insufficient free graphite Low oil content

... ... No Yes No Yes Yes No

Metallography of the worn sample revealed a cementite network in the product (Fig. 14). This is a clear indication that excess carbon had gone into solution, which resulted in the solid lubricant being depleted. Corrective Measures. Sintering was carried out at 1080 C for 20 min, which resulted in excess combined carbon. The temperature was reduced so that the combined carbon was 50.8%. The sintering temperature for the product was fixed at between 1040 and 1060 C Results. Bushes sintered at 1060 C for 20 min resulted in a pearlitic-steaditic microstructure with no cementite network (Fig. 15).

Case Study 6: Fracture of Steam-Treated Part Valve plates for compressors were subjected to grinding after steam treatment. Fracture of a steam-treated valve plate was observed during grinding after steam treatment. The parameters

Case Study 5: Wear after Sintering Bushes were produced with Fe-0.45%P– 2%Cu-2.5%C and supplied in the oil-impregnated condition. These bushes were deliberately undersintered so that some free graphite was retained, which acted as a solid lubricant that in turn minimized wear in application. The

Fig. 14

Pearlite-steadite microstructure with cementite network


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Fig. 15

Pearlite-steadite microstructure free from cementite network

that can cause the fracture of a steam-treated part are: Operation


Steam treatment

4M parameter

Observations

Not applicable Not applicable Low density Low sintering temperature Low soaking time Oxidizing atmosphere Decarburizing atmosphere High dewaxing temperature Overloading of parts High temperature Insufficient steam Low soaking time Overloading of parts Masking of parts

... ... No No No No No No Yes No No No No No

were subjected to standard metallographic analysis. The product with low hardness revealed a predominantly ferritic structure with free copper (Fig. 16), even though the blend graphite addition was 0.8% (Fig. 17). Thus, undersintering was suspected to be the possible cause. However, the sintering control charts revealed that the temperature as well as the belt speed were meeting the specifications. The loading pattern was then checked. It was observed that the loading pattern followed that which was recommended for a 46 cm (18 in.) belt, whereas the belt width was 30 cm (12 in.). This resulted in overloading of parts, which consequently resulted in undersintering. Corrective Measures. Suitable one-point lessons were imparted to the sintering operators to ensure that the right process plans were followed. Case Study 7: Oxidation after Sintering Exhaust valve seats for internal combustion engines are produced with a chromium-rich hard-phase alloy. The hard-phase alloy is essential for wear resistance at elevated temperatures, to which the exhaust valve seat is subjected. Oxidation of the hard phase was observed after sintering. The parameters that can cause oxidation after sintering are: Operation

Hardness measurements revealed very low hardness. Product with hardness conforming to specification and product with low hardness

Fig. 16

Low-hardness sample with almost no combined carbon. Only ferrite and free copper are observed.


Fig. 17

4M parameter

Observations

Chromium-rich alloy Not applicable Rusted compacts High rapid burnoff temperature High cooling zone dewpoint

Yes ... No Yes Yes

Typical microstructure with ~0.7% combined carbon


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Chromium-rich particles are susceptible to oxidation when the product is heated and cooled in the range of 650 to 1050 C. Oxidation of the hard-phase particles is favored by:

High rapid burnoff temperature High cooling zone dewpoint

Figure 18 shows the microstructure with oxidized hard-phase particles. The oxide film around the hard phase does not allow the hard phase to interact with the matrix. As a result, the matrix will not have chromium carbides, which helps the valve seats to retain wear resistance at high temperatures. The rapid burnoff temperature was observed to be 650 C, and the cooling zone dewpoint was between 20 to 30 C. These are potential causes for oxidation. Corrective Measures. The furnace was allowed to stabilize for a period of 4 to 16 h. The gas ratio was maintained at 70% N2 and 30% H2 during this period, and the gas volume was increased by ~25%. The rapid burnoff temperature was restricted below 600 C. The valve seats were sintered when the operating dewpoint in the hot and cooling zones is 35 C or lower. Results. The modification in sintering parameters resulted in the production of oxidationfree valve seats (Fig. 19).

density of 7 g/cc. Subsequently, the gear was sintered, and a sizing operation was then carried out for dimensional correction. Induction hardening of the gear was carried out after sizing, and cracks were observed after induction hardening. The parameters that can cause cracks after induction hardening are: Operation

Powder chemistry Blending Compaction Sintering Sizing Induction hardening

4M parameter

Observations

High carbon Not applicable Low density Low sintering temperature Low sintering time Not applicable High heating time Severe quench

Yes ... No No No ... No No

Figure 20 shows the crack. This crack originates in the middle of the root of the gear and

Case Study 8: Cracks after Induction Hardening A PM transmission gear was subjected to induction hardening for an application that demanded high wear resistance as well as high contact and bending strength. The gear had an Fe-Cu-C chemistry and was compacted to a

Fig. 18

Exhaust valve seat with oxidized hard-phase particles

Fig. 19

No oxide layer observed in the hard-phase particles after implementation of corrective measures

Fig. 20

Crack in the middle of the root of the primary driven gear


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does not originate at any sharp corners. Hardness profiles from the root indicated that the root has been hardened to a very high hardness (Fig. 21), compared to the earlier sample in which no crack was observed. Microstructure near the root revealed untempered martensite, which gradually became a mixture of martensite and bainite and subsequently ferrite and pearlite. This is in contrast to the earlier sample where the root was not hardened and had a bainitic structure. Secondly, the combined carbon estimated from the pearlite content in the sintered microstructure was 0.7% compared to 0.6% in the product that was successfully induction hardened. The strain caused by the martensitic transformation increases in proportion to the combined carbon in the microstructure. Thus, the possible causes for the crack were:

Too high a soaking time, which resulted in a severe quench High combined carbon in the sintered microstructure

Corrective Measures. The formation of martensite in the root was a result of quenching from very high temperatures. The time was reduced, and the coil design and quench were modified. However, the martensitic microstructure could not be avoided in the root, and cracking was still observed. It was concluded that the product was unable to withstand the stresses associated with the transformation. The

Fig. 21

Hardness traverse from root

carbon content was reduced in the blend to 0.6%, which resulted in a lesser volumetric strain due to martensitic transformation in the product. Induction hardening with this chemistry yielded no cracks, and the hardness traverse matched that obtained in the rapid prototype samples. Results. Products with a lower combined carbon content did not crack when induction hardened. Case Study 9: Cracks after Quenching A flyweight for a governor assembly in a diesel engine was carbonitrided to meet the product requirements for strength. This is a complex part wherein multiple punches are used to form the sections of different thicknesses. The product was observed to crack after carbonitriding. The parameters that can cause cracks after quenching are: Operation

Powder chemistry Blending Compaction Sintering Carbonitiriding

4M parameter

Observations

Not applicable Not applicable High density difference Underfilling of powder Low sintering temperature Low sintering time High hardening temperature Severe quench

... ... No Yes No No No No

The parts were subjected to magnetic particle inspection prior to heat treatment. No cracks were observed. Metallography was then carried


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out at the location where the crack occurred in heat treatment. This is shown in Fig. 22. Insufficient powder filling was observed (Fig. 23). The underfill results in a localized region of low strength, which gives way during subsequent quenching and results in crack formation. Corrective Measures. The localized underfilling was due to the presence of a sharp corner in the tool. A trial was conducted wherein the sharp corner was smoothed by filing the parts. No crack was observed after carbonitriding. It was resolved to increase the radius in the tool so that a sharp corner and consequently an underfill are avoided. Products thus produced were heat treated, and no cracks were observed. Results Products produced with an improved radius in the tool resulted in elimination of product rejections due to quench cracks. Case Study 10: Dimensional Change in Carbonitriding An angle lever for a diesel injection pump was produced with FD-02-00 chemistry and carbonitrided to a case depth of 0.2 to 1.2 mm. When the carbonitriding was established at a new heat treatment source, it was observed that the bore did not shrink sufficiently during carbonitriding. The parameters that could cause a dimensional change in carbonitriding are: Operation

4M parameter


Not applicable Not applicable Bore undersized due to tool wear High sintering temperature High sintering time Low soaking time Insufficient carbon and nitrogen enrichment High quench oil temperature Insufficient transformation

Observations

... ... No No No Yes No

Carbonitriding was carried out at 840 C for 45 min with a carbon potential of 0.8% in the atmosphere. In the case of the vendor, better furnace gas sealing ensured that this case depth was achieved in 20 min. The hardness traverse indicated that the longer soaking time had a higher carbon martensite in the case as well as a higher core hardness, which in turn resulted in a case depth at the lower end of the specification for a lower soaking time. This is shown in Fig. 24. Corrective Measures. Higher shrinkage is ensured by a higher amount of martensitic transformation. This was ensured by increasing the soaking time from 20 min to 45 min. The increased soaking time resulted in better exposure of the product to carbon and nitrogen enrichment, which in turn resulted in higher martensite content and the desired level of dimensional shrinkage. Results. The increased martensite content resulted in increased volume expansion. Thus, the shrinkage in the bore is higher with a soaking time of 45 min. Products soaked at 840 C for 45 min were found to shrink sufficiently, and dimensions conformed to specifications.

Case Study 11: Low Surface Hardness after Carbonitriding Low surface hardness was observed in an iron-copper part after carbonitriding. The

No Yes

Fig. 23 Fig. 22

Quench crack in flyweight

Metallography of the region where the crack occurred in the sintered sample, showing a pre-existing defect caused by poor powder filling

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Fig. 24

Hardness traverse in bell crank lever

parameters that can cause low surface hardness after carbonitriding are: Operation


4M parameter

Not applicable Not applicable Low density Low sintering temperature Low sintering time Low soaking time Insufficient carbon and nitrogen enrichment Excess carbon and nitrogen enrichment Delayed quench Pearlite/bainite in case High retained austenite

Observations

... ... No No No No No Yes No No Yes

After ensuring that the density conformed to specification, the low-hardness products were subjected to metallographic analysis. The case microstructure revealed a high amount of retained austenite, which is the cause of low hardness (Fig. 25). High retained austenite can result if the carbon potential as well as the nitrogen potential are high. The oxyprobe, which senses the carbon potential, was examined and found to be covered with soot. This resulted in improper sensing of carbon potential, leading to an excess of active carbon in the furnace atmosphere, which in turn resulted in high retained austenite. Corrective Measures. The temporary corrective measure adopted to salvage the parts was to carry out a deep freeze, because reprocessing may result in scrapping of the product

due to dimensional distortion. On a permanent basis, a periodic preventive maintainence schedule for the oxyprobe was drawn up and adhered to. Results. Correcting the oxyprobe yielded products with a martensitic/510% retained austenite structure that met the requirements for surface hardness. Case Study 12: Variation in Bore Diameter after Heat Treatment A valve retainer for a shock absorber was produced from an FC-02-08 blend that was hardened for wear resistance. High ovality was observed after heat treatment. This resulted in ~10% of the products being out of specification.

Fig. 25

High retained austenite in the case

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The parameters that can cause a wide variation in bore diameter after heat treatment are: Operation


4M parameter

Observations

Excess copper content Elemental additions Low density Low sintering temperature Low sintering time Volume expansion due to martensitic transformation Thermal contraction during quenching

No No No No No Yes Yes

The ovality was checked after sintering and after heat treatment. The data revealed that hardening contributes significantly to distortion, and the contribution of sintering is negligible. Thermal contraction due to quenching and the accompanying martensitic phase transformation contributed significantly to the distortion. Corrective Measures. To minimize the distortion, it was resolved to eliminate the hardening operation and achieve the required hardness in sintering. Prealloyed powder containing nickel and molybdenum, blended specially to produce products with negligible distortion after sintering, was adopted to achieve both the hardness and dimensions after sintering. Results. Changing the process from conventional hardening to sinter hardening reduced the dimensional spread to 50% of the variation observed in products with a separate heat treatment operation. Case Study 13: Low Breaking Load after Carbonitriding A component for a gearbox application produced with a chemistry of FD-02-00 and having an overall density of 7.2 g/cc was subjected to carbonitriding to meet the product requirements for wear and mechanical strength. During assembly, the product failed by fracture. The parameters that can cause low breaking load after carbonitriding are: Operation


4M parameter

Not applicable Not applicable Low density Low sectional density Low sintering temperature Low sintering time High carbonitriding temperature High time of carbon potential attainment Variation in batch quantity Martensite in core No case depth inspection in thin section

Low breaking load indicated poor toughness of the product. Poor toughness is related to both microstructure and density. Lower densities and a through-hardened microstructure with martensite in the core can cause premature fracture. Metallography was carried out in different sections of the failed product. Metallography of the sections indicated that the thinner sections in the product were through hardened. Figure 26 shows the through-hardened microstructure. Sectional density measurements indicated that the thicker sections had a density of 47.2 g/cc, and the thinner sections had a density of 7.1 g/cc. However, this through hardening of thinner sections was not observed in the initial samples that were tested. To determine the cause of this variation across different heat treated batches, the carbon potential attainment was studied for various batches. During establishment of the carbonitriding process, the process parameters were fixed for ~500 parts. The same was subsequently fixed for the full batch quantity of 3500. In this particular case, the carbon potential pickup in the furnace was completely different for the trial and bulk. Figure 27 shows the difference in carbon potential pickup. There is a significant difference in the time for attainment of a carbon potential of 0.8%. Longer times between 0.6 and 0.8% in the bulk lots result in additional carbon pickup, which resulted in through hardening of the thinner sections. Corrective Measures. The carbonitriding temperature and time were reduced. The batch quantity was fixed. The density in the component was increased to 7.2 g/cc min in all

Observations

... ... No Yes No No Yes Yes Yes Yes Yes

Fig. 26

Martensite in the core of thin sections in the product

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Fig. 27

Difference in carbon potential attainment for change in batch quantity

sections, which ensured closure of interconnected pores. Inspection standards for sectional density and case depth measurements in the thinner section were introduced so that any nonconformance would be detected and not reach the customer. Results. The corrective measures resulted in a pearlitic-bainitic microstructure in the core of the thin section, and the toughness and, consequently, the breaking load were achieved consistently. Case Study 14: Dimensional Change in Steam Treatment Pistons having a density of 7 g/cc were steam treated for wear resistance and corrosion resistance. During steam treatment, it was important that the bore dimensions be maintained within specification. It was observed that after steam treatment, the bore was undersized. The parameters that can cause dimensional change in steam treatment are: Operation

Powder chemistry Blending Compaction Sintering Machining Steam treatment

4M parameter

Observations

Not applicable Not applicable Not applicable Not applicable Bore undersized High temperature Thick oxide layer

... ... ... ... No Yes Yes

After ensuring that the machined dimensions of the bore were in accordance with specification

prior to steam treatment, metallography was carried out on the samples with the undersized bore. An oxide layer of up to 15 mm was observed (Fig. 28). Typically, Fe3O4 layers are 3 to 5 mm, because the lattice of Fe3O4 does not allow further diffusion of oxygen molecules once this layer is formed. Such a thick layer is due to the formation of FeO. The steam treatment set temperature was 560 C for this product. Because the Fe3O4 reaction is exothermic, the temperature of the products exceeded 570 C, resulting in the formation of FeO. Corrective Measures. The steam treatment temperature was subsequently reduced to 530 C, where an oxide layer of 4 to 6 mm was obtained (Fig. 29). This resulted in the bore dimensions meeting the specification.

Fig. 28 560 C

Thick oxide layer indicative of FeO formation that was obtained at a steam treatment temperature of

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pg 415


Fig. 29 530 C

Oxide layer of 6 mm, indicative of Fe3O4 obtained on reduction of the steam treatment temperature to

Case Study 15: Low Core Hardness after Steam Treatment Crankshaft sprockets with an Fe-Cu-C chemistry and a density of 6.6 to 6.8 g/cc were supplied in the steam-treated condition. The customer demanded that the hardness in the surface and core meet the specification of 140 HV10 min. The parameters that can cause low core hardness after steam treatment are: Operation


Sizing Steam treatment

4M parameter

Observations

Low carbon Not applicable High density Low sintering temperature Low sintering time High ferrite content Closed surface pores Oil in pores Insufficient time Low temperature High temperature (4570 C) Masking of products Nonuniform steam circulation

Yes ... No No No Yes No Yes Yes Yes No Yes No

It was observed that within a batch, some products had a core hardness greater than 140 HV10. Low and high core hardness samples were subjected to metallography. It was observed that higher-hardness samples had better steam penetration compared to low-hardness samples. Corrective Measures. To ensure higher core hardness, the products required greater exposure to steam. This was achieved by changing the steam treatment temperature and time from 530 C/30 min to 550 C/2 h. Masking of products can also potentially prevent uniform exposure to steam. To eliminate masking, the

loading pattern was changed from a double layer to a single layer. However, even after increasing the steam exposure time and improving the loading pattern, variation of hardness was observed within the same batch. A trial was conducted wherein oil-free sintered samples were steam treated with sized and oiled parts. It was observed that the sintered samples have uniform steam penetration, and the thickness of the steam oxide layer is uniform throughout the part. This is in contrast to the sized products with low core hardness, wherein the oxide layer is minimal. The presence of oil or a related organic residue was suspected of interfering with the steam penetration. An oil that has a higher volatile content and a lower organic residue content was used for rust prevention, and the same was used as a lubricant in sizing. The change of oil resulted in a core hardness conforming to specification. However, the hardness obtained ranged from 145 to 190 HV. The presence of products with hardness marginally above 140 HV10 indicated that the hardness must exceed 160 HV to avoid the chances of core hardness going below specification. This additional improvement in hardness was achieved by having a completely pearlitic microstructure instead of a ferrite+pearlite microstructure, which was effected by a blend addition of 0.8% C. Results. Implementation of the corrective measures resulted in the core hardness being achieved consisently in accordance with specification, as follows: Properties

Specification

Surface hardness Core hardness

140 HV10 min 140 HV10 min Sintered ... microstructure Steam-treated ... microstructure

Before

After

152–208

185–235

110–173

170–207

Ferrite+ 100% pearlite pearlite Complete steam Incomplete penetration with steam oxide layer of penetration with low-oxide 4–6 mm in core layer52 mm in core

REFERENCES

1. Asm Handbook, Volume 7, Powder Metal Technologies and Applications, 1998 2. E. Klar and P. Samel, Powder Metallurgy Stainless Steels: Processing, MicroStructures, and Properties, ASM International, 2007


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Induction Hardening Janez Grum, University of Ljubljana

INDUCTION HEATING, in most applications, is used to selectively heat only a portion of the workpiece that requires treatment. This usually means that the process can be accomplished in a relatively short time and with high efficiency, because energy is applied to the workpiece only where it is needed. Induction surface hardening is applicable to axisymmetric or near-axisymmetric machine parts in steel or cast iron that are produced in substantial volumes. There are two basic techniques for induction hardening machine parts: single-shot and scanning. The former employs selective heating and quenching to harden a specific area or areas of the machine part in one operation. The latter is usually applied to harden progressively long, continuous sections, such as shafts and spindles. In this instance, the scanning inductor traverses the length of the section, heating only a relatively small area at any given time, and is followed closely by the quench arrangement, which is often an integral part of the inductor. These advantages make it possible for induction hardening to be fully automated and are especially suitable for a large series of workpieces. The induction-hardening procedure enables an engineer, by simply adapting the shape of the induction coil, to ensure the desired shape of the hardened profile of the surface layer. Likewise, the engineer can surface harden only that part of the surface (local hardening) on which a certain increased level of hardness and wear resistance are desired. One of the main advantages of induction hardening is the ability to harden a surface layer only in certain places at a defined penetration depth and shape. For dynamically loaded machine parts, it is very important to ensure the total compressive stresses in the thin, most-loaded surface layer. The total stresses are a sum of residual stresses in a machine part and of load stresses produced by

the action of external forces and moments. To ensure a long life of the machine part, knowledge of the residual stresses in the machine part and how to adjust the size and distribution of the residual stresses by means of the selection of an appropriate production technology are very important. In surface hardening, compressive residual stresses always occur in the thin surface layer due to martensite transformation. The size and variation of the residual stresses depend primarily on carbon content and less on the type and content of alloying elements in heat treatment and surface-hardening steels. The variation of residual stresses in the surface layer can be modified by varying the induction-heating conditions and by a quenching method. Induction surface hardening creates very desirable residual stresses in the hardened surface layer. Residual stresses are always of a compressive nature and are usually present to the depth of the induction-hardened layer. Residual internal stresses, that is, the so-called residual stresses, are the stresses present in a material or a workpiece when there is no external force and/or external moment acting on it. The residual stresses in metallic machine parts have attracted the attention of technicians and engineers only after manufacturing processes improved to the level at which the accuracy of the manufacture exceeded the size of deformation, that is, distortion, of a workpiece/product. Thus, it was almost 150 years ago that the effect of internal stresses on plasticizing, that is, destruction, distortion, and plastic deformation, of workpieces was already known. It was then that experts introduced measurement of individual dimensions of products. For a given type of machining process, they connected the influence of the selected machining conditions with the size of dimensional deviations. This was also the beginning of an expert approach to the selection of the most suitable machining and/or heat

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treating conditions based on the criterion of minimum dimensional deviations, that is, minimum workpiece distortion. Currently, measurement of individual workpiece dimension is a very practical, uncomplicated, and reliable method of product quality assessment. Through the appropriate selection of the grinding wheel and grinding conditions and taking into account the physical and mechanical properties of the workpiece material, very favorable compressive residual stresses in the hardened surface layer can be retained. How is it possible to assure a desirable surface and surface layer quality after induction hardening and fine grinding? Finding an answer to this question requires a very good knowledge of the grinding process on the microlevel as well as knowledge of mechanical and heat effects acting on the layer of the workpiece, including the type of grinding wheel and the grinding conditions. An all-inclusive consideration of the numerous influences of the tool type and condition on the changes on the surface and in the surface layer of the workpiece in the given machining conditions is described by the term surface integrity. This is a scientific discipline providing an integral assessment of the surface and subsurface layer. It was defined at the beginning of the 1960s. For high-quality machine parts and parts subjected to heavy thermomechanical loads, different levels of description of the surface integrity were defined. A basic level of the surface-integrity description includes measurement of roughness and analysis of the microstructure and microhardness in the thin surface layer resulting from the machining process under given machining conditions. The second level of the surface-integrity description includes studies of residual stresses in the surface layer and of mechanical properties of the given material. The third level of the surface-integrity description includes tests making clear the behavior of the given part under the operating conditions. As essential advantage of induction surface hardening is that it is possible to achieve a sufficient repeatability of the hardened layer thickness on the workpiece as well as a desirable or even prescribed hardened layer profile, ensuring sufficient hardness and favorable distribution of residual stresses in the hardened layer. A variety of steels and a whole range of inductionhardening methods provide the possibilities for very accurate planning of residual stress size and distribution. This is of growing importance,

since manufacturers are frequently required to produce machine parts that, among other surface properties, must have quite specific residualstress distribution along the depth of the hardened layer. It has become a proven fact that high compressive stresses ensure high fatigue strength of machine components and reduce the danger of crack occurrence and growth on the surface of components. In thermal hardening, the surfaces of suitable materials, usually plain carbon, low-alloy steels, or cast irons, are austenitized and then quenched to produce a hard martensitic case that is usually tempered in a subsequent operation. Case depths are normally in the range of 0.5 to 5 mm. Case hardness is typically approximately 700 HV on hardening and 600 to 650 HV after tempering at 200 C (Ref 1). Heating processes include electrical induction and resistance (Ref 2–4), and direct impingement methods using flames, (Ref 5), lasers (Ref 6–11), and electron beams (Ref 12). Of these, induction heating is the most widely used. Laser and electron beam heating have recently become established in a number of applications, mainly where distortion was a problem with the induction method. Thermal treatments are mainly employed when only local areas on machine parts require hardening and wear resistance. Consequently, such methods are energy efficient compared with thermochemical methods where bulk heating of batches of components is undertaken. Normally, the surface-hardening process also introduces compressive stresses into the surface layers, leading to an improvement in fatigue properties. For example, the drive shafts of heavy lorries and buses are induction surface hardened to improve their fatigue properties (Ref 10, 11). The skills and experience required for manual flame hardening have been largely superseded by automated flame techniques or by the induction method. However, for one-off machine parts, the simplicity of manual flame hardening and the extremely low capital investment ensure that the method is still used on a regular basis (Ref 5, 13). From a heat treatment point of view, laser can be considered as a versatile and flexible highintensity heat source that can operate in air. It is capable of undertaking a range of processes, essentially simultaneously, since the laser beam can be directed through air by metal mirrors and switched and shared among a number of work

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pg 419

Induction Hardening / 419

stations. Manipulative techniques using mirrors allow the beam to be directed to areas not accessible by other techniques, for example, the bores of tubes. Set against these considerations are the high capital cost and low energy efficiency of the technique (Ref 1, 6, 10).

Steels for Surface Hardening Heat treatable steels contain between 0.2 and 0.6% C and can be carbon steels or low-alloyed steels (Table 1) (Ref 14, 15). In practice, engineers choose these steels because, in the soft condition (normalized or soft

annealed), they are more suitable for machining, while their strength properties and hardness can be subsequently refined by additional heat treatment. Soft steels have good machinability, while additional heat treatment improves their mechanical properties to a desirable level. Figure 1 presents a schematic of the heat treatment methods used to achieve the desired mechanical properties (Ref 15, 16). Besides the mentioned surface hardening, other thermochemical methods can be used to create a desirable wear resistance. The steels for surface hardening are from the group of heat treatable steels, except that they usually contain between 0.35 and 0.5% C. Heat treatable steels can be

Table 1 Suitable steels for induction surface hardening Chemical composition % Type of steel

1 2 3 4 5 6 7 8 9 10 11

C

Si

Mn

P, max

S, max

Cr

Mo

Ni

0.33–0.39 0.38–0.44 0.43–0.49 0.48–0.55 0.50–0.57 0.42–0.48 0.34–0.40 0.38–0.44 0.38–0.44 0.38–0.44 0.37–0.43

0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40 0.15–0.40

0.50–0.80 0.50–0.80 0.50–0.80 0.60–0.90 0.40–0.70 0.50–0.80 0.60–0.90 0.60–0.90 0.50–0.80 0.70–1.00 0.50–0.80

0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035

0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035 0.035

... ... ... ... ... 0.40–0.60 0.90–1.20 0.90–1.20 0.90–1.20 0.40–0.60 0.60–0.90

... ... ... ... ... ... ... ... 0.15–0.30 0.15–0.30 0.15–0.30

... ... ... ... ... ... ... ... ... 0.40–0.70 0.70–1.00

Source: Ref 14, 15

Fig. 1

Heat treatment methods for carbon and low-alloyed steels. Source: Ref 15


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supplied in soft or refined conditions. Carbon steels with less than 0.35% C are in the normalized condition, and the products made from them can be in the soft or refined condition. For improvement of the wear resistance of these steels, additional surface hardening can be applied (Ref 14, 17, 18). Below 0.35% C, the treated surface is insufficiently hard, and above 0.6% C, there is a danger of surface cracking after quenching. To obtain a satisfactory hardening response in induction surface hardening, it is necessary for the carbides to dissolve when the steel is heated. Thus, steels that have been quenched and tempered and have small carbide particles respond better than steels with large spheroidized carbides. Similarly, the rate of solution of carbides in steels other than plain carbon and low-alloy steel is too slow for them to be satisfactorily induction hardened. Steels must also have sufficient hardenability to achieve the required hardness at the specified depth (Ref 17). The procedures are similar for fabrication of machine parts made of carbon steels with more than 0.35% C and alloyed steels, expect that because of the higher machinability required, they must be in the soft-annealed condition. For induction surface hardening, it is recommended to use steels with an appropriate carbon content (0.35 to 0.45%) and a careful selection of alloyed elements. Usually, these steels have an increased silicon (0.15 to 0.40%) and manganese (0.50 to 1.00%) content and other alloying elements, such as chromium, molybdenum, and nickel. More highly alloyed tool steels (O1, D2, D3, A1, and S1) and some martensitic stainless steels (AISI 416, 420, and 440C) are also sometimes induction hardened. Among alloyed steels, chromium, chromiummolybdenum, and chromium-molybdenumnickel steels prevail. With given combinations of alloying elements in induction-hardened machine parts, the internal stresses in heating as well as in cooling can be controlled. The steels are normally quenched in water. In certain cases, the alloy steels can be cooled by means of an oil emulsion. Steels can also be oil quenched. Having selected the right shape of the product and the right choice of technology, it is possible to expect only minimal distortions of the workpiece after completion of the heat treatment (Ref 14). If the starting points are the hardness and residual-stress profiles, then the induction-heating conditions should be adapted to the selected steel. Induction heating of a machine-part

surface to the austenitic zone above TA3 is of major importance, since not only the case depth but also the right through-depth residual-stress and microhardness profiles of the machine part are to be provided. The through-depth microhardness profile of the surface layer depends, with the given steel, on the induction-hardening conditions. Induction heating may be controlled by infrared thermometers. An infrared thermometer is placed close to the induction coil, which means that only the maximum surface temperature may be measured. This technique is quite simple and practical. It allows a user to determine the required, that is, optimal, power density for heating. When a multiturn coil is used, for example, with longer workpieces, several infrared optical-fiber thermometers may be employed. Thus, with sufficient spacing between two adjoining turns, the temperature may be measured at different locations on the machine part. A difficulty may arise due to nonuniform heating along the coil movement. By collecting and processing the gathered measurement data, the process quality may be assessed. In scientific research, temperature measurement with thermocouples may be applied too. In induction surface heating, temperature measurement using thermocouples is quite exacting, since they are embedded in the machine part. The hardened state of steel is to be obtained at an appropriate depth. The thermocouples embedded at a particular location in the workpiece permit temperature measurements at this location only. In progressive hardening, however, the momentary temperature is measured. Thus, the so-called temperature cycles in heating, as well as in cooling, are obtained. The temperature cycles make it possible to predict the efficiency of surface hardening. The choice of energy input, that is, power density, in an appropriate nomogram, that is, the choice of power and heating time with reference to the shape of the induction coil and case depth, applies only to stationary heating and hardening of machine parts. This means that in progressive hardening the data obtained on the energy input are of informative character only.

Main Features of Induction Heating Induction heat treatment is a segment of the much larger technical field of induction heating, which combines many other industrial processes using the phenomenon of heating by induction (Ref 1, 16, 18–22).

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Induction heating is often one of the most effective heat treatment processes available for a variety of applications, including:

Surface hardening Through hardening Tempering and stress relief Annealing and normalizing Grain refinement Precipitation hardening or aging Sintering of powdered metals

In most of these applications, induction heating is used to selectively heat only the portion of the workpiece that requires treatment (Ref 2). This usually means that the process can be accomplished in a relatively short time and with high efficiency, because energy is applied to the workpiece only where it is needed. One of the main features of induction heating compared with conventional heating procedures is that heat is generated in the workpiece itself. In conventional heating procedures, the heat input achieved is only 5 to 200 kJ/m2s energy, whereas in induction heating this energy input is 300 MJ/m2s. In induction heating, heat penetrates into the workpiece by the aid of

Fig. 2

high-frequency alternating current, the choice of frequency depending on heating requirements. Induction heating power supplies are frequency changers that convert the available utility line frequency power to the desired single-phase power at the frequency required by the induction heating process. They are often referred to as converters, inverters, or oscillators, but they are generally a combination of these. The converter portion of the power supply converts the line frequency alternating current input to direct current, and the inverter or oscillator portion changes the direct current to single-phase alternating current of the required heating frequency. Many different power supply types and models are available to meet the heating requirements of a nearly endless variety of induction heating applications (Ref 21, 22). The specific application will dictate the frequency, power level (Fig. 2), and other inductor parameters such as coil voltage, current, and power factor (cos Q) or Q factor (Ref 15, 19). Advantages in Surface Hardening of Machine Components. Induction hardening is most often used for surface hardening of machine components and has the following

Typical power-frequency regions of induction heat treatment applications. Source: Ref 15, 19

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advantages over other procedures (Ref 2, 19–22):

Relatively short heating times Heating procedure is not strictly governed by hardening temperature. All that matters is that the heating process does not end at too low a temperature, because sometimes it is necessary for the transformation into austenite. The heating temperature is limited by the solidus-line temperature, since the process is carried out while the material is in the solid condition. Due to a short heating time, there is no danger that the austenite grains would grow at higher austenitization temperatures, which also means there is no danger of coarse and brittle martensite formation. The quenching procedure is easy to perform, contributing to short surface-hardening times. In progressive induction hardening, a spray coil is located directly below the inductor and quenches the heated surface. In single-shot hardening, the inductor is designed to perform the function of heating as well as quenching. The coil around the workpiece functions as an inductor in the heating phase. After the austenitization temperature has been reached, the current is interrupted, and the coil functions as a spray for quenching. Induction hardening is a short procedure that does not require any additional protection against oxidation. Thus, compared with other similar procedures, such as cementation, it does not require much subsequent machining. Due to the nature of the procedure, the workpieces, especially if symmetrically shaped, are less susceptible to undesirable deformations after induction hardening. The volume changes in the workpieces after hardening the surface layer can be very accurately predicted or estimated. The volume changes after induction hardening of thin layers are so small that quite often the function of the machine component is not affected. Especially in induction hardening of thin layers and workpieces with low mass, it is possible to achieve the desired critical rate of cooling by self-cooling in air alone, that is, by heat conduction from the heated surface layer into the remaining cold part of the workpiece. With thicker layers and workpieces of greater mass, it is necessary to use

quenching agents that move the actual cooling rates close to the critical cooling rate. These requirements can be met with the right selection of quenching oils or polymer water solutions. Practical experience has shown that polymer water solutions are very suitable for quenching of induction-heated surfaces, since optimal quenching can be ensured with the right choice of concentration of the polymer water solution. The induction-hardening procedure enables the engineer, by simply adapting the shape of the induction coil, to ensure the desired shape of the hardened profile of the surface layer. Likewise, the engineer can surface harden only that part of the surface (local hardening) on which a certain increased level of hardness and wear resistance is desired. One of the main advantages of induction hardening is the possibility of hardening a surface layer only in certain places, at a defined penetration depth and shape. These advantages make it possible for induction hardening to be fully automated and are especially suitable for a large series of workpieces. Induction hardening always leaves compressive residual stresses in the surface layer, which makes machine components more resistant to dynamic loads. Compressive residual stresses in the surface layer after induction hardening prevent the occurrence of cracks in dynamically loaded components and prevent the growth of existing cracks on the workpiece surface, if these are present due to hardening or grinding. Induction hardening is appropriate for smallsized workpieces, since, by a well-chosen technology of heating and cooling or quenching, a hardened surface layer and a refined core can be ensured. Thus, the required wear resistance of the machine component at a certain location as well as the required loadbearing capacity of the component can be created, experiencing only a slight loss in toughness of the core.

Induction Hardening of Machine Parts Working Coil and Procedure during Induction Hardening. Methods of induction hardening of thin surface layers are always adapted to the product size and shape and the

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requirements of location, size, and shape of the hardened layers profile. Designing an induction coil, however, is very specific to a given application and can be considered the most important aspect of the engineering of the system. There are several major functions that an induction coil must perform to make a job successful:

Induce current in the load so that the proper heating pattern is obtained Accomplish the proper heating pattern with as great an efficiency as possible Provide an impedance match to the generator so that adequate power can be transferred to the load Have a geometry that will accomplish the aforementioned three major functions and permit easy loading and unloading of the part being heated

First, the outer and inner surface on machine components must be distinguished. The size and shape of the induction coil must be adapted to the size and shape of the product, and during the heat treatment process, the distance should be adapted to the workpiece. It is important that the design of the machine component be planned for induction-hardening specifics. A uniform gap size between the induction coil and the workpiece surface is very important, since the energy penetration depends on it. This is why singleshot hardening, where the workpiece and the coil are at a standstill, is rarely used. To ensure a uniform gap, single-shot hardening is almost

Fig. 3

always combined with a rotating motion of the workpiece, or, on long workpieces, progressive hardening is used. Figure 3 shows different types of coils for induction hardening designed for external surfaces (Fig. 3a), internal surfaces (Fig. 3b), and front surfaces (Fig. 3c) of the workpiece. In all three cases, the shape of the induction coil is adapted to the size and shape of the workpiece surface that is to be hardened (Ref 2, 15, 16, 19–22). Workpiece sizes depend on the manufacturing possibilities of high-frequency coils and their types. This restricts the use of induction hardening primarily to small-sized machine components. Induction hardening of internal surfaces requires a great deal of knowledge and experience in the design and manufacturing of induction coils of small diameters. The smallest internal diameter of a machine component that is to be induction hardened depends on the manufacturing possibilities of making small-diameter coils, the capacity of the high-frequency generator (power, frequency), and the positioning accuracy of the induction coil inside the workpiece. The induction coils are made of materials with the highest possible electrical conductivity. Consequently, copper or silver is usually chosen to produce the induction coils. The induction coils are adapted to the shape of the machine part at the location where heating, that is, surface hardening, is performed. The choice of the size and shape of the induction coil depends on the

Significant types of coils for induction heating. Source: Ref 15, 16

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Fig. 4

Influence of shape and arrangement of induction coil around machine part on heating, that is, hardened-pattern profile

method of heating, that is, hardening. With the single-shot heating method, the induction coil may be adapted by choosing an adequate pipe cross section and by adapting the number of turns used so as to provide simultaneous heating of the entire area at the machine part. Figure 4 shows the influence of the induction coil shape on the heating profile (Fig. 4a) and that of the distance between the individual turns of the induction coil and the gap on the heating profile (Fig. 4b). The concentration of the heat generated in the machine part heated by the induction coil is symmetrically and uniformly distributed across the cross section, so that no distortion of the part may occur after quenching. The symmetry of the hardened-pattern profile thus reduces the risk of distortion of the machine part. The induction coils are made of pure copper, which can be nicely shaped as required to produce coils. Because the electrical conductivity of the material is reduced due to plastic deformation of the material, it is necessary to adequately anneal the coils to soften the material. Thus, the highest possible electrical conductivity is ensured. It is also necessary to prevent the formation of copper oxide at the coil surface to preserve highest possible electrical conductivity of the material. This will provide a high efficiency of induction heating. In single-shot induction heating, that is, induction surface hardening, it is important to provide the same gap size all around the cylindrical machine part. Because such an arrangement can be very exacting, the heating process is usually performed with the rotation of the machine part. Figure 5(a) shows single-shot hardening performed with a different number of

Fig. 5

(a) Influence of multiturn coil on hardened-pattern profile. (b) Influence of shape and gap size of induction coil turns on hardened-pattern profile. (c) Influence of number of induction coil turns on hardened-pattern depth

inductor turns along the specified hardened pattern. This means different divisions among the individual coil turns. Under the same heating conditions, considerable difference in the hardened-pattern depth will be obtained with the same heating time. Figure 5(b) shows different examples of shaping a multiple-turn induction coil around the cylindrical machine part. The figure indicates that the gap size along the cylindrical part changes, which produces different hardenedpattern profiles. Figure 5(b) shows the influence

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of a multiple-turn coil with too-large divisions of the individual coil turns; therefore, the same heating depth, that is, hardened-pattern depth, along the coil length is not achieved. Thus, individual coil areas showing a smaller hardened-pattern depth are obtained at the machine part. Figure 5(c) shows the influence of the cross section of a single-turn coil on the hardenedpattern profile. For hardening of axles, shafts, and similar lengthy machine parts, the scan-hardening process with a single-turn coil of an adapted cross section is chosen. Such a hardening process is necessary, otherwise, high-energy densities and very long multiple-turn induction coils would be needed. Figure 6 shows a diagram of heating, that is, through hardening, of a cylindrical rod of different diameters (D, in centimeters) with different inductor powers (P, in kilowatts), and a given frequency ( f, kilohertz) and inductor movement (v, in cm/s). The diagram shows a linear dependence between the connected power at the coil and the depth of hardening, which means that the choice of a higher power produces hardening at a greater depth, the movement velocity being the same. It can also be established which diameters of the axles or shafts can be through hardened if the same power is available and different movement velocities are chosen. It is possible to distinguish between:

Single-shot hardening Scanning or progressive hardening Tooth-by-tooth hardening or gap-by-gap hardening

Fig. 6

Quenching techniques are an important design feature of induction-hardening equipment. The important questions to be answered when determining quenching systems include (Ref 2, 17, 23, 24):

Workpiece size and geometry Hardenability of steel Type of austenitizing operation (surface or through hardening) Type of heating method (single-shot or scanning) Type of quenchant

The two most common types of systems consist of spray quench rings and immersion techniques. When quench rings are used for round bars, their shape, like the coil, is generally round (Ref 2, 23, 25). Figure 7 shows different ways to single-shot harden the surface layer. Common to all of them is that heating is performed along the length of the layer, so that the induction coil embraces the particular part of the workpiece where heat treatment is to be performed (Ref 15, 16). The ring may be located concentric with the coil (Fig. 7c) or directly underneath or alongside it (Fig. 7b), as in the single-shot induction hardening setup. In any of these ways, it is important to ensure a constant gap size between the induction coil and the workpiece surface. This is usually achieved by a rotating motion of the workpiece. In the first case, shown in Fig. 7(a), the cooled induction coil embraces the object at a certain height. After heating to the austenitization temperature is

Selection of heating conditions in scan hardening to provide through hardening of a cylindrical rod

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completed, the workpiece is quenched in a quenching bath, or quenching is done by moving the workpiece into the quenching spray located beneath the induction coil (Fig. 7b). Figure 7(c) shows a method of heating and cooling in the induction coil, which is made so that it enables heating of the workpiece surface for a certain period, and then, on switching the current off, the same coil is activated as a spray for quenching. The induction coil is usually made of copper tubing with additions of alloying elements that do not affect the electrical conductivity of copper but contribute to good machinability of the material. Good machinability is required because it is necessary to make the openings for the quenching spray. Thus, during the entire

Fig. 7

heat treatment, the workpiece is rotating, and in this way, uniform heating and quenching are ensured. The heating and quenching conditions created in this way produce a uniform hardened surface layer and ensure repeatability of thickness and shape of the hardened profile. In induction scanning or progressive hardening, workpieces move through the quench ring and coil, with quenching occurring immediately after heating (Fig. 8). For nonsymmetric workpieces, the quenching system, like the coil, is generally the same shape as the workpiece (Ref 15, 16). The second method of induction hardening involves a group of procedures of progressive or progressive-rotating hardening. This type of

Single-shot induction surface hardening of a cylindrical workpiece. Source: Ref 15, 16

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hardening is employed when the required length of the hardened surface is longer than the length (size) of the coil. In this case, the correct heating mechanism can be ensured by the optimal induction coil design (number of coil turns, length of the coil) and suitable energy inputs for a given size and profile of the hardened layer (Ref 23, 25, 26). To ensure the uniformity and repeatability of the hardened layer, progressive motion of the induction coil must be provided, or the workpiece is fed progressively through the coil while rotation of the workpiece takes place at the same time. This method of induction hardening is presented in Fig. 8. It is possible to design the induction coil separately from the spray (Fig. 8a) or combined so that the induction coil has downward-directed spray openings for quenching in its bottom part (Fig. 8b). Gear wheels belong to very demanding machine elements, from the point of view of mechanical machining as well as heat treatment. In hardening smaller gears, that is, gears with small modules or small diameters, the method of single-shot hardening can be used to ensure a varying thickness of the hardened layer (Ref 15, 16). The thickness is maximum on the tooth tip and then decreases toward the tooth root.

Fig. 8

Figure 9(a) shows single-shot hardening of tooth tips, where the coil encircles the whole gear. This method is very simple and suitable for highfrequency heating of gears with a module up to 3 mm and for smaller gear diameter. In this case, it must be ensured that the tooth flanks are hardened at least along the length that is otherwise activated in the mesh, since the thickness of the hardened layer varies, being highest at the tooth tip and gradually decreasing toward the tooth root. This method of hardening is suitable only for gears subjected to low loads, just to increase the wear resistance of the gear (Ref 16, 23, 26, 27). Figure 9(b) shows induction surface hardening, where the whole tooth and a certain area below the tooth root are hardened. This method of induction surface hardening is appropriate for gears with a module up to 5 mm. Heating is achieved by two systems: first, with a current of medium frequency, and then in the final phase for a relatively short period by heating with a high-frequency current. By quenching, a desirable hardness of the tooth flank surface and increased strength of the gear tooth can be achieved. In the root of the tooth, due to a refined microstructure, an increased material fatigue

Progressive induction hardening of a cylindrical workpiece. Source: Ref 16

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strength can be noted. Gears heat treated in this way are suitable for highly loaded machine parts (Ref 28–30). Figure 9(c) shows hardening of individual gear teeth, where the coil encircles the tooth and the individual teeth are heated in turn and quenched directly after heating. In the case of large-module gears, simultaneous heating of the left and right flank of the same gear tooth can be performed (Fig. 9d). Induction hardening of all gear teeth (Fig. 10) can be done only for gears with a module smaller than 3 mm and a gear diameter up to 100 mm. This kind of hardening requires high-power frequency generators, for example, 200 kW. For gears exposed to higher loads, where refinement of the whole tooth is desirable, a power of 30 to 40 kW and longer heating times can be selected. To ensure a desired penetration depth and profile of the induction-hardened layer, the following heat treatment parameters can be varied in a single-shot or progressive induction hardening by choosing a suitable power density and feed rate:

The size and shape of the induction coil adapted to the workpiece Kind of steel and its thermal properties Size and mass of the workpiece on the location where induction hardening is to be performed Quenching agent and method of quenching

Fig. 9

A number of graphs and nomographs are available for this purpose, offering a selection of heat treatment conditions for heating and cooling. The most important data are the starting point data on energy input and frequency of the current, and the temperature to achieve in induction hardening. From these data, the time necessary for heating, in the case of single-shot hardening, or the feed rate of the workpiece or the rate at which the coil should move along the workpiece can be defined (Ref 2, 18–20, 23, 27, 31, 32). Figure 11(a) illustrates induction surface heating with a medium-and high-frequency current. The procedure is known as doublefrequency heating (Ref 27). Here, the gear is first placed into the coil fed by the mediumfrequency current. Then, the gear is moved into the high-frequency coil, where only the surface layer of the gear tooth is reheated with a highfrequency current. When both phases are completed, the gear is dropped or moved into a quenching bath. In this method of heating with a double frequency, the progressive motion of the workpiece can be combined with additional rotation. Additional rotating motion ensures a uniform reheating of the surface layer and results in uniform microstructural changes on the left and right side of the tooth. In this way, undesirable and nonuniform dimensional deviations between the left and right tooth flanks are prevented. The same method of induction surface

Different methods of induction surface hardening of gear wheels. Source: Ref 15, 16

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heating is shown in Fig. 11(b), where the same coil is used for heating with a medium- and highfrequency current (Ref 27). The second method is hardening of individual gear wheels or individual gear gaps. Figure 12 shows induction heating of a gear wheel toothby-tooth and Fig. 13 shows induction heating with the gap-by-gap method. In both methods, the tooth surface is first heated with the induction coil and then quenched with a specially adapted spray system. Figure 14 shows the shape and position of the massive inductor placed around the gear tooth (Ref 27). The coil is shaped so that the gap between the coil and the tooth surface varies. Only in this way is it possible to ensure a uniform thickness of the layer in the middle and on the edge of the tooth (Ref 33–37). Difficulties with this method of gear wheel hardening occur when the tooth gaps are too small and the coil heats up the adjacent flank. This method of induction surface heating of gear teeth is often not appropriate, since after heating the adjacent tooth, the next step does not ensure

Fig. 10

the desirable hardness of the adjacent tooth. To avoid this, special protection made of thin copper sheets is used to prevent heating of the adjacent tooth flanks. The conditions are presented in Fig. 15(a–c) (Ref 27). In addition to protecting adjacent tooth flanks, the shape of the coil around the tooth has other effects. By changing the shape and position of the coil around a particular tooth, it is possible to achieve equal hardened layer profiles, as in the case of single-shot hardening. The thickness of the hardened layer is at a maximum at the tooth tip and then gradually decreases toward the tooth root. In Fig. 15(c), the coil is placed slightly lower, heating only the tooth flank along the entire height from the root to the tip. Figure 15(b) shows the lowest position of the coil while still reaching below the tip of the tooth. The coil positioned in this way does not heat the tip of the tooth but only the surface of the tooth flank from the tip to the root and yields maximum thickness of the hardened layer in the middle of the tooth. In this case, too, small gears are heated with a high-frequency current,

Single-shot surface induction hardening of gear wheel tooth tips

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Fig. 11

Fig. 12

Induction heating system for gear wheel heating by double frequency (medium/high). Source: Ref 27

Tooth-by-tooth induction surface heating process of installing the induction coil around the gear wheel tooth for heating, followed by shifting spray quenching

Fig. 13

Gap-by-gap gear wheel induction surface heating by moving the induction coil for heating and moving or installing spray quenching

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whereas large-sized gears are first heated with a medium-frequency current, and then heating is performed with a high-frequency current (Ref 2, 16, 18, 19, 27, 32). The third method of induction surface hardening is appropriate for large gear modules and is known as tooth-gap hardening, which belongs to the progressive hardening methods. In this case, the coil is placed so that it ensures a uniform gap between the coil and the flanks of two adjacent teeth. The tooth-gap hardening method is very demanding and requires much

Fig. 14

Relative position of the massive induction coil to the gear wheel tooth at induction heating. Source: Ref 27

Fig. 15

Influence of induction coil height on profile of induction heating surface layer in an individual tooth. Source: Ref 27

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experience and knowledge to achieve the desirable properties of the gear. This method is also known as contour hardening. It is an ideal method for heat treatment of gears, because it increases the hardness on the tooth surface only slightly while decreasing the load-bearing capacity in the root of the tooth. Gears heat treated in this way exhibit very good behavior in operation, because compressive residual stresses are present in the root of the tooth. Gears with induction-hardened flanks, given that the dimensioning is carefully carried out, can achieve the highest fatigue strength. To verify the results of induction surface hardening, it is necessary to take certain measures for controlling the quality of the hardened layer. For this purpose, hardness and microhardness measurements, supported by microstructural analysis, are commonly used. A disadvantage of this procedure is that, due to the method of heating and quenching (nonuniformly overheated left and right tooth flank), slightly higher dimensional deviations may be obtained than in the case of simultaneous hardening of both flanks of the same tooth (Fig. 15d) (Ref 27). Metallurgical Aspect of Induction Surface Heating. Prior to transformation hardening, an operator should calculate the processing parameters for the given power system. The procedure is as follows. Some of the processing parameters are chosen, some calculated. The choice is usually left to the operator and his experience. Optimization is then based only on the selection of power density and scan speed. The correctly set parameters of transformation hardening ensure the right heating rate, heating to the right austenitizing temperature, TA3 , and a sufficient austenitizing time, tA. Consequently, with regard to the specified depth of the hardened layer, a temperature a little higher than the transition temperature, TA3 , should be ensured. Because of a very high heating rate, the equilibrium diagram of, for example, steel, is not suitable; therefore, it is necessary to correct the existing quench temperature with reference to the heating rate. Thus, with higher heating rates, a higher austenite transformation temperature should be ensured in accordance with a timetemperature-austenitizing (TTA) diagram. The diagram in Fig. 16(a) is such a TTA diagram for 1053 steel in the quenched and tempered state, whereas Fig. 16(b) is for the same steel in the normalized state (Ref 38). Because the steel shows a pearlitic-ferritic microstructure, a sufficiently long time should

be ensured to permit austenitizing. In fast heating, austenitizing can be accomplished only by heating the surface and subsurface to an elevated temperature. For example, with a heating time, t, of 1 s, for total homogenizing, a maximum surface temperature, Ts, of 880 C should be ensured in the first example and a much higher surface temperature, 1050 C, in the second example. This indicates that approximately 170 C higher surface temperature, DTs, should be ensured in the second example (normalized state) than in the first example (quenched and tempered state). Figure 17 shows a space TTA diagram including numerous carbon steels with different carbon contents. The TTA diagram gives particular emphasis to the characteristic steels, that is, 1015, 1035, 1045, and 1070 steels, and their variations of the transition temperature, TA3 , with reference to the given heating rate and the corresponding heating time (Ref 38). Such a temperature difference ensures, with regard to the heating and cooling conditions of the specimen, the time required for austenite homogenizing, tA, in the given depth. Figure 18 shows a shift of the transformation temperature, which ensures the formation of inhomogeneous and homogeneous austenite within the selected interaction times (Ref 39). A shorter interaction time will result in a slightly higher transformation temperature, TA1 , and also a higher transformation temperature,

Fig. 16

Time-temperature-austenitizing diagram for steel 1053 in various states. Source: Ref 38

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TA3 . To ensure the formation of homogeneous austenite with shorter interaction times, considerably higher temperatures are required.

Fig. 17 Ref 38

Fig. 18

Influence of heating rate and carbon content on austenitic transformation temperature. Source:

Figure 18(a) shows a temperature-time diagram for austenitizing of Ck 45 steel. The isohardnesses obtained at different interaction times in heating to the maximum temperature ensure that partial or complete homogenizing of austenite is plotted. Figure 18(b) shows the same temperature-time diagram for austenitizing of 100Cr6 hypereutectoid alloyed steel. The diagram indicates that with short interaction times, which in laser hardening vary between 0.1 and 1.0 s, homogeneous austenite cannot be obtained; therefore, the microstructure consists of austenite and undissolved carbides of alloying elements that produce a relatively high hardness, even up to 920 HV0.2. After common quenching of this alloyed steel at a temperature of homogeneous austenite, a considerably lower hardness, only 750 HV0.2, but a relatively high content of retained austenite were obtained. Retained austenite is unwanted, since it will produce unfavorable residual stresses and reduce wear resistance of such a material.

Temperature-time-austenitizing diagrams with lines of resulting hardness for various steels. Source: Ref 39

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The distribution of residual stresses in heat treatment procedures where only the surface of the workpiece is heated (induction hardening, flame hardening) differs greatly from the procedures where heating is performed throughout the entire volume (nitriding, cementation). In nitriding and cementation, the aforementioned second layer in the subsurface does not appear at all, because the direction of the heat flow is opposite to the direction of the heat flow in induction and flame hardening. The resultant operating tensile stresses on the surface or in the surface layer can thus be considerably smaller. Due to the surface hardness, induction and flame hardening lowers the fatigue strength of machine components; therefore, care should be taken to diminish all detrimental effects in the surface layer. A typical example of induction surface hardening is surface hardening of gears that are heated with a low heating rate and relatively low current frequency. The outer hardened zone includes almost the entire height of the gear teeth, whereas the second zone is in the tooth root area. A gear heat treated in this way will meet the wear-resistance requirements expected of the gear tooth, while the strength of the other part of the tooth is of minor importance. The fatigue strength in this case will be relatively low due to high tensile residual stresses in the tooth root, that is, in the second zone where the operating or load tensions and the tensile residual stresses are summed up. Figure 19 shows another example of an induction-hardened gear where the frequency of the current was so high that the gear tooth is heated along the flank surface and tip, as is the case in cementation. The energy input in heating a gear tooth or the whole gear was such

Fig. 19

that the second zone has not shown up. A similar heat treatment can be applied to the spline inside the gear. A gear heat treated in this way is more resistant to wear and corrosion and should have high resistance to fatigue in bending because of a smaller thickness of the layer in the second zone. Many induction-hardened gears are treated in the tooth gap, that is, in simultaneous heating and subsequent quenching of two adjacent flanks of the left and right tooth of the gear with a hardened root area. In view of the variety of methods to induction harden gears and the other possible ways of gear hardening, it is unwise to make a hasty decision in selecting the procedure. It is necessary to make a thorough analysis, including answers about the expected quality of the hardened layer and analysis of the operation loads of the machine components. Highly loaded gears can be successfully induction hardened if a high-frequency current and high-input power are used. High frequency is necessary in order to obtain a sufficient thickness of the hardened layer on the tooth flanks and a fine hardened layer in the tooth root area. A high-input power is necessary to increase the heat gradient, which makes the size of the second zone smaller, resulting in a thinner second layer with tensile residual stresses. On gears with a small diameter, the inductionhardening equipment should be able to harden the gear throughout its volume. Induction hardening of entire gears has some advantages, such as rapid heating, no danger of decarbonization and oxidation, high productivity, and repeatability in the gear quality. Thus, the hardened layer is only defined by the first zone to the depth that is greater than the height of the gear tooth.

Typical examples of induction surface hardening of (a) carbon steel and (b) alloyed steel gears produced from carbon steel (a) and alloyed steel (b)

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In induction hardening and quenching, lowering or even the disappearance of compressive residual stresses is achieved in the tooth root area, causing a considerable decrease in fatigue strength and a higher propensity for fatigue. In flame heating, lower temperature gradients are reached than in induction heating. This results in increased thickness of the layer in the second zone. Possible harmful effects due to the disappearance of the second zone of the hardening layer can be avoided by heating the entire gear, which is possible to do in commonly used furnaces. When the induction coil has stopped heating, an austenitic microstructure in the surface layer should be obtained. Then, the cooling process for the austenitic layer begins. To accomplish martensite transformation, it is necessary to ensure a critical cooling rate that depends on the material composition. Figure 20 shows a continuous cooling transformation diagram for EN19B steel, including the cooling curves (Ref 38). Because carbon steels have different carbon contents, their microstructures also show different contents of pearlite and ferrite. An increased carbon content in steel decreases the temperature of the beginning of martensite transformation, TMS as well as of its finish, TMF

Fig. 20

Figure 21 shows the dependence between carbon content and the two martensite transformations (Ref 38). Consequently, an increase in carbon content in steel results in the selection of a lower critical cooling rate. In general, the microstructures formed in the surface layer after transformation hardening can be divided into three zones:

A zone with completely martensitic microstructure A semi-martensitic zone or transition microstructure A quenched and tempered or annealed zone with reference to the initial state of steel

Transformation hardening of steel starts from its initial microstructure, which is ferriticpearlitic, pearlitic-ferritic. or pearlitic. In steel heating, transformation into a homogeneous austenitic microstructure should be ensured. Figure 22 shows the dependence of the maximum surface temperature obtained in induction heating with a machine part made of steel with 0.8% C and a pearlitic microstructure (Ref 40, 41). How the through-depth heating of the machine part will proceed depends on the maximum surface temperature obtained and the power density, which means that the throughdepth hardness profiles will differ.

Continuous cooling transformation diagram of EN19B steel. Source: Ref 38

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With the same power density, this means that a higher maximum surface temperature will be accomplished with a longer heating time. Since the heating times are usually short, austenitic

grains have little time left to grow. It can be assumed that the grains remain fine and do not affect through-hardenability, that is, the throughdepth hardness profile. The diagram shows the heating conditions that provide, in all cases, a completely homogeneous austenitic microstructure to a certain depth and, consequently, constant hardness. Then follows a transition zone consisting of homogeneous austenite and some inhomogeneous austenite and pearlite. Consequently, hardness in the transition zone will gradually decrease to that of the parent metal. It will be approximately 240 HK. It is important that the transition zone consists of a mixture showing different ratios of the microstructures concerned. The different microstructure ratios in the transition zone, however, define the hardness profile in this zone. Thus, at the maximum surface temperature, a maximum hardness of approximately 700 HK, without a constant part with homogeneous martensite, is obtained, and then it decreases immediately to the hardness of the parent metal. With a maximum temperature of 800 C, there will be constant hardness to a depth of 0.6 mm; from that point to a depth of 0.85 mm, the hardness will slowly decrease to that of the parent metal, that is, 240 HK. The highest hardness, 850 HK, is obtained with the maximum austenitizing temperature, 850 C, and is found to a depth of 1.10 mm. From that point to a depth of 2.70 mm, the hardness slowly decreases to that of the parent metal. With higher maximum surface temperatures, the lowest

Fig. 21

Influence of carbon content in steel according to start and finish temperature of martensitic transformation

Fig. 22

Hardness profiles for an induction-hardened 0.8% C steel for various maximum temperatures. The initial microstructure of the steel was pearlite. Source: Ref 40, 41

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hardness is obtained at the surface, and then, with a greater depth, it slowly decreases. With steel having a pearlite-ferrite or ferrite-pearlite microstructure, a microstructure consisting not only of inhomogeneous martensite but also of pearlite-ferrite grains will occur in the transition zone. Such a microstructure results in a stronger decrease in hardness than with the steel having a pearlite microstructure. It is essential for efficient induction surface hardening that constant hardness is obtained to a sufficient depth and that the hardness profile of the transition zone is adequate. It is only in this way that notch defects in the hardened layer may be prevented and better operation of the machine part under dynamic load may be ensured.

Magnetic Flux Concentrators In induction surface heating, unwanted areas on the workpiece are often heated or even hardened. This type of problem appears when the shape of the product surface is very complex, and it is therefore difficult to adjust the coil for local heating and quenching. For more demanding shapes, shields from materials with good heat conductivity, for example, copper, are often used on the product to prevent heating of the workpiece. Heating the copper shield prevents undesirable heating of certain areas as well as loss of heat. Such an example can be seen in Fig. 15, which shows the protection of two adjacent heat flanks while heating the tooth in the middle with an induction coil (Ref 15, 19–21, 28, 42). Heat losses occur due to local heating of the workpiece surface that is not to be hardened; however, this is done because it is not possible

Fig. 23

to adjust the shape of the induction coil. The heating temperature is higher than the hardening temperature, so that after quenching, unwanted increased surface hardening is obtained. In cases of treatment of two local surface areas on the workpiece that are in direct proximity, if heat is applied two times in sequence, the heating of the second area may result in tempering of the previously treated area. In these cases, the area hardened first may have a tempered microhardness with a lower hardness. Therefore, research has been done on how to form and adapt the coil to offer a more concentrated magnetic flux. By adjusting the concentration of the magnetic flux, it is possible to achieve localized heating of only those areas on the workpieces that are to be hardened. In the last decade, the development of induction coils has been directed toward achieving localized concentration of magnetic flux (Ref 19, 21). The purpose of magnetic flux concentration is to improve the efficiency of surface heating and reduce heat losses. The use of a magnetic flux concentrator enables selective local heating on workpiece/product areas with complex geometry. Figure 23 shows a straight conductor with current density distribution in points “A” and “B” for three cases (Ref 19, 42):

Current density distribution in a straight conductor (Fig. 23a) Current density distribution in a straight conductor when the conductive material (workpiece) is approached (Fig. 23b). The current density is greater in point “A”, closer to the workpiece material, than in point “B”. As a result, the workpiece is locally heated over a longer length—longer than the width of the conductor. This is referred to as current redistribution due to the proximity effect.

Current distribution in an inductor without/with a magnetic flux concentrator and its effect on the heating profile of the workpiece. Source: Ref 19, 42

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By placing the conductor into the magnetic flux concentrator made of powder metals based on iron, nickel, cobalt, and other powder materials, it is possible to achieve a higher density of the magnetic field (Fig. 23c).

The concentrated magnetic field results in localized heating on only those areas that are meant to be hardened. The magnetic flux concentrator surrounds the conductor, so that the current density on the conductor surface is redistributed, as is the heating of the workpiece material. The effects of the magnetic flux concentrator depend on:

Workpiece material and shape of the local area to which heat is applied Workpiece material and concentrator shape

The current density in this case is the highest in point “A” of the conductor and is considerably higher than in point “B”. This results in effective local heating, where the length of the heated area is only slightly greater than the width of the conductor, but, due to the high density of the magnetic flux, a considerably greater depth of the heated area is achieved. As for concentrator material, different applications require the use of different materials. The material for the concentrator must be chosen after considering several factors:

Relative magnetic permeability Magnetic reluctance Flux density in saturation Losses under magnetization Resistance to high temperatures Cooling abilities Resistance to cooling effects of fire extinguishers Good machinability Adjustment to different shapes and sizes of coils Ease of assembly and disassembly Manufacturing costs, depending on the kind of material, induction heating parameters, and geometry features between the concentrator and the workpiece

Figure 24 shows a local surface area with a complex shape that is induction heated with a single coil. Current density is in the inductor, which results from the position of the inductor with respect to the workpiece. This results in differences in the current density, which are

reflected in the different heating/workpiece depth profiles created by the differences in the power density distribution (Ref 19, 42). Figure 25 shows the same area and shape, but heating was performed with a magnetic flux concentrator (Ref 19, 42). This is placed on the left and right sides of the surrounding coil, which contributes to higher concentration of the magnetic flux to prevent heating of the sides of the workpiece at this place. A redistribution of power density took place, and a desirable local heating profile of the workpiece was achieved. Advantages offered by the magnetic flux concentrator are:

Smaller consumption of power Improved efficiency of heating Better exploitation of equipment due to shorter heating times More selective heating of the workpiece areas

Fig. 24

Heating profile on rotational workpiece with induction coil. Source: Ref 42

Fig. 25

Heating profile on rotational workpiece with flux concentrator. Source: Ref 42

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Achievement of desirable heating profiles of hardened layer profiles High repeatability of the procedure in terms of hardened profile size and hardened layer microstructure Efficient protection from unwanted heating of adjacent areas and successful prevention of reheating and softening the already hardened layers Elimination of detrimental effects on the operator’s health due to exposure to the magnetic field Improved operation lifetime of the heating equipment and higher productivity Less workpiece distortion and lower costs in straightening and/or final grinding subsequent to hardening Reduction of the quantity of rejected parts in terms of required size and shape of the hardened layer and workpiece distortion and cracking

Figure 26 shows two hardening procedures for tooth profiles of gear wheels using appropriate induction coils with a ferritic core concentrating the magnetic field (Ref 19). The first procedure is known as gap-by-gap hardening (Fig. 26a), because the induction coil with the ferritic core is moved in the tooth space so that a suitable gap between a tooth profile and the inductor is provided. The advantage of using magnetic flux concentrators rather than the conventional coil is that, with the same energy input, heating time is shortened. Unfortunately, the use of concentrators also shows some deficiencies, since the transition zone between the hardened and nonhardened microstructures, which is very important for hardening, is lost due to rapid heating. This results in less favorable residual-stress and microhardness profiles of the hardened layer. Thus, a greater risk of failure

Fig. 26

Gap-by-gap and tooth-by-tooth induction hardening of gear wheels. Source: Ref 19

is incurred at the point where teeth or other machine parts are most strongly loaded, which is at the transition from the hardened subsurface to the nonhardened core, where an increased stress concentration will occur. Such heating circumstances may be avoided by choosing lowerenergy inputs or power densities. In this case, however, the case depth is more difficult to control. If a hardened-and-tempered microstructure is to be provided at the tooth inside as well, heating with two frequencies is applied. The hardened-and-tempered microstructure may be accomplished with only one frequency in heat treatment of smaller gear wheels when the induction coil encompasses the entire gear wheel. The second procedure is tooth-by-tooth heating (Fig. 26b). In this case, the induction coil is moved against a tooth so that an appropriate gap is provided between the concentrator and the surface of both flank profiles of the same tooth. The difference between the two techniques is that in gap-by-gap heating, two adjacent tooth profiles and the root section of a gear wheel are heated, whereas in tooth-by-tooth heating the entire gear tooth is heated. Thus, in the first example, only hardened gear wheel tooth profiles are obtained, whereas in the second example, hardened tooth profiles of the gear wheel and quenched and tempered tooth inside are obtained. Figures 27(a to d) show various automobile parts that were induction surface hardened using the single-shot or scan-hardening technique (Ref 19). They were prepared for macroscopic and microscopic examinations of the hardened layer. The segments of the individual specimens of various characteristic and exacting automobile parts after grinding, polishing, and macroetching permit identification of the profile of the hardened pattern, for example, in the crosssectional and longitudinal direction of the part, respectively, as well as an analysis of the microstructure and microhardness. Each figure also indicates which technique of heating, that is, single-shot hardening or scan hardening, and which heating conditions (P, f ), using singleturn or multiple-turn induction coils, were employed. The quality of the surface-hardening process can often be efficiently assessed by measuring the hardened-layer depth at different locations on the individual parts. If the achieved depth is very uniform, regardless of the location of measurement, then the part was not subjected to distortion. A sufficiently high initial hardening

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because the part must be cut in the longitudinal and transverse directions respectively, so that the information required on the hardened pattern may be obtained. Such methods are appropriate only for periodical control, particularly statistical control with periodical sampling. Based on the results obtained for the characteristics of the selected hardened pattern, the quality of an article is confirmed or accepted and uninterrupted production provided. These types of nondestructive testing are long-standing; therefore, nondestructive methods of testing the material condition after induction hardening are gradually being introduced. Of all the methods, magnetic and magnetic-induction methods have established themselves because they are very fast, reliable, and provide reproducible results. The nondestructive methods of testing these parts allow on-line supervision of part quality, since, due to the speed of testing, all parts may be tested, which is a general tendency in mass production today.

Conditions in Induction Heating and Quenching of Machine Parts

Fig. 27

(a) Single-shot inductors used for both track (lobes) and shaft of this automotive component. The part is sectioned and acid etched to show the hardness pattern. All tracks are hardened at the same time using 250 kW/30 kHz; the stem is hardened using 135 kW/10 kHz. (b) Automotive right and left cam shafts that have been selectively induction surface hardened. The cam shaft was forged, heat treated, then final ground. No premachining was necessary. The equipment used a dual-spindle transfer mechanism; the coil was a four-turn coil that heated four lobes per spindle at a time. Power was applied for each set of four lobes: 150 kW, 10 kHz. The result is a 4.2 mm deep case depth in the base circle of the cam. (c) Hardness patterns on carbon steel crankshaft journals resulting from the stationary inductionhardening process. (d) (Left) An unacceptable nonuniform hardness pattern due to the nonuniformity of the workpiece, scanned with a single-turn inductor. (Right) An acceptable hardness profile achieved with a single-shot inductor. Source: Ref 19. Courtesy of Inductoheat, Inc.

temperature provides an adequate through-depth microhardness of the hardened profile. The analysis is considered a destructive method,

Heating of workpieces is done so that a magnetic field is created in the inductor, which is connected to a high-frequency generator. When a ferromagnetic material or workpiece is introduced into a magnetic field, eddy currents are induced. The distribution of eddy currents in the workpiece is specific, their density being highest on the surface and decreasing toward the inside. This phenomenon is known as the surface effect or skin effect. Due to resistance offered by the workpiece material, heating takes place mainly in the thin surface layer, whereas the inner core remains cold or is only slightly heated (low-mass workpieces). The depth of current penetration depends on workpiece permeability, resistivity, and the alternating-current frequency. Because the first two factors vary comparatively little, the greatest variable is frequency. Depth of current penetration decreases as frequency increases. High-frequency current generally is used when shallow heating is desired; intermediate and low frequencies are used in applications requiring deeper heating. Most induction surface-hardening applications require comparatively high power densities and short heating cycles to restrict heating

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to the surface area. The principal metallurgical advantages that may be obtained by surface hardening with induction are the same as for flame hardening. The drop in magnetic permeability of steels depends on the temperature line, TA2 , where steel transforms from magnetic into nonmagnetic ferrite. The larger the effect of the magnetic permeability change on the temperature line, TA2 , the smaller the carbon content in the steel (the larger the proportion of ferrite in the steel) and vice versa. Due to rapid heating, phase transformation moves upward toward higher temperatures. The temperature-time curves of heating along the depth of a cylindrical component depend on the kinetics of the magnetic transformation, TA2 , and the effects of other phase transformations during induction heating. A thickness of 1.0 to 1.5 mm is reached with a medium-frequency current. The temperaturetime variation over the cross section of the steel workpiece is a function of the following factors:

Penetration depth of eddy currents Heat conduction of the material Heating rate of the surface Initial temperature of the surface Size and shape of the workpiece

The depth of penetration of the heat is governed mainly by the power and frequency employed. The normal power density is 0.1 to 2 kW/ cm2 of the heated surface. The relationship between depth of penetration and frequency can be calculated approximately by using simplified expressions, which are valid for the temperature rise in steel up to the hardening temperature (Ref 16): 20 dCS = pffiffiffi f

cold state (20 C)

500 dHS = pffiffiffi f

hot state (800 C)

where dCS is the depth of penetration in the cold state, measured in millimeters; dHS is the depth of penetration in the hot state, measured in millimeters; and f is the frequency, measured in hertz. Due to heat conduction in the material during heating, the overall depth of penetration is larger. It is possible to calculate the additional penetration due to heat conduction from the expression: pffi dHC =0:2 t

where t is time, measured in seconds; and dHC is the depth of penetration for heat conduction, measured in millimeters. The total depth of penetration is obviously dT = dCS +dHC. It should be stressed that these expressions give only a rough estimate of the depth of penetration, and they have been included here only to show the fundamental effects of frequency and time. In flame heating, the temperature achieved on the surface at equal energy input is considerably higher than in induction heating, the overheating and the hardened layer thickness being dependent on the heat conduction of the workpiece material. Figure 28(a) shows the temperaturetime variation over the cross section of the workpiece in flame heating (Ref 15). Characteristic of this variation is that the temperature rapidly changes with time, and therefore, the conditions for the formation of a homogeneous austenitic microstructure are not fulfilled. Figure 28(b) shows the temperature–time variation over the cross section of the workpiece in induction heating. The temperature variation is very similar to that in flame heating up to magnetic transformation, that is, to line A2. At temperatures higher than line A2, eddy currents grow characteristically, and the rate of heating decreases sharply. This slows down the heating above temperature line A2. A reduced rate of heating on the surface provides the conditions for faster heating into the depth of the workpiece. This figure shows that a relatively thin layer is heated up, but the layer has a rather homogeneous austenitic microstructure. The temperature-time variation on the workpiece cross section, or the temperature field, depends on the workpiece size and shape. Thus, in heavymass workpieces, faster heat abduction into the remaining cold part of the workpiece is achieved, and that is why the actual variation of temperature over the cross section is steeper. This means that in heavy-mass workpieces, a higher surface temperature than in low-mass workpieces must be ensured to grant the same penetration depth. The microstructural changes in induction hardening depend to a large extent on the rate of heating and subsequent cooling. The rates of heating range from one to a few seconds, which means that the diffusion processes may become jeopardized. In steel, transformation of pearlite into austenite takes place in induction heating at almost the same temperature as in conventional heating. In subeutectic steels suitable for surface hardening, it is important that the induction

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Fig. 28

Temperature profile across the workpiece diameter in (a) flame surface heating and (b) induction surface heating. Source: Ref 15

heating procedure ensures enough time for the diffusion of carbon for transformation of ferrite into homogeneous austenite (Ref 2, 16, 27, 29, 30). Figure 29 shows the temperature line of through hardening of subeutectic steels with different carbon contents versus different heating rates (Ref 15, 27). The graph shows that austenitization is clearly influenced by the heating rate, especially when carbon concentrations are low to medium. Likewise, in the case of rapid heating for through hardening, a considerably higher temperature is needed than for normal hardening. Thus, in surface hardening carbon steels, it is difficult to ensure enough homogeneous martensite, whereas in the transition temperature range (TA1 to TA3 ), a higher homogenization of martensite is achieved but with a presence of ferrite. The proportion of ferrite in the transition temperature range is thus higher with faster surface heating and smaller carbon content in the steel. In surface hardening alloyed steels, a better homogenization of the austenite is achieved in the heating phase, and when quenching is completed, a very homogeneous martensite with a uniform microhardness along the depth of the hardened layer is derived. Unfortunately, there is a transition temperature range with a smaller ferrite content, which causes a sharp drop in

Fig. 29

Influence of induction surface heating rate on hardening temperature for subeutectic steels. Source:

Ref 15, 27

hardness in the transition area with a hardened and nonhardened microstructure. A very important heating requirement in the hardening procedure, besides the austenitization temperature, is the time necessary for austenitization, since both of these heating parameters

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affect the size of austenitic grains. Austenitic grain size is, on the other hand, dependent on the martensite formed subsequent to quenching and is reflected in the toughness of the surface layer. That is why heat treatment conditions are sought that would ensure the finest and most homogeneous austenitic microstructure in the heating and overheating phases, leading to the formation of very fine martensite with the highest possible toughness of the hardened layer after quenching (Ref 27, 29). This microstructural condition and the achieved mechanical properties ensure good wear resistance of machine components (gears) and their good response under dynamic loads. Figure 30 shows the growth of austenitic grains for different rates of induction heating for a steel with 0.4% C that can be used for induction hardening (Ref 15, 27). With increasing rates of heating, the austenitic transformation moves toward higher temperatures. A higher heating temperature creates a higher formation rate of austenitic crystal and therefore fine grains of austenite. Subsequent to quenching, these fine grains of austenite ensure a very fine martensitic microstructure. Figure 31 shows the relationship between the hardened layer hardness and the heating rate and temperature for a steel with 0.45% C (Ref 15, 27). For each steel, there exists a certain temperature range after hardening that yields the best microstructural condition and thus the best properties. With higher rates of heating, this

Fig. 30

Influence of surface heating rate on austenitic grain size. Source: Ref 15, 27

range moves to higher temperatures. This means that quenching from lower temperatures leads to imperfect hardening, while higher temperatures yield medium or rough needles of martensite. The heating rate of 50 C/s is sufficient to reach a hardening temperature range between 850 and 925 C. In the case of a higher-energy input that heats up the thin surface layer with a rate higher than 140 C/s, the required temperature range becomes 870 to 970 C. In both cases, a surface hardness of 60 HRC is reached subsequent to quenching. Figure 32(a–e) shows the entire process of induction hardening a cylindrical component with a small diameter or cross section, which, subsequent to hardening, leads to self-tempering (Ref 15, 27). Figure 32(a and b) show induction heating of a thin surface layer to the austenitic temperature range, ensuring, a sufficient thickness of the austenite layer d2 subsequent to quenching. The process of quenching or self-tempering is shown in Fig. 32(c and d), where, due to heat conduction into the remaining cold part of the workpiece, the temperature on the workpiece surface layer increases, and thus, the thickness of the austenite layer increased to d3 or d4, respectively. Due to a small workpiece cross section or low workpiece mass causing heat conduction, the temperature in the middle of the workpiece rose to a point corresponding to the tempering temperature of the given material. Since the process of tempering takes place in the workpiece with the available heat needed for austenitization, self-tempering of the workpiece is indicated (Ref 27, 31, 32, 43). Specific properties of the hardened layer can be described by analyzing the microstructure

Fig. 31

Hardness reached after induction surface hardening at various heating rates in steel with 0.45% C. Source: Ref 15, 27

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Fig. 32

Individual phases in induction heating and spray quenching in the workpiece surface layer and corresponding temperaturediameter diagrams. Source: Ref 15, 27

with hardness measurements on the surface or microhardness measurements in the cross section of the hardened layer, and by measuring residual stresses (Ref 27, 44, 45). Figure 33 shows the hardness profile in the hardened layer subsequent to induction surface hardening (line 1) as a function of different carbon content in the steel after conventional hardening (line 2) (Ref 15, 27). Induction surface-hardened layers normally have on the average 3 HRC higher hardness on the surface than that achieved in the same kind of steel after conventional hardening. This is primarily due to a finer martensite and compressive residual stresses present in the induction surfacehardened layer (Ref 2, 27, 46).

Time-Temperature Dependence in Induction Heating The time variation of temperature in induction heating of a thin surface layer depends on the

Fig. 33

Influence of carbon content on steel hardness after various heat treatments. Source: Ref 15, 27

type and shape of the induction coil used, the alternating-current frequency, f (Hz), the and power density Q (W/cm2). Power density is defined by the selected power of the highfrequency generator and the surface layer of the workpiece. Surface heating depends on the coupling between the induction coil and the workpiece. Figure 34 shows the influence of

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the selected power density and frequency on the reference depth of the skin effect in a ferromagnetic material (Ref 2, 20). A higher power density results in a greater reference depth of the skin effect and a greater depth of the heated layer with the same maximum temperature obtained at the workpiece surface. Figure 35 shows the interdependence between the heating parameters, that is, power density and generator frequency, as a function of the specified depth of the hardened layer and the heating time required for single-shot techniques of surface induction hardening (Ref 2, 20). The data supplied by the diagram can have a character of information only, but they make the selection of an optimal surface induction condition easier. With the scanning technique of surface induction hardening, however, the speed of the workpiece movement, rather than time, assures the depth of the required hardened layer and should be defined. Generally, longer heating times are required with smaller power densities and vice versa. For the same depth of hardened layer, longer heating times with lower current frequencies are also required. With regard to the depth of hardened layer selected between 0.5 and 10.0 mm, generator frequencies of 450, 10,

and 3 kHz can be selected in the single-shot surface-hardening technique, in which case appropriate power densities between 2 and 50 MW/m2 are obtained. With lower high frequencies, such as 10 kHz, the same depth of surface-hardened layer, that is, 2.0 mm, can be ensured only when the power density is changed to 50 MW/m2. The lowest generator frequency, 3 MW/m2, shown in Fig. 35 cannot ensure the depth of a hardened layer smaller than 2.5 mm. Immediately after tempering, an intensive inverse heat flow is expected as well. The power density and frequency used in induction hardening are usually based on the shape and size of the machine part to be surface hardened; the case depth is specified also. In addition to the depth of hardening, the case pattern along the entire length of the machine part is important. Regardless of the complexity of a workpiece shape, case depth and transition that are as uniform as possible as well as regular-shaped hardened ends should be provided. Inadequate pattern transitions may produce high stress concentrations related to the given loads in its vicinity, so the material cannot resist dynamically loaded parts.

Fig. 34

Fig. 35

Reference depth of skin effect as a function of power density and selected generator frequency for ferromagnetic steel. Source: Ref 2, 20

Influence of high-frequency generator on selection of power density and heating time with given thickness of surface induction-hardened layer. Source: Ref 2, 20

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Various heating conditions that are defined by the power density and frequency provide different pattern depths. The third influencing parameter is heating time. Thus, different hardness and residual-stress profiles may be achieved. The steel grade to be hardened, the loop shape, and the gap size between the inductor and the workpiece should be considered. The data found in the diagram apply only to stationary hardening; therefore, with scanning surface hardening, these values should be suitably corrected. For progressive hardening, the hardening conditions should be slightly corrected to allow for the loop movement. With the surface heat treatment processes, studies are often conducted on the influence of the selected heating and quenching conditions on the depth of the hardened layer and the size of the transition zone between the hardened and nonhardened microstructures. One simple and practical procedure to control surface heat treatment is to measure the time variation of temperature from the beginning to the end of the heating process and also from the beginning to the end of the quenching process. The heat process may be changed by changing the power density and the generator frequency, whereas the quenching process may be changed by the selection of different quenching agents and quenching processes. Figure 36 shows the measured and calculated temperature cycles for the surface, the core, and in a radius, r, of 7 mm in a depth of 1 mm in a cylindrical specimen (Ref 47). A comparison of the temperature cycles shows that in surface induction hardening, a thermal flow of 3 MW/mm2 in heating and 5.8 MW/mm2 in quenching was selected. Under

Fig. 36

Comparison of variations of calculated and measured temperature cycles for a cylindrical specimen 16 mm in diameter. Source: Ref 47

such heating conditions, a maximum temperature of nearly 1000 C was attained, while heating above a temperature of 800 C was somewhat slowed down. The data in the diagram show that the time required for heating the specimen from the ambient temperature to 800 C is equal to the time required for heating from the latter to the maximum temperature obtained at the surface, that is, 1.6 s. A temperature cycle at the surface takes 3.2 to 3.3 s. The temperature differences between the surface and the core in a given moment are the greatest during the heating process, that is, DTmax ffi 600 C. During the quenching process, however, they can reach up to 360 C. Temperature gradient changes are much stronger in heating than in quenching. In material heating, there is also a great difference in yield stress of the material, which can produce plastic deformation. Another very important finding (Ref 47) is that the theoretical model is appropriate, since the results obtained were confirmed by the standard experimental methods such as temperature measurement with thermocouples, diamond pyramid hardness test, and measurement of residual stresses with x-ray diffraction. The difference between the measured and the calculated temperature cycles is very small. It occurs mainly in heating and reaches up to 60 C at maximum, not taking into account the losses due to eddy currents. Figure 37 shows the variation of temperature from the surface toward the core with various volume power densities, Q, that is, 0.4 · 109, 1.2 · 109, and 2.4 · 109 W/m3 (Ref 48).

Fig. 37

Calculated variation of temperature through specimen cross section in induction heating up to hardening temperature with different volume power densities. Source: Ref 48

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With smaller volume power densities, the maximum temperature differences between the surface and the core also become smaller and the austenitizing times longer. Because of higher temperatures attained in the core, the inverse heat flow may even be so high as to produce a change in microstructure and, consequently, a reduction of hardness in the surface layer. The induction heating industry has standardized power supply frequencies, and probably 99% of the power used will be at the frequencies listed in Table 2. Also included in Table 2 is the type of equipment used to a change 50 Hz (60 in some countries) to a higher frequency and the conversion efficiency (Ref 23). Melander (Ref 49, 50) first treated single-shot surface induction hardening of low-alloy steel with 0.4% C, 0.7% Mn, and 1.1% Cr for tempering and hardening as well as surface hardening. For an analysis, a representative-sized machine part that is most often used, that is, a cylindrical specimen 40 mm in diameter, was chosen. Induction-heating conditions were Table 2 Power sources, frequencies, efficiency, and power for induction heating equipment Type

Vacuum tube oscillators Motor generators Frequency multipliers Frequency inverters Source: Ref 23

Fig. 38

Power (P), kW

Efficiency (g), %

Frequency ( f ), kHz

5–600

50–60

200–450

7.5–500 100–1000 50–1500

75–80 90–95 85–95

1, 3, 10 180 and 540 0.5, 1, 3, 10

selected so that the temperature of the diameter, TA1 , was exceeded to a depth of 5.0 mm. This means that a change of microstructure and hardness was expected even to the depth of 5.0 mm, where only partial austenitization was obtained. Induction-heating conditions were chosen that subjected the surface layer to heating for up to 35 s. Time variations of temperature at the surface of the cylindrical specimen and in its subsurface at depths of 2.0, 4.0, and 10.0 mm are shown in Fig. 38 (Ref 49, 50). Time-temperature diagrams differ from the previous ones, since a distinctive deviation occurs in heating the specimen material when the temperature of magnetic domain, TA2 , is reached and exceeded. The course of surface heating indicates that the transformation temperature TA2 was obtained in 10 s. In spite of the same power density, further heating of the surface up to a temperature of 850 C was very slow due to the nonmagnetic character of the steel surface layer; it took another 28 s. It is difficult to assess which models of induction heating are more suitable than others. However, with the heating process suggested, it is possible, after quenching, to ensure a homogeneous, fine austenitic microstructure exhibiting the finest martensite with the highest possible hardness of the given steel. Due to the presence of alloying elements and high cooling rates of the surface layer, in addition to fine martensite, up to 3% residual austenite also appears. Based on the time-temperature variation of heating, the depth of the hardened layer ranges between 2.0 and 4.0 mm. From the

Time-temperature cycle during single-shot surface induction heating and quenching. Source: Ref 49, 50

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time-temperature variation at a depth of 10 mm, it can be assumed that the specimen was heated through the entire volume, that is, to the very core of the specimen. Because of a different time variation of temperature in the fourth temperature cycle, it may be concluded that the maximum temperature obtained in the depth of 10 mm is lower than the magnetic transformation temperature of the given steel. Because of strong overheating of the cylindrical specimen toward its center, lower temperature gradients occur, which result in a reduction of thermal stresses during the heating process. According to a comparatively high temperature in the core, temperature gradients between the surface and the core are generally lower, which produce a decrease in axial internal stresses generated during quenching and also a decrease in axial residual stresses. Figure 39 shows the time variations of through-thickness temperature in single-shot induction heating with high-frequency generator powers of 40, 60, 100, and 180 kW and a medium-frequency current (Ref 51). Characteristic temperature transformations, TA1 and

Fig. 39

TA3 , for equilibrium heating are plotted (Ref 51). Induction heating is a very fast process; therefore, the temperature transformations shift to higher temperatures. To obtain homogeneous austenite in the surface layer, it is necessary to heat the surface layer to the hatch-marked temperature range. The results shown in the four diagrams make it possible to draw the following conclusions:

The heating curves differ strongly. Only with the power of 40 kW does the temperature not reach the hatch-marked temperature range in 30 s. With all other powers higher than 60 kW, the hatch-marked temperature range is reached in a shorter heating time. With all powers higher than 60 kW, more intensive heating of the surface layer occurs, which is shown in a steeper hardness curve in the hatch-marked temperature range.

The first question raised by an engineer would be how to evaluate the through-thickness variations of hardness and residual stresses with reference to the time variation of temperature.

Time-temperature variations in single-shot heating at various powers. Source: Ref 51

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The second question would be how to select induction-heating parameters to obtain optimal properties of the surface-hardened layer with a minimal energy input. In answer, the initial steel microstructure must be taken into account and an appropriate quenchant selected to ensure a cooling rate in the surface layer equal to or a bit higher than the critical cooling rate required. Figure 40 shows the time and through-thickness temperature variations of the specimen on heating with different powers and different velocities of workpiece movement, that is, v1 = 140 mm/min (Fig. 40a), v2 = 220 mm/ min (Fig. 40b), v3 = 370 mm/min (Fig. 40c), and v4 = 680 mm/min (Fig. 40d), in scan hardening (Ref 51). In the first case, with v1, the hardened layer is obtained with the powers of 74 and 56 kW. In the second case, with a higher velocity of movement, v2, the hatch-marked temperature range is reached only in heating with the powers of 78 and 59 kW. In the third case, with v3, the hatch-marked temperature range is reached only in heating with the power of 112 kW, whereas in the fourth case, with the highest velocity, v4, this temperature range is

Fig. 40

not reached in spite of a very high power, 151 kW.

Quenching Systems for Induction Hardening Control of the quenching process of a gear wheel from high hardening temperatures ensures martensite microstructure. The cooling rate should be high enough to prevent formation of unwanted softer microstructures, such as pearlitic or bainitic ones; therefore, it is very important that in the development of new systems of induction hardening, quenching systems are designed properly. Process parameters should be accurately controlled to ensure permanent and reproducible results after hardening of machine parts. It is also indispensable to define individual parameters and determine their permissible deviation in operation, presuming that the deviations have a negligible influence on the results of heat treatment. The quenching system for induction hardening is defined by eight parameters (Ref 17,

Time-temperature variations in scan hardening at various scan speeds. Source: Ref 51

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52): heat time/scan rate, power level, power frequency, part position/rotation, quenchant concentration, quenchant flow, quenchant temperature, and quenching time. Important simultaneous changes of one or more of these parameters can produce unwanted effects on the workpiece, which, in extreme cases, result in an unsuitable microstructure, deviations in the depth of the hardened profile, unsuitable hardness variation (too low hardness, soft spots), and exceeding the distortion of the machine element. In practice, there are machine parts of different shapes and sizes, requiring different depth of the hardened layer. In these cases, the type of material chosen and its through hardenability should be considered. Thus, with alloyed steels having good through hardenability and/or machine parts with a comparatively thin hardened surface layer, the martensite microstructure can be obtained without the application of a quenchant. In such cases, heat sinks from the surface into the cold workpiece core, so that the critical self-cooling rate obtained at the surface is higher than the critical cooling rate (Ref 24, 53). Figure 41(a) shows the time variation of the temperature measured at the individual measuring points of the gear-wheel tooth during quenching with a pressurized water jet (Ref 52). Data on the temperature variation show that the cooling temperatures were very similar at measuring points 1 and 2, somewhat lower at measuring point 3, and the lowest at measuring point 4, where, after 40 s of quenching, it still equaled approximately 350 C. Figure 41(b) shows the calculated time variation of temperature at the same measuring points of the gear-wheel tooth during quenching, taking into account selected heat-transfer coefficients, a0, of 5000 and 10,000 kcal/m2 Ch (Ref 52). The calculated time variations of temperature differ from the measured ones only at those measuring points with higher cooling rates. The time variation of temperature refers only to the selected measuring points; therefore, assessment of the stress state during the quenching process and the distortion of gear-wheel teeth or a tooth is very difficult. The authors focused on an analysis of the conditions during the quenching process by means of a calculated distribution of temperature, shown as isotherms (a0 = 5000 kcal/m2 Ch) at the half-cut of the gear-wheel tooth. Figure 42 shows the variation of isotherms during the

quenching process after 0.1, 1.0, and finally 10 s of cooling (Ref 52). In quenching, it is very important that the temperature gradients are high enough to prevent plastification, that is, distortion, of the gear wheel. Such quenching conditions, including high-temperature gradients, result in a decrease

Fig. 41

Temperature profiles of tooth gear at given measuring points during cooling process. (a) Measured. (b) Calculated. Source: Ref 52

Fig. 42

Temperature distributions process. Source: Ref 52

during

the

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in favorable compressive residual stresses after hardening. It follows that a sufficiently high cooling rate must be ensured at the tooth surface as well as at the required depth of hardening. That means that temperature gradients in the gearwheel tooth must be attained so that gear-wheel teeth are hardened with the minimum required cooling rate, ensuring the required martensite transformation. A sequence of images showing isotherms in very short time intervals from the beginning of quenching allows assessment of the size of distortion of the tooth and residual stresses to be expected after quenching. Figure 43(a) shows the measured volumefraction distribution of martensite and Fig. 43(b) the calculated volume-fraction distribution of martensite when different values of the heattransfer coefficient, a, that is, 5000, 7500, and 10,000 kcal/m2 Ch, are taken into account (Ref 52). The main difficulty encountered in the calculation is how to determine the heat-transfer coefficient to obtain a description of real quenching conditions. Quenching is very intensive, since it is carried out under a pressurized waterjet; the gear wheel has been induction heated, whereas the core is cold. Two common quenching methods that use a quenchant are spray quenching and immersion quenching. Particularly popular spray quenching techniques that offer different possibilities are:

Spray with progressive scan heating (scan hardening) Spray after heating in position (single-shot hardening)

Fig. 43

Spray quenching out of location after heating

Spray quenching is carried out immediately after heating with a short pause. It is used with machine parts made of steels with good hardenability and with a comparatively thin hardened layer. Immersion quenching is carried out after the inductor has been disconnected from the high-frequency generator. Gradual hardening is very efficient and prevents the influence of inverse heat from the core toward the surface, so that no reverse heat flow and no heating of the already cooled surface occur. This problem is characteristic of large parts, where the workpiece surfaces should be hardened selectively. In the past, straight oils and water-soluble oils were used for quenching after surface induction heating. Straight oils and water-soluble oils produce mild quenching effects, which reduce difficulties due to distortion and/or crack initiation. Straight oils require immersion quenching to minimize the risk of oil ignition, whereas water-soluble oils are more suitable for spray quenching. Other quenchants used are water, polymeric water solutions of different concentration, water-soluble oils, saltwater, and so on. Polymeric water solutions are inflammable quenchants. They are prepared in various concentrations to be suitably adapted to various cooling rates. For spray quenching, various flow rates of quenchants can be selected. They depend on the size of inlet openings and the

Martensite distribution of the hardened tooth gear. Source: Ref 52

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number of outlet nozzle boreholes in one, two, or more rows at the quenching ring. The fluid flow rate can also be regulated by pressure. Technical literature provides data on shaping of quenching nozzles as well as on the selection of the quenchant flow rate required. It is very important that a sufficient flow rate of the quenchant is chosen to ensure sufficient heat removal from the workpiece surface layer during quenching. The flow rate is also controlled by the number of boreholes for the flow and spraying of the quenchant, respectively. An appropriate arrangement of the boreholes in the ring permits a uniform heat removal after spray and proper and uniform cooling. The borehole cross sections should be only 15% of the total available ring surface and of the inductor in the singleshot-type application, respectively. Rotation velocities of the workpieces in surface induction heating are relatively high and can be selected between 800 and 1000 rotations per minute, whereas the rotation velocity of the workpieces during quenching is considerably lower, that is, 40 to 60 rotations per minute. Very high rotation velocities of the workpieces are required when a uniform and reproducible depth of hardened layer is to be ensured, and particularly with relatively short heating times. Figure 44 shows the dependence of cooling rate changes on the momentary temperature at the workpiece surface for four different concentrations of polymeric water solutions (Ref 24). The selection of an appropriate concentration of the polymeric water solution for the selected steel and the required depth of hardened layer should ensure minimum distortion of the workpiece, which is the final purpose of any heat treatment. Thus, the required heat removal from the heated surface layer of the workpiece, considering the thickness of the heated layer, should be ensured by an appropriate quenching system.

control. An appropriate machining technology and appropriate heat treatment processes should be selected, and internal stresses lower than the yield stress should be ensured at any moment and any location of a machine part during heating and/or cooling. Analytical methods provide an insight into heat treatment conditions if the time variation of internal stresses is monitored and the dependence between the cooling time and the specimen temperature at each point is known. With regard to the specimen temperature determined in this way, the specimen yield stress at each point can be determined as well. Consequently, with different heat treatment conditions, different values of physical quantities can be selected. They are reflected in the changed conditions in the material, which make it possible to study distortion of the machine part during cooling and to determine the magnitude of the residual stresses. Figure 45 shows the time variation of axial stresses at the surface and in individual depths, that is, from 2.0 to 4.0 mm under the surface and in the middle of a cylindrical specimen (Ref 50). In surface heating, tensile stresses occurred at the surface due to thermal extension of the

Time Variation of Stresses and Residual Stresses With a certain heat treatment performed, the required microstructural changes and an appropriate magnitude and variation of hardness are obtained. It is also required that distortion of a machine part be as small as possible so that the final size of a machine part can be obtained with minimum precision machining. Consequently, it is very important that distortion be kept under

Fig. 44

Polymer additive ratio effects on the cooling rate. Source: Ref 24

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Fig. 45

Axial stress distribution at various depths below the surface during single-shot induction surface hardening. Source: Ref 50

surface layer, but when the transformation temperature of the pearlitic- to austenitic microstructure was exceeded, additional compressive stresses occurred at the surface. The compressive axial stresses transform into tensile stresses in the zone of undercooled austenite. In the transformation of undercooled austenite into martensite, the compressive stresses increase with the increase in martensite fraction. In the core, the opposite sign of the stress was obtained, that is, the tensile stress. At the end of quenching, the compressive residual stresses obtained in the surface-hardened layer were approximately 1600 N/mm2, and the tensile stresses in the core were approximately +870 N/mm2. Internal stresses are induced in heat treatment by temperature and microstructural changes. Residual stresses in the induction surfacehardened layer are always of a compressive nature, are relatively high, and have a good effect on dynamically loaded components. The existence of residual stresses in the radial direction, that is, into the depth of the hardened layer, is very important, as is the absolute value of residual stress on the surface and the stress profile in the transition from compressive into tensile stresses (Fig. 46) (Ref 15). In the case of induction hardening, a maximum compressive residual stress in the surface layer is achieved, which is very desirable for dynamically loaded components. The transition from compressive into tensile stresses should be

Fig. 46

Residual stress profile below the surface after induction surface hardening. Source: Ref 15

as gentle as possible, lessening the effect of stress concentration in loaded components. This contributes to the fact that a machine component is less susceptible to overloads in operation. It has been shown that residual stresses are closely linked to hardness variation and microstructure in the transition zone of the hardened layer, that is, in the narrow range between the hardened and nonhardened microstructure. Some examples of the dependence between microhardness and residual stresses are shown in

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Fig. 47 (Ref 15). Figure 47(a) shows a very steep microhardness profile in the transition zone and the highest compressive stress in the surface, and, related to this, a very steep transition of residual stresses into the tensile range. The change from compressive into residual stresses happens at the transition between the hardened and nonhardened areas. Figures 47(b and c) show graphs of broader transition zones, which produce a change in the microhardness and also the residual stresses (Ref 15). Thus, Fig. 47(c) shows a very slight drop in hardness in the transition zone for the same microhardness on the surface. This is reflected in lower compressive stresses in the surface, accompanied by a slight change of stresses in the transition zone into lower tensile residual stresses. In induction surface hardening, the engineer should choose the kind of heat treatment conditions that will result in the microhardness and residual-stress profiles shown in the last two examples in Fig. 47. Investigations of residual stresses after induction surface hardening have confirmed that when the hardened layer is 2 mm thick, the change of compressive into tensile stresses happens in compliance with the transition zone, that is, the achieved depth of the hardened layer. When the thickness of the hardened layer is greater than 2 mm, the transition from compressive into tensile stresses happens in the hardened zone, that is, the martensite microstructure. This means that induction surface hardening is much

Fig. 47

more difficult if the thickness of the layer is above 2 mm. Due to intensive cooling, the internal stresses in the surface layer may become so high that they cause failure of the component. This failure is very typical, because the surface layer separates from the core due to high radial stresses. Important features on the curve are:

Maximum value of compression residual stress in the hardened surface layer Maximum value of tension residual stress in the hardened surface layer Transition width of compression to tension of residual stress in the hardened surface layer Transition steepness of compression to tension of residual stress profile Layer depth with transition microstructure

The examples shown in Fig. 46 and 47 were based on a presumption that the maximum hardness and the maximum value of the residual stresses obtained were independent from the heating power density and the heating time (Ref 15). In the case of induction surface hardening, a maximum compressive residual stress in the surface layer is achieved. The transition from compressive into tensile residual stresses should be as gentle as possible, lessening the effect of stress concentration in loaded components. This contributes to the fact that a machined component is less susceptible to overloads in operation. It has been shown that residual

Various residual-stress and hardness profiles below the surface. Source: Ref 15

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stresses are closely linked to hardness variation and microstructure in the transition zone of the hardened layer and base material, that is, in the narrow range between the hardened and nonhardened microstructure. On the basis of the heating temperature cycle, microstructural changes in the surface layer after quenching can be predicted. Consequently, the temperature cycles at the surface, that is, a through-thickness heated layer, have a decisive influence on the variation of residual stresses after quenching. The variation of the temperature cycles in heating can be adjusted by the selection of adequate energy input with reference to the steel microstructure and its thermal properties. Figure 48 shows three examples of induction surface hardening, each with different energy input during heating, whereas the quenching process was the same in all cases (Ref 15, 54). The variation of residual stresses after induction surface hardening is affected by the energy input through the workpiece surface and thermal conductivity of the workpiece material. The energy input in the workpiece depends on the power chosen, the duration of high-frequency current, the shape of an induction loop, the size of a gap between the induction loop and the workpiece, and the area exposed to the energy input. In progressive hardening, the velocity of progression of the high-frequency loop along the heated area must be considered. The transition from tensile to compressive residual stresses presents a serious danger for catastrophic failure. From the aforementioned, it can be understood that the workpiece must be slowly heated and also quenched with the correct cooling rate. The actual cooling rate is very important throughout the martensitic transformation and must be as close as possible to the critical cooling rate. The quenching process must be carried out very carefully in the martensitic transformation temperature range to ensure internal stresses lower than the yield point. The only barely-seen difference is in the transition area, when the compressive residual stress starts to fall rapidly. In the transition area, the residual-stress curve is steeper and less desirable as far as notch effects under dynamic loads are concerned. In general, efforts should be made to achieve low or moderate residual stresses in the transition area after induction surface hardening. Manufacturing engineers are often confronted with the question of how to

obtain this kind of residual-stress profile. It should be emphasized that because of their lower heat conductivity, alloyed steels for surface hardening are very difficult to treat, and it is difficult to ensure desirable residual stresses in the transition area. This is why special alloyed steels are available that display a more favorable behavior in heating and quenching. Because of these difficulties, delamination of the surface to the very core may occur. Figure 48(a) shows that that the energy input was too weak, since only a partial transformation of the steel matrix into austenite occurred (Ref 15, 54). After quenching, this will produce only a small portion of martensitic microstructure, which is followed by a measured variation of residual stresses. The second example, shown in Fig. 48(b), indicates that heating with a high-power density was very intensive. The result of a too-high power density is a distinct influence of heat conduction from

Fig. 48

Residual-stress profiles after induction surface hardening at various input energies. Source: Ref 15, 54

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the surface to the subsurface, which, in turn, resulted in a higher temperature measured in the subsurface than at the surface after the delayed interruption of heating at the beginning of the quenching process. Because of such a heating regime and energy input at the lower limit of the required energy, considerably lower compressive residual stresses were achieved at the surface. In this case, the highest value of compressive residual stress was obtained at a depth of 1.0 mm. The variation of the stress was not the most adequate, since induction surface hardening can provide considerably higher compressive residual stresses in the surface layer. The anticipated variation of the residual stresses, that is, very high compressive residual stresses ranging between 800 and 1000 N/mm2, is shown in Fig. 48(c) (Ref 15, 54). A very favorable variation of residual stresses can be found next to the transition zone between the quenched and unquenched surface layers, which ensures favorable behavior of a machine part under dynamic loads. In practice, the progress of heating and quenching can be monitored by physical modelling, in which case the temperature cycle at the workpiece surface as well as the temperature cycles in the individual depths are obtained. The variations of the temperature cycles in heating and the known temperatures, that is, temperature ranges, make it possible to determine changes of a pearlite-ferrite microstructure into an austenitic one. The latter makes it possible to predict the microstructure after surface hardening.

Fig. 49

The progress of induction surface heating can be supervised by measuring the temperature at the workpiece surface using a pyrometer. The temperature measurement is a determination of temperature cycles in individual depths during induction heating and subsequent cooling. There are three basic and important metallurgical conditions that will affect a successful induction surface-hardening operation: Lower critical temperature, where metallurgical phase transformation begins Upper critical temperature, where the formation of austenitic grain microstructure is complete Surface temperature, which affects the resultant grain size of the heat treated microstructure For a typical AISI 140 material, the steel reference books will list TA1 =750 C, TA3 =795 C, and Tsmax =900 C. However, because of the shorter heating times with induction heating, the actual values required for a satisfactory microstructure and process may be 30 to 60 C higher than for conventional furnace heat treatment processes. Figure 49 shows the temperature cycles in the surface and in the individual depths during induction heating and quenching (Ref 15, 54). The temperature range of the pearlite-ferrite transformation into austenite, which is a condition for the formation of the martensitic microstructure after quenching, is plotted also. On the basis of the heating temperature cycle, microstructural changes in the surface layer after

Temperature cycles during single-shot induction surface hardening, at various depths. Source: Ref 15, 54

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quenching can be predicted. Consequently, the temperature cycles at the surface, that is, a through-thickness heated layer, have a decisive influence on the variation of residual stresses after quenching. The variation of the temperature cycles in heating can be adjusted by the selection of adequate power density with reference to the steel microstructure and its thermal properties. Figure 50 shows graphic representations of temperature cycles at the surface at various power densities in single-shot induction hardening (Ref 15, 54). Each temperature cycle includes induction heating, which is followed by quenching a short while after the interruption of heating. Four temperature cycles of induction heating and quenching measured at different heating power densities using high-frequency current were plotted. With the same inductor having the same gap width between the inductor and the workpiece surface, but with different power densities, different temperature cycles of heating were obtained. With a higher power density, the heating curve was steeper and the heating period shorter (Ref 54). In all cases of heating, however, the same depth of the hardened layer was ensured. With the given power density, however, it is more difficult to ensure the same depth of the hardened layer, since the heating times are considerably reduced. With very short heating

Fig. 50

times, it is much more difficult to ensure the same depth of the hardened layer, which indicates a worse repeatability of the results. The shorter heating times with the same cooling times contribute to a shorter hardening cycle and result in higher productivity. This higher productivity, however, entails a high consumption of electric energy for heating and less favorable through-thickness microhardness and residualstress variations in the hardened layer. Figure 51 shows microhardness and residual-stress variations with the highest power density (Q1) and different heating times, tH1 . . . tH6 (Ref 15, 54). The influence of heating time affects the microhardness and residual-stress variations. With a correctly chosen condition of induction heating of the thin surface layer, adequate values of the maximum compressive residual stresses are obtained at the surface. The variation of residual stresses across the hardened surface layer, that is, in the transition zone between the hardened and the unhardened material, is more important. The variation should not be steep, which means that the gradient of residual stresses in the transition zone between the hardened and unhardened material should be very gently sloping. The microhardness variation in the transition zone, with a gradient of residual stresses as gently sloping as possible, can be accomplished by selecting the given power density and an appropriate heating time, since

Temperature cycles at the surface in induction surface heating and quenching at various power densities. Source: Ref 15, 54

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full use of compressive residual stresses in dynamically loaded parts can be made in this way. Figure 52 shows a cylindrical workpiece or sample with initial diameter DI and initial height HI. Its purpose is to show the process of residual stresses after induction surface treating in the central part of the workpiece (Ref 15, 27). With energy input to the thin surface layer, the initial diameter increases to the value of DA. Change of diameter is associated with thermal expansion of the material and is due to ferrite and/or pearlite transformed into austenite. After quenching, a thin surface layer with a martensite microstructure is obtained that has a different specific volume than the initial microstructure. The hardened surface layer is ready to receive a greater diameter (DH), which is resisted by the initial microstructure. Of interest are the residual stresses in the radial and axial directions. Residual stresses are compressive in the hardened layer, then they change into tensile. Residual stresses in the axial direction are compressive in the middle part of the hardened layer and turn into tensile below the hardened layer. In this way, very high tensile residual stresses are achieved at the bottom of the

Fig. 51

hardened layer. Compressive residual stresses prevail in the remaining part of the nonhardened surface layer. A residual-stress profile is unfavorable in the transition zone between the hardened layer and the rest of the nonhardened part of the workpiece. The transition from tensile to compressive residual stresses presents a serious danger for catastrophic failure. From the aforementioned it can be understood that the workpiece must be slowly heated and also quenched with the correct cooling rate. The actual cooling rate is very important throughout the martensitic transformation and must be as close as possible to the critical cooling rate. The quenching process must be carried out very carefully in the martensitic transformation temperature range to ensure internal stresses lower than the yield point. Figure 53 presents a crankshaft formed by hot forging, and Fig. 54 shows the manufacturing procedure from blank to crankshaft (Ref 15, 36). The procedure of forming should be carefully prescribed, including the initial and final temperatures of forging and the uniform plastic deformation rate for the entire volume.

Microhardness and residual-stress profiles at various heating times, tH1–tH4. Source: Ref 15, 54

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Fig. 52

Radial and axial residual-stress profiles after induction hardening the surface layer in the central part of a cylinder steel rod. Source: Ref 15, 27

This will ensure a fine-grained, banded microstructure contributing to a favorable relationship between strength and toughness of the product. The forging procedure is followed by annealing to remove the residual stresses incurred by the refinement procedure applied to forgings. Deformation of the crankshaft occurs after heat treatment, as early as in the phase of heating to the austenitization temperature. Due to stress relief as well as the effects of nonuniform cooling of the product, it is necessary for the forgings to be subjected to straightening prior to mechanical treatment. If necessary additional annealing can be prescribed to remove internal stresses induced by straightening. This is followed by turning and rough grinding to approach the final dimensions of the product. The technology of manufacturing the crankshaft involves careful selection of the conditions of turning and subsequent rough grinding to avoid the occurrence of internal stresses that would remain in the material even after induction surface hardening and grinding, thus reducing the fatigue strength of the material. Induction hardening may be preceded by stress annealing if the depth of the surface hardening is smaller than the depth of the damaged layer, since, in this way, it is possible to change the unfavorable stress state in the surface layer induced by machining. In this case, the depth of the induction surface

Fig. 53

Schematic presentation of a crankshaft with marked main bearing locations. Source: Ref 15, 36

hardening was greater than the depth of the damaged surface layer; therefore, machining could immediately be followed by induction surface hardening. After induction surface hardening, the size and distribution of residual stresses contribute to toughness and fatigue strength of the material. In these tests, induction surface hardening was followed by finish grinding and nondestructive magnetic inspection of the surface to reveal the possible existence of cracks on the product surface. Crankshafts were taken from production after induction surface hardening with the heat treatment and machining conditions as specified in the technology sheet. The residual stresses on the main crankshaft bearings were measured on the bearing location in the middle (sample A in Fig. 53), on the extreme left side (sample C) and on the extreme right side (sample G).

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Fig. 54

Machining and heat treatment procedure from blank to crankshaft. Source: Ref 15, 36

Figure 55 shows residual-stress distribution after induction surface hardening in the central bearing location (sample A) and on the extreme left side (sample C) (Ref 15, 20, 44). For both locations, residual stresses were measured on two samples. The distribution of residual stresses on location A is very similar on both samples, as expected, with the highest compressive stress ranging between 1020 and 1060 N/mm2 at a depth of approximately 250 mm and then slowly dropping to a depth of 3.5 mm. The residual-stress distribution after induction surface hardening on bearing location C is very similar to that in the central bearing location A, except that its absolute values are slightly lower, and a distinct decrease in the residual stresses can be noted as early as a depth of approximately 3 mm, reaching its minimum value at a depth of approximately 5.0 mm. The residual-stress distribution is just as favorable as in the central location, except that its absolute values are slightly lower. The difference in the residual-stress distribution can be related to the period of overheating in the austenitization temperature, which resulted in a thinner layer in austenitization and also a thinner hardened surface layer. Figure 56 shows a zone of the measured values of residual stresses at four bearing locations of the knee shaft (Ref 54). The arithmetic mean value for characteristic depths of the hardened layer was determined, and the highest and lowest values of residual stresses at the given depth were established, respectively. The upper and lower values of the residual stress with the individual depths of the hardened layer were defined by the range of scatter of residual-stress values. The upper and lower confidence limits were defined using a statistical

Fig. 55

Residual-stress profile after induction surface hardening on sample A of the mean bearing location in the middle of the crankshaft and on sample C on the extreme left side. Source: Ref 15, 20, 44

data analysis of the measured values of the residual stresses through the hardened-layer depth. Figure 57 shows a zone of the scattered values of residual stresses, the calculated variation of mean values in the individual depths, and the zone determined by the upper and lower confidence limits for the hardened-layer depth, to a depth of 5.5 mm (Ref 54). The zone between the upper and lower confidence limits is important and requires a very definite gradient of residual stresses. Where the limit, that is, the depth, should be set and where the gradient should be controlled depends on the chosen induction-heating conditions. The heating conditions applied in the study indicated that the specified gradient of residual stresses should be ensured at a depth ranging between 3.0 and 5.0 mm. A practical application showed that the stress gradient between the surface and a depth of

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3.0 mm was not problematic, since the values obtained lay in this range in all cases. Difficulties may be encountered in induction heating and induction hardening, respectively, if the residual-stress value obtained at the surface lies at the lower limit of the scattered results. This means that the measured compressive residual stresses are at the lower limit as well, which may result in a steeper gradient to the depth of 3.0 mm. Consequently, the required residualstress variation cannot be achieved through the entire hardened layer. The cylindrical specimen was surface induction heated to a temperature of 980 C and then quenched in saltwater. Figure 58 shows the calculated distribution of individual

microstructural phases from the surface to the center of the cylindrical specimen at the end of cooling (Ref 20, 47). The initial microstructure was preserved to a radius, r, of 5.5 mm. In the radii between 5.5 and 6.6 mm, the pearlitic-ferritic microstructure and low-carbon martensite appeared. With the radii exceeding 6.6 mm, a fine martensitic microstructure with approximately 6.0% residual austenite was obtained due to intensive cooling. With the selected conditions of induction heating and fast cooling of the specimen, no homogeneous martensite was formed in the surfacehardened layer. Figure 59 shows the calculated and measured variations of individual components of residual stresses (Ref 20, 47). The

Fig. 56

Residual-stress profiles for six measurements on four bearing locations after induction hardening. Source: Ref 54

Fig. 57

Determination of upper and lower confidence limit and arithmetic mean values of residual stresses through the hardened depth. Source: Ref 54

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calculated variations of residual stresses represent high compressive stresses at the surface, that is, the axial component of residual stresses, sz, equals 803 N/mm2, and the tangential one, sT, is 588 N/mm2. On the contrary, the tensile residual stresses were calculated after hardening in the core, with the preserved pearlitic-ferritic microstructure. Thus, the axial component of residual, stresses, sz, calculated for the core equalled to +370 N/mm2 and the tangential one, sT, was +62 N/mm2. The diagram in Fig. 59 indicates that the maximum tensile stresses were attained in the transition zone between the hardened layer and the unhardened one (Ref 47). At a greater depth, very low stress gradients occurred, and in the opposite direction, that is, in the thin surface layer to a depth of 2.5 mm, very high gradients of residual stresses occurred. The variations of residual stresses were determined experimentally by the x-ray diffraction method. The large gradient changes of the measured residual stresses in the thin surface layer can also be confirmed by measurement. The results of the measured and calculated variation of residual stresses in the surface-hardened layer sufficiently agree with small local deviations. To determine the local deviations of the variation and of residual stresses, numerous calculations were made with varying physical parameters of the material as well as different process parameters.

Fig. 58 Ref 20, 47

Calculated distribution of microstructures along the cylinder radius at the end of cooling. Source:

A particular problem with steels having ferritic-pearlitic and pearlitic-ferritic microstructures, respectively, is that heating of short duration does not ensure complete homogenization of austenite. Figure 60 shows the influence of inhomogeneity of austenite on the level of residual stresses, which is particularly noticeable in the martensite zone (Ref 20, 47). The initial inhomogeneous austenitic microstructure resulted in the appearance of a martensitic transformation with a small fraction of residual austenite. With regard to the volume fraction of martensite and residual austenite, plastification of the material occurred, which produced internal stresses and the variation of residual stresses, particularly in the thin surfacehardened layer. Figure 61 shows the calculated distribution of residual stresses due to heating of the specimen with heating rates of 200 and 800 C/s to a temperature of 1050 C, followed by cooling with a cooling rate of 1500 C/s (Ref 20, 47). With the high heating rate, 800 C/s, a very steep transition of residual stresses from the compressive to the tensile zone was obtained, which resulted in a decrease in fatigue strength. With the considerably lower heating rate, 200 C/s, the variation of residual stresses in the thin surface layer was essentially more favorable than that with the high heating rate. The variation of

Fig. 59

Calculated and measured residual-stress profiles after induction surface hardening. Source: Ref 20, 47

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the tangential and axial components of residual stresses permits the following observations:

With the lower heating rate, the residual stresses at the surface are lower by 100 to 200 N/mm2. The stress gradient for the tangential and axial components, st and sz, is very small in the subsurface, from a depth of 0.7 to 4.5 mm. The transition from compressive to tensile residual stresses does not occur before a

depth of 2.6 mm due to a small stress gradient. The radial component of residual stresses is 0 at the surface. In the subsurface, it is of the tensile character. Thus, it equals approximately 50 N/mm2 with the higher heating rate, 800 C/ s, and is 400 N/mm2 with the lower rate, 200 C/s. Figure 62 shows simulation of the variation of residual stresses with a favorable rate of surface induction heating, 200 C/s, which gives the

Fig. 60

Residual-stress profiles after induction surface hardening for heterogeneous and homogeneous austenite at austenitizing temperature. Source: Ref 20, 47

Fig. 61

Simulated residual-stress profiles at maximum surface temperature (Tmax = 1050 C) with various heating rates (VH1 = 200 C/s and VH2 = 800 C/s) and at a given cooling rate, VC, of 1500 C/s. Source: Ref 20, 47

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maximum temperature obtained at the surface, that is, 1050 C (Ref 20, 47). This was followed by quenching with two cooling rates, 1500 and 300 C/s. The variation of the tangential and axial components of residual stresses permits the following observations:

With the higher cooling rate, compressive residual stresses are obtained at the surface, that is, the axial component of 800 N/mm2 and the tangential component of 1050 N/mm2. With a low gradient, the axial and tangential components of the stresses vary; they change their sign to the tensile zone only at a depth between 2.6 and 2.8 mm. With the lower cooling rate, considerably lower compressive residual stresses are obtained at the surface, that is, an axial component of 275 N/mm2 and a tangential component of 390 N/mm2. A comparison of the axial and tangential components of residual stresses indicates that because of the considerably reduced cooling rate, the latter is lower in the thin surface layer by a factor of 3. The gradients of residual stresses with a higher or lower cooling rate are favorable, since a slow decrease of compressive to tensile residual stresses results in a minor susceptibility of a machine part to fatigue under dynamic loads.

Fig. 62

Residual stresses due to surface induction hardening of the gear wheels were calculated along the tooth surface during the heating process as well the quenching process. The series of images in Fig. 63(a) show distributions of internal stresses during the heating process after 20, 60, 100, and 187 s (Ref 52). The profile variation of internal stresses indicates that:

In the initial heating phase, the internal stresses in the tooth root are of the compressive character and reach up to 700 N/ mm2; with further heating, they change into tensile internal stresses ranging from +200 to +300 N/mm2. They are considerably lower in the zone reaching from the root to the pitch circle. Up to a heating time of 60 s, they are of the tensile character and reach up to 700 N/mm2. With further heating, they gradually change to compressive stresses attaining 100 N/mm2. They are obviously very low and insignificant in the upper part of the gear-wheel tooth, from the pitch-circle diameter to the tip of the tooth, therefore, they are not plotted. For this reason, deformations are the greatest at the tip of the gear-wheel tooth.

The series of images in Fig. 63 show changes of internal stresses during the heating process (Fig. 63a) and the cooling or quenching

Simulated residual-stress profiles at maximum surface temperature (Tmax = 1050 C) with heating rate, VH, of 200 C/s and at a various cooling rates (VC1 = 1500 C/s and VC2 = 300 C/s). Source: Ref 20, 47

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process (Fig. 63b) (Ref 20, 52). It is characteristic of the quenching process that considerably higher internal stresses occur and are more important in the tooth root. The figures show the magnitude and variation of internal stresses after the cooling times of 0.1, 1, 10, and 300 s, when the gear wheel is finally cool. The series of graphic representations of the internal stresses at the tooth surface indicate that:

In the initial quenching phase, at 0.1 and 1 s, the internal stresses are of the tensile character along the entire tooth height. They are the highest in the tooth root and gradually decrease toward the tip of the tooth. Between the quenching times of 1 and 10 s, the sign of the internal stresses changes, becoming gradually of the compressive character and ranging between 100 and 150 N/mm2 in the upper part of the tooth and up to 600 N/mm2 in the tooth root. At the end of quenching, after the cooling time of 300 s, the internal stresses change considerably only in the tooth root and attain up to 1500 N/mm2.

20 s

60 s

0.1 s

1s

100 s

187 s

10 s

300 s

The magnitude of residual stresses was also measured with the x-ray diffraction method and strain gages. Figure 64 shows the results of stress measurements in the pitch circle, that is, at the middle of the tooth surface for both teeth and in the tooth root where critical residual stresses occur (Ref 20, 52). By means of the x-ray diffraction method, it was found that:

The highest residual stresses occur in the root. At the surface they equal approximately 540 N/mm2, then increase to 750 N/ mm2 at a depth of 30 mm to reach 870 N/ mm2 at a depth of 60 mm. The residual stresses are a bit lower in the middle of the tooth surface, that is, at the pitch circle. At the surface, they equal approximately 90 N/mm2, then gradually increase with a greater depth to reach 400 N/ mm2 at a depth of 60 mm.

A general conclusion can be drawn that the experimental as well as theoretical results agree very well and provide useful information, especially to a technologist in the manufacture of gear wheels. Figure 65 shows a part of a gear in cross section, with residual stress distribution in the surface-hardened layer of the gear tooth (Ref 15, 20). Also in this case, a residual-stress distribution typical of induction hardening is obtained, that is, compressive residual stresses in the hardened layer with a martensite microstructure, followed by tensile residual stresses, and, at greater depth, compressive residual stresses once again. The size and distribution of residual stresses can be influenced by power density.

(a)

Stress σ 0 (b)

daN/mm2

50

Fig. 63 Ref 20, 52

Principal stress during (a) induction surface heating process and (b) quenching process. Source:

Fig. 64

Residual-stress measurements on the root below the tooth surface. Source: Ref 20, 52

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A number of authors have done research into the effects of power density or energy input used in heating, with the purpose of determining optimal induction-hardening conditions that lead to good mechanical properties of gears. Lower power density at the same current frequency requires longer heating times at the same depth of the hardened layer and results in higher compressive residual stresses with a moderate transition of residual stresses from the compressive into the tensile region. In induction hardening of gears, it is necessary to ensure the most uniform depth of the hardened layer to achieve a symmetric distribution of residual stresses in the gear tooth cross section.

Workpiece Distortion in Induction Surface Hardening Dimensional changes of the workpiece are closely related to the internal stresses occurring in heating and cooling. During heat treatment, when internal stresses in the workpiece are higher than the yield point, distortion of the workpiece takes place. When the workpiece has an axisymmetric shape, it is very important to

Fig. 65

Residual-stress distribution in the induction surfacehardened layer of the gear tooth. Source: Ref 15, 20

place the induction coil so that symmetric heating and quenching is achieved, resulting in uniform thickness of the hardened layer with a martensitic microstructure. Figure 66 shows induction surface hardening of a cylindrical rod and tube. When the gap between the coil and the workpiece varied from the maximum gap, r1, to the minimum gap, r2, that is, r14r2, the hardened layer thickness varied also (Ref 15, 27). This means that the side of the workpiece with a thicker hardened layer will suffer greater changes in volume than the side with the thin layer. A changing thickness of the hardened layer will result in distortion or curvature of the cylindrical rod in the direction of the thinner induction-hardened layer. Figure 66(a) shows a cylindrical tube with an induction-hardened inner surface. Asymmetrical placement of the coil inside the tube (hole, bore) causes differences in the gap between the coil and the bore surface inside the workpiece. The conditions of heating are very similar to those mentioned earlier, and the cylindrical rod bends in the direction of the thicker hardened layer. Similar distortion happens on prismatic bodies where the specimen dimensions are a · h · l. Distortion of workpieces subsequent to induction surface hardening on rectangular crosssectional rods, strips, plates, and similar forms is defined or estimated with respect to the size of the object and thickness of the hardened surface layer (Ref 27). In unilateral surface hardening of prismatic bodies, the distortion will be greater with a smaller height (h) and greater length (l). Distortion also grows with the depth of the hardened surface layer. Figure 67 shows distortion or, more precisely, bending of a steel rod with a rectangular cross section, where the rod has the same cross section a · h and the same length, l (Ref 15). Figure 67(a) presents specimens of equal size and shape that have been induction hardened to different depths. For an ideal comparison, choose heat treatment conditions that will compare easily, that is, equal power (P) and frequency (f ), and change the time of heating. Changing the induction heating time allows the identification of different depths of the hardened layer and very similar microhardness or microstructure in the transition zone. Figure 67(b) shows the conditions in induction surface heating when the temperature transformation is not exceeded, TA1 . This means that on the completion of quenching, the hardened layer has not

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Fig. 66

Influence of nonuniform thickness of surface-hardened layer on distortion for cylindrical steel rod and tube. Source: Ref 15, 27

Fig. 67

Bending of a steel rod of rectangular cross section as a function of the thickness and location of the hardened layer. Source: Ref 15

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been reached (d0B = 0), and this is reflected on only a partial depth of the overheated layer, d00B. Despite surface quenching, volume changes in the surface layer are not affected, which means that there is no distortion of the workpiece, so the size of bending is f = 0. Bending distortion may be noted quite rarely, and only if very intensive cooling is applied with temperature differences between the surface and subsurface defined by the depth at which the temperature in the specimen is equal to the temperature of the environment. However, such cases can be neglected, since the goal is to achieve the austenitization and homogenization temperature, which is much higher than the temperature of the ferrite/pearlite-to-austenite transformation, TA3 . In practice, the cooling rate is adapted to the quenching medium, which ensures the formation of a hardened microstructure with minimum internal stresses. Figure 67(c) illustrates conditions after induction surface hardening achieves a very small depth of the hardened layer, d0C , and a correspondingly small depth of the heat-affected zone. The achieved depth of penetration (conduction) of heat after induction hardening is equal to d0C +d00C =dC . The volume changes in the specimen are influenced only by the hardened layer, that is, the volume of the hardened layer versus the entire volume of the specimen. If the volume changes are small, then the bending of the specimen is negligibly small. For a given hardened layer depth, it is possible at each point during the cooling process to ensure a higher yield point of the material (Rp, 0:2 ) than the internal stress in the axial direction of the specimen. When the depth of the hardened layer is equal to or smaller than one-twentieth of the 0 41/ h, there is no bendspecimen height, dC 20 ing deformation of the specimen, only a slight increase in residual stresses in the axial direction, despite a minimum volume change of the hardened surface layer. Figure 67(d) illustrates the conditions after induction surface hardening of a specimen with a slightly greater depth of the hardened layer, dC0 , than on the specimen in Fig. 67(b), dC4dB. The achieved depth of penetration (heat conduction) after induction heating is also slightly greater and is dC or dC4dB. If the achieved hardenedlayer thickness is dC41/10 h, then hardening is accompanied by deformation of the specimen, described as bending deformation in the middle of the specimen ( f40). Surface hardening is followed first by annealing to remove the

stresses and then by mechanical straightening and finally grinding. The procedures of subsequent treatment of machine components are chosen mainly with respect to the size of the bending deformation. Figure 67(e) shows the conditions after induction surface hardening of the specimen on both sides. Several different procedures can be chosen to ensure two-sided hardening. The first procedure involves hardening of the top surface to a certain depth, d0E , and then of the bottom surface to the same depth. In the second procedure, both top and bottom surface are hardened at the same time. The third procedure involves hardening of the top surface from the left to the right side and the bottom surface from the right to the left. The most suitable procedure is the second one, where the top and bottom surfaces are hardened simultaneously to the same depth. The achieved depth 00 of heating is equal to dE , whereas the depth of the hardened layer is equal to d0E . When equal heat treatment conditions are ensured on both sides of the specimen, the depth of the hardened layer on the top surface, dEU, is equal to the depth of the hardened surface on the bottom surface, dED. This induction surface-hardening procedure ensures that the bending deformation in the middle of the specimen is equal to zero ( f = 0). Figure 68 shows two crankshafts that differ in the position of the crankpin with respect to the main journal. The first example presents a crankshaft where the distance, a, between the axes of the crankpin and the main journal is greater than the sum of the radii of both the crankpin and the journal: a4

dGL dKL + 2 2

In the second example, the crankshaft is slightly different, having the distance between the two axes equal to the sum of the radii of the crankpin and journal: a=

dGL dKL + 2 2

This must be considered in the analysis of the deformation of the crankshaft after hardening cylindrical parts that alternate in following one another (crankpin and main journal). For the conditions after induction surface hardening of the crankshaft described in the first

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example, Fig. 68(a) shows induction surface hardening of the crankpin, which causes the deformation of the main journal (Ref 15). Then follows induction surface hardening of both main journals on the left and right side. Due to tensile residual stresses in the web between both main journals, the web sections come closer to one another, and the length of this crankshaft section shortens by 2Dl. Considering the number of crankpins or the number of main journals, the entire length of the crankshaft decreases by Dl. This value depends on the number of the main journals, N, and the size of contraction per one embedment (Dl). The entire contraction of the crankshaft is thus DLC = N Dl. In the second example (Fig. 68b), the crankshaft is designed differently, with the axial distance between the crankpin and the main journal being equal to the sum of both radii (Ref 15). As in the previous case, induction surface hardening of the crankpin induces tensile stresses in the web, causing distortion of the main journal. Both adjacent main journals come closer to one another, but the distortion is smaller than in the first case. After induction surface hardening of both adjacent main journals, any temporary distortion that may have occurred vanishes, and the crankshaft has equal length, just as prior to heat treatment. Research has confirmed that the stress state in the workpiece is more favorable

Fig. 68

and distortions are smaller in progressive hardening than in single-shot hardening. The only problem is whether the size and shape of the workpiece allow the application of progressive hardening. Fujio et al. (Ref 52), in their third report on induction hardening, focused on studies of distortion of gear-wheel teeth and residual stresses in gear wheels. The authors measured the outer diameter and the root diameter of the gear wheel across two opposite teeth and the root parts of the teeth with a micrometer before and after quenching. The same dimensions were also calculated theoretically, taking into account the volume changes due to phase transformations, and serve as a basis for determination for deviations. Figure 69 shows a change of the outer diameter of the gear-wheel tooth for gear wheels 2, 3, and 4 (Ref 20, 52). The gear wheel has 26 teeth, which means that the deviation was measured between the first and twenty-sixth tooth, the third and sixteenth, and so on, so that thirteen measurements were made in total. The results of the measurements are shown as points in the diagram. The dotted line represents the theoretically calculated increase in diameter due to heat treatment. Figure 70 shows a change of tooth height for gear wheels 2, 3, and 4, as in the previous

Crankshaft distortion after surface induction hardening of individual journal locations. Source: Ref 15

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example (Ref 20, 52). The dotted lines in the individual diagrams represent the theoretical deviations of height of the individual teeth by 0.18 mm and the deviations found by measurement, which are mainly greater than the theoretical values. The greatest deviation can be found with gear wheel 4. It exceeds 0.2 mm. Figure 71 shows changes of the root diameter. Calculations show that after heat treatment, the root diameter decreases by approximately 0.02 mm (Ref 20, 52). The measured values of the same diameter, however, remained

Fig. 69

Change of tip gear-wheel diameter after quenching. Source: Ref 20, 52

Fig. 70

Change of whole depth of gear wheel after quenching. Source: Ref 20, 52

unchanged for gear wheels 2 and 3 and decreased by 0.1 mm maximum for gear wheel 4. A comparison of the outer tooth diameter and the tooth root diameter, also taking into account the tooth height, shows that distortion of individual gear wheels is considerably more complicated and cannot be described by the selected measurement methods; therefore, changes of tooth profiles along with tooth height were measured. Figure 72 shows tooth profile error curves after induction surface heating and after quenching (Ref 20, 52). Prior to quenching, the left tooth surface was marked with the letter “a” and the right tooth surface with the letter “b.” Measurements of the height profile are shown for two teeth, those marked 1 and 14, of gear wheels 2 and 3. Considering that the tooth height is 10 mm and the deviations are plotted in millimeters, their absolute value can be evaluated. It ranges between 9 and 100 mm at each tooth surface concerned. One method of induction surface hardening appropriate for large gear modules is known as gap-by-gap hardening. It belongs to the progressive hardening methods. In this case, the coil is placed so that it ensures a uniform gap between the coil and the flanks of two adjacent teeth. The tooth gap-hardening method is very demanding and requires much experience and knowledge to achieve the desirable properties of the gear. This method is also known as contour

Fig. 71

Change of root gear-wheel meter after quenching. Source: Ref 20, 52

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Fig. 72

Tooth profile error curves (a) after induction surface heating and (b) after quenching. Source: Ref 20, 52

hardening. It is an ideal method for heat treatment of gears because it increases the hardness on the tooth surface only slightly and decreases the load-bearing capacity in the tooth root. Gears heat treated in this way exhibit very good behavior in operation, because compressive residual stresses are present in the root of the tooth. Gears with induction-hardened flanks, given that the dimensioning is carefully carried out, can achieve the highest fatigue strength. To verify the results of induction surface hardening, it is necessary to discuss certain measures for controlling the quality of the hardened layer. For this purpose, hardness and microhardness measurements supported by microstructural analysis are commonly used. A disadvantage of this procedure is that, due to the method of heating and quenching (nonuniformly overheated left and right tooth flanks), slightly higher dimensional deviations may be obtained than in simultaneous hardening of both flanks of the same tooth (Ref 27). Figure 73 shows the deviation in the dimensions of the tooth after induction surface hardening by heating with a coil that encircles the gear tooth (Ref 15, 20, 55). Measurements of gear teeth and gear gaps after induction surface hardening show an increased volume in the tooth root and thus a smaller gap and increased volume at the tip of the tooth (increased gear diameter). These

Fig. 73

Distortion of individual tooth shape after induction hardening caused by volume changes. Source: Ref 15, 20, 55

volume changes result in a slightly smaller width of the tooth above the pitch circle and slightly increased tooth width below the pitch circle of the gear. These dimensional deviations are relatively small and negligible in the case of gears with small diameter or small module, but may become more important in gears with larger modules and greater tooth width (Ref 55, 56). Surface hardening of gears is one of the most frequent applications of induction surface hardening. In the first part of this chapter, different methods of induction heating of gears was discussed, depending on gear size but also on the properties expected of gears after heat treatment.

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Gears are axisymmetric machine elements; therefore, volume changes due to the transformation into a martensitic microstructure induced by heat treatment are to be expected. Teeth themselves are symmetrical, which means that careful application of uniform heating of particular gear teeth and uniform quenching, regardless of the method applied, should not cause distortion after heat treatment. In induction surface hardening, compressive residual stresses are created in the surface, which, in addition to hardness and wear resistance of the surface, create a considerably high fatigue strength and resistance to bending loads. Figure 74 shows shape distortion after singlefrequency and dual-frequency induction surface hardening (Ref 20, 57). Figure 74(a) shows a simulation of distortion by showing the initial shape of the tooth half-cut prior to hardening (dotted line) and the modified shape after dualfrequency induction surface hardening (continuous line). Figure 74(b) shows a similar simulation of distortion but for conventional single-frequency induction surface hardening. In Fig. 74, the calculated actual size of distortion is 50 mm, with an equivalent geometrical scale of 2.0 mm. A comparison of both cases of hardening from the viewpoint of tooth distortion clearly shows that distortion in single-frequency surface induction hardening is twice that obtained after dual-frequency hardening. This difference in tooth distortion can be attributed to the fact that in single-frequency

surface induction hardening, the martensitic transformation occurs in the whole tooth. Consequently, the volume change of the tooth is greater and so is the distortion. In dualfrequency surface induction hardening, only the tooth contour becomes hardened. Consequently, the volume change of the tooth is smaller and so is the distortion, that is, it equals only half of that obtained after single-frequency hardening.

Residual Stresses after Induction Surface Hardening and Finish Grinding The last phase in the manufacturing of crankshafts is fine grinding, where achieving the desirable condition of the surface and the surface layer requires:

Suitable dimensions of the particular bearing locations with respect to the allowable deviations Suitable surface roughness Grinding stresses are compressive or the lowest tensile to maintain the favorable stress profile obtained by induction surface hardening of the surface layer Smallest changes possible in the microstructure and also smallest changes in the hardness and microhardness profiles in the heat-affected zone after grinding

How is it possible to assure a desirable surface and surface layer quality after induction surface hardening and fine grinding? Finding an answer to this question requires a very good knowledge of the process of grinding on the microlevel as well as all mechanical and heat effects acting on the layer of the workpiece, including the type and condition of the grinding wheel. An allinclusive consideration of the numerous influences of the kind and condition of the tool on the changes on the surface and in the surface layer of the workpiece in the given machining conditions can be based on the descriptions of surface integrity (Ref 58–63). For the grinding process, the following conditions have been selected: Fig. 74 Ref 20, 57

Distortion after induction surface hardening with (a) dual frequency and (b) single frequency. Source:

Different kinds of grinding Different grinding conditions normal, abusive)

(gentle,

Because of thermomechanical loads in the thin surface layer during the grinding process,


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very complex physical and chemical processes take place on the microlevel. For an accurate description of the conditions in the contact zone for machining a given material, it is necessary to consider the grinding method as well as the kind of material from which the grinding wheel is made, its structure, and the grinding conditions. The heat conditions in grinding are a result of the contact conditions between the individual grinding grains and the workpiece material as well as deformation work in chip formation in the shear zone. The amount of generated heat is strongly dependent on the chosen machining conditions and is abducted mainly through the chip, while a smaller part of the generated heat is transferred through heat transfer into the thin surface layer of the workpiece. Heating of the chip does not cause any particular difficulties, but heating of the thin surface layer of the workpiece creates the conditions for different mechanical and thermokinetic processes, which cause microchemical changes. Heating up the thin surface layer of the workpiece can leave certain undesirable effects that change the properties of the part surface layer and thus harm its operational abilities. The generated friction heat is transferred through the heat-transfer phenomenon from the contact between the grinding grain and the workpiece into the grinding grain. The increased amount of heat on the grinding grain or grinding wheel intensifies the wear processes and damage of particular grinding grains, which follow very different and complex mechanisms and, in the final phase, affect the serviceability and operational life of the tool. The grinding tool consists of grinding grains, each representing a process. Grinding grains are interconnected with an appropriate binder, which is defined by different degrees of porosity of the grinding wheel structure. Therefore, it is necessary to also know the force needed to break the grinding grain from a given structure, expressed by the grinding wheel hardness. Simultaneous changes in volume proportions of the grinding grains and binder can create different structures on the grinding wheel that do not behave in the same way. This means that by changing the kind of material for the grinding wheel and the binder, it is possible to achieve equal effects by changing the structure of the grinding wheel. It follows that by a suitable combination of influential parameters, it is possible to achieve a longer life and wear resistance of the wheel in equal kinematic conditions of the wheel and the workpiece. The wear of the

grinding grains is a result of mechanical and thermal effects, which are reflected in reduced cuttability. Figure 75 shows the basic forms of wear and damage on the grinding wheel grains, expressed by characteristic changes on the grinding grains (Ref 15, 64). Due to mechanical loads on particular grinding grains, short-lived but intensive heat effects are created in the contact or on the friction surfaces with the workpiece. In these cases, blunting of the grinding grains (Fig. 75a), breaking off of the grinding grains (Fig. 75b), or splitting of the grinding grains (Fig. 75c) may happen. Requirements are often set that a worn-out grain should fall from the grinding wheel at a certain moment. A wornout grain typically has an increased contact surface, which causes the forces on the grinding grain during the cutting process to be higher than the binding forces between the grains, and therefore, the grain falls out (Fig. 75d). It is expected that the manufacturing engineer will carefully choose all the parameters of the process so that the worn-down grains fall out. This ensures more efficient machining and reduced thermomechanical effects in the thin surface layer of the workpiece material. To create a suitable quality of the workpiece surface and surface layer, it is necessary to ensure that grinding grains will fall out, and that the activation of new, sharp grinding grains is made possible. It is very important to choose a cutting condition that would make the grains fall out only when their cuttability is reduced. Figure 75(e) illustrates characteristic blunting of the grinding grain due to chemical reactions at high temperatures. Chemical reactions at high temperatures are frequently followed by filling the pores with overheated, highly plastic chips of the workpiece material (Fig. 75f). When the pores

Fig. 75

Basic forms of wear and damage on the grinding wheel grain. Source: Ref 15, 64

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on the wheel surface become filled, the temperature in the contact zone rises and changes the progress of the temperature cycle into the depth of the workpiece. The maximum temperature of the temperature cycle on the surface rises, as does the temperature in the particular depths of the workpiece material. This results in greater depth of the heat-affected zone in the material, which may have fatal consequences in terms of surface layer properties of the workpiece material. Generally, blunting of the grinding grains makes the contact surfaces between the grains and the workpiece larger, which creates the conditions for increased mechanical effects in the contact zone and higher heat input, accompanied by stronger heat effects, in the surface layer of the workpiece. A plastically deformed layer is created because of interaction between the grinding tool and the workpiece on the place where the chip separates from the base. A result is hardening of the thin surface layer of the workpiece material and occurrence of internal stresses that may lead to failure of the thin surface layer and/or deformation of the workpiece, with the presence of residual stresses at the end of machining. Macro- and microanalysis with optical and/or electronic microscopes shows microcracks and/ or other damage on the surface caused by an inadequate grinding method or procedure. Microscopic assessment of the surface state and damage on the surface quite often points to inadequate grinding conditions. The most frequent surface damage includes hollows, mars, torn-off areas, built-up edges of the workpiece or the tool, and so on. It is often necessary to consider the generated heat effects that cause microstructural and/or chemical changes accompanied by dimensional changes. The damage on the workpiece surface should be taken very seriously, since this may give rise to very detrimental friction conditions during operation with another element in the mating pair. In the analysis of microstructural changes in a thin surface layer of the material after finish grinding, it is possible to evaluate the size of the heat-affected layer. Figure 76 shows the time variation of the maximum temperature on the surface at the particular depths with respect to the workpiece speed, VW (work speed, VW) (Ref 15, 65, 66). By knowing the melting temperature and the austenitization temperature of the discussed steel, the depths of the remelted layer and the heat-affected layer can be defined. Under

different grinding conditions, different temperature cycles were obtained on the surface and in the depth of the heat-affected layer, which has effected microstructural changes and changes in the microhardness and residual stresses (Ref 64–66). Figure 77 presents the temperature cycles on the specimen surface as a function of depth in the hardened steel at given grinding conditions (Ref 15, 65). The temperature cycles can be treated separately as a heating phase and then a cooling phase. The maximum temperatures achieved at the surface and in the surface layer, respectively, are also very important (Ref 64). A distinction can be made between three characteristic cases of temperature cycles:

The maximum temperatures at the surface and in the surface layer, respectively, are higher than the melting temperature of the specimen material, depending on the temperature cycles (Fig. 78a) (Ref 65). Such conditions may occur due to very sharp grinding conditions or the selection of an inappropriate grinding wheel with regard to the specimen material. The depth of the remelted layer is only a few micrometers and makes a very fine ledeburite microstructure containing fine cementite spread in residual austenite. The newly formed microstructure has a slightly lower hardness than martensite. The residual stresses in the thin surface layer will be tensile, due to plastic deformation of the surface layer in grinding caused by tensile forces in the contact zone of the specimen material. To this should be

Fig. 76

Maximum temperature drop as a function of depth in the hardened steel during grinding with various work speeds, Vw. Source: Ref 15, 65, 66

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added the tensile stresses induced by the occurrence of residual austenite. The maximum temperature in the contact zone is lower than the temperature required for the beginning of melting of the given material and higher than the austenitizing temperature. The lower temperature, that is, the austenitizing temperature, will shift because of a high rate of heating of the specimen toward higher temperatures, as known from transformation diagrams (Fig. 78b) (Ref 65). Provided the previous microstructure of the surface layer was martensite-cementite-carbide, a finer martensitic microstructure with a higher carbon content, obtained at the expense of cementite-carbide phases and with the possibility of a lower content of residual austenite, may be expected in the thin surface layer after grinding. The modified content of the cementite-carbide phase depends on the heating conditions, whereas the content of the residual austenite depends on the cooling conditions.

Fig. 77

Grinding temperature cycles in different depths in the hardened steel at given grinding conditions. Source: Ref 15, 65

Fig. 78

The maximum temperature in the contact zone is lower than the temperature required for the beginning of austenitization and higher than the lower temperature that is limited by the temperature of steel tempering. It is equal to approximately 200 C (Fig. 78c) (Ref 65). The grinding conditions are very mild, so that with the selection of the correct type of grinding wheel, no important changes are expected in the surface layer. Only martensitic tempering may occur in the surface layer if it was not performed during the induction surface hardening of the specimen.

The engineer should be aware that surface integrity depends on the tribological conditions in the operation of a component/assembly. Therefore, adequate knowledge for the assessment of tribological conditions of a component in operation is very important for the prescription of machining that would result in the desired surface condition and subsurface layer. In general, there are two tribological systems. The first is the one present during machining, and the second is acting in operation. In both tribological systems, however, it is the properties of the workpiece material and its condition prior to and after machining that play a key role. Under different machining conditions of grinding, different temperature cycles were obtained on the surface and in the depth of the heat-affected zone, which has effected microstructural changes and changes in the microhardness and residual stresses. Thus, on the surface, a maximum temperature higher than the temperature of melting of the workpiece material was obtained. The depth of the remelted layer is only a few micrometers and makes a very fine ledeburite microstructure containing fine cementite spread in residual austenite. The newly formed

Maximum temperature drop as a function of depth in the induction surface-hardened steel at various speed, Vw. Source: Ref 65

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microstructure has a slightly lower hardness than martensite. The residual stresses in the thin surface layer will be tensile, due to plastic deformation of the surface layer in grinding caused by tensile forces in the contact zone of the workpiece material as well as the tensile stresses induced by the occurrence of residual austenite. The relative grinding stress is obtained by measuring the residual stress after induction surface hardening, followed by measuring the same spot after induction hardening and grinding, then calculating their difference. Figure 79 shows the measured absolute residual-stress profile after induction surface hardening and the measured residual-stress profile after induction surface hardening and grinding (Ref 15). A relative grinding stress represents the difference between measured residual stress on specimens that were induction surface hardened and induction surface hardened and grinded. A relative grinding tensile stress equals +425 N/mm2, and then the sign changes in the depth at approximately 175 mm, as shown in Fig. 80. Figure 80 shows the absolute residual-stress profile after induction surface hardening and grinding on bearing location “A” as well as the average residual stress after induction surface hardening. From this, the relative grinding stress can be calculated (Ref 15). Figure 81 shows a zone of relative grinding stresses occurring under gentle grinding conditions (Ref 54). With gentle grinding conditions, compressive relative grinding stresses will prevail at the

Fig. 79

Subsurface residual-stress profile after induction surface hardening and grinding (absolute stress) on bearing location “A”. Source: Ref 15

surface. It is only in the subsurface that the tensile influence can be felt. The relative grinding stresses are low and practically negligible, although they exert a positive influence on the original residual stresses after induction hardening. The zone was defined by the upper limiting values characteristic of greater depths under normal grinding conditions and by the lower limiting values characteristic of gentle grinding conditions. Generally, care should be taken when choosing grinders and grinding conditions. The chosen grinder should ensure regeneration of the abrasive grains, so that the worn-out grains may fall off and the new ones grind the surface. The selected grinding conditions should ensure that a suitable force be applied to the worn-out abrasive grains so that they may fall off and new, sharp abrasive grains become active. Figure 82

Fig. 80

Subsurface profile of relative grinding stress on bearing location “A”. Source: Ref 15

Fig. 81

Relative grinding stress region through thin surface layer after gentle grinding conditions. Source: Ref 54

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shows a zone of relative grinding stresses occurring under abusive or normal grinding conditions (Ref 54). The zone of grinding stresses was determined from the profile of residual stresses after induction hardening and also after induction hardening and grinding (Ref 64). By deducting the two residual stresses, the socalled relative grinding stresses were determined. The abusive and normal grinding conditions produce remelting of the thin surface layer, that is, its rehardening. With the abusive and normal grinding conditions, the tensile relative grinding stresses are always predominant at the surface. At greater depths, there is still the influence of remelting under the abusive grinding conditions, whereas rehardening occurs under normal grinding conditions. For the sake of clearness, the zone was defined by the upper limit of residual-stress values, which is characteristic of the abusive grinding conditions, and the lower limit, which is characteristic of the normal grinding conditions (Ref 64). The results confirm as predominant the residual stresses, stresses induced by the plastic deformation of the material, and as a lesser influence, tensile stresses caused by the formation of residual austenite. On the basis of the measurements of residual stresses after induction surface hardening or induction surface hardening and grinding, it can be concluded:

For residual stresses after induction surface hardening and grinding, the conditions of abusive grinding are a more favorable choice. They lower to a lesser extent the desirable compressive residual stresses after induction surface hardening.

Grinding conditions can be chosen so that the melting temperature of the workpiece material (gentle grinding conditions) is not exceeded. Then, the favorable compressive stresses after induction surface hardening are lowered due to plastic deformation of the workpiece material during the process, and thus, relatively low tensile residual stresses are obtained. However, this will significantly lower the productivity. Special attention should be paid to the selection of the type of grinding wheel in terms of grinding wheel material, binding agent, hardness, and pore density, since a correct selection can contribute to higher cutting efficiency concerning the plastic deformation of the workpiece material. In this way, the grinding tensile stresses are kept as low as possible and make the compressive residual stresses induced by induction surface hardening of the prevailing variety.

Induction surface hardening creates a very desirable residual stress state. Residual stresses are always of a compressive nature and are usually present to the depth of the induction surface-hardened layer. However, a major difficulty in induction surface hardening is ensuring a very slight/slow variation in microhardness and the existence of compressive residual stresses in transition areas to the microhardness of the core material. By gently grinding and varying the hardness and existence of compressive stresses in the transition area, it is possible to diminish the notch effect induced by stress concentration. Additional grinding of an induction surface-hardened surface deteriorates the stress state in the surface layer, since grinding has always induced tensile stresses. By the correct selection of machining conditions and grinding wheel properties, the engineer will contribute to less tensile residual stresses and will avoid deteriorating the favorable residual-stress state after induction surface hardening (Ref 44, 59).

Hardness Profiles in the Induction Surface-Hardened Layer

Fig. 82

Relative grinding stress region through thin surface layer after abusive or normal grinding conditions. Source: Ref 54

The induction surface-hardened layer was analyzed by measuring the hardness and microhardness and their relation to the microstructure (Ref 44). To establish some relationships between microstructural changes versus

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hardness across the entire hardened layer and the microhardness in the very thin surface layer, the latter were measured by the Vickers method. Thus, Fig. 83 and 84 show the hardness profiles versus the induction-hardened layer depth according to Vickers at a load of 20 N or measurements of microhardness according to Vickers in a very thin surface layer to a depth of 150 mm at a load of 2 N (Ref 44). The hardness of the surface layer after induction surface hardening is very uniform in all the investigated main bearing locations and equals approximately 520 to 550 HV2.0. The hardness profile highly conforms to the residualstress profile, which is confirmed by a decrease in hardness in the transition zone. The hardness profile in the transition area is likewise very steep and points to high stress concentrations in this location when the crankshaft is in the loaded state.

Figure 83 presents the hardness profile in an induction surface-hardened layer to a depth of 5.0 mm and the microhardness profile to a depth of 150 mm on bearing location “A” (Ref 44). The hardness measurements show that the quenched and tempered steel has a hardness of approximately 220 to 260 HV2.0, and the surface-hardened layer has a hardness of approximately 540 HV2.0. In the surface-hardened layer, a slight increase in hardness as a function of depth is evident, which is conditioned by microstructural differences due to varied cooling rates of the surface layer. The surface is cooled under the effects imposed by the cooling medium, yet at a greater depth, the effects of the medium are accompanied by a more expressed effect of the cold mass of the core, resulting in the formation of very fine martensite and greater hardness at greater depth (Ref 64).

Fig. 83

Hardness profile in the induction surface-hardened layer and microhardness profile in a very thin surface layer for bearing location “A”. Source: Ref 44

Fig. 84

Hardness profile in the induction surface-hardened layer and microhardness profile in a very thin surface layer for bearing location “C”. Source: Ref 44

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Kosel and Kosec (Ref 67) conducted an investigation on the cracks formed at the surface-hardened and ground bearing location of a crankshaft. Crankshafts are made of highquality chromium-molybdenum heat treatable steel. The shafts were forged from square shafts with a cross-sectional area of 120 by 140 mm. In the course of mechanical and heat treatments, the workpieces were tested with various nondestructive methods, for example, visual inspection, and, if required, magnetic particle and/or penetrant inspections. A peculiarity of shaft production is that during forging, the material flows from the middle of a slab ingot to the outer section of a forging due to plastic deformation. The steel in the middle is usually of lower quality than the steel at the surface of a cast iron-works blank and that of a formed steelworks semiproduct. The character of forming the steel-works semiproduct tool, steel-works blank, and crankshaft was such that a reverse material flow occurred from the middle to the surface. The authors reported the presence of weblike cracks at the surface of the crankshaft bearing location (Fig. 85) and gave two reasons for the occurrence of cracks immediately after forging (Ref 67):

Inhomogeneity, that is, nonuniform crosssectional chemical composition of steel Inadequate conditions of bearing-location grinding

After macrostructural and microstructural examinations as well as a microchemical analysis were performed, it was found that, from the metallurgical point of view, the forging showed quality, and the defect may be attributed to the grinding process alone. Figure 86 shows a macrosection of the forging at the bearing, and

Fig. 85

Weblike surface cracks at bearing location of crankshaft. Source: Ref 67

Fig. 87 shows a forging made of low-quality slab ingot (Ref 67). The main purpose of the investigation was to find the cause of the weblike cracks after induction surface hardening and the final grinding on the bearing location on the crankshaft. Grinding was studied in the same way as at the cut-out bearing locations. From the steelworks slab ingots prepared for forging, specimens of a suitable length were cut out and turned to size.

Fig. 86

Macrograph of etched favorable billet cross section. Source: Ref 67

Fig. 87

Macrograph of etched unfavorable billet cross section. Source: Ref 67

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The specimens prepared in this way were hardened and tempered at 560 C and then surface hardened, with a maximum surface temperature ranging between 820 and 900 C. The case depth ranged between 3 and 4 mm. The macrograph in Fig. 88 clearly identifies the bearing location and shows a considerable difference in case depth. Generally, it was found that:

There were insignificant differences in chemical compositions. There were segregations in the central part of the crankshaft. There were shrinkage cavities mainly on the inside and also on the outside part of the forging. There were microsegregations. Some of the hardened bearing locations showed dimensional variations with reference to the anticipated thickness. The microhardness obtained at the surface was within the expectations and somewhat above a value of 700 HV2.0. The depth of the hardened bearing locations ranged between 2.0 and 2.5 mm. No differences in the microstructures could be found using a common optical microscope analysis or a scanning electron microscope. The microhardness at the cracked cylindrical specimens was somewhat lower, approximately 600 HV2.0.

Figure 91 shows two temperature cracks, with a pore at the crack location being visible as well. An analytical model was elaborated to clearly demonstrate the causes of crack formation at the hardened surface after grinding (Ref 67). The stress condition, sR, occurring at the hardened-layer surface after heat treatment and grinding was known. In several cases, the surface-hardened layers took the shape of an eccentric ring, which means a nonuniform thickness of the hardened layer. The cracks occurred at the location of the smallest layer thickness. As expected, residual stress occurs at the specimen surface where cracks initiate when stress occurs due to phase transformation to martensite, thermal stresses during hardening, stresses due to martensitic tempering during grinding, and thermal stress after grinding (Ref 67). In the phase transformation from austenite to martensite, steel volume will increase. The volume strain can be calculated using the equation: eV =

VM 7VA VA

where VM is the volume of an elementary cell of martensite, which can be calculated for the given steel: VM =(2:861 0:013% C)2 (2:861+0:116% C) =23:7262 103 nm3

Figure 89 shows the microhardness profile of the cross section of an induction-hardened layer for a favorable bearing location, and Fig. 90 shows the profile for an unfavorable bearing location without cracks (Ref 67).

Fig. 88

Macroscopic examination of case depth at cross section of crankshaft bearing. Source: Ref 67

Fig. 89

Hardness characteristic at cross section of weblike cracked neck as a function of depth, z. Source: Ref 67

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and VA is the volume of an elementary cell of austenite of the same steel: VA =0:5(3:548+0:044% C)3=22:68242 103 nm3

then eV = 0.046. A stress condition with such a volume strain would exceed the material strength. Table 3 indicates that the smallest hardened-layer thickness required is obtained under the condition that the resulting stress shall be higher than the breaking stress, according to Hook’s model, and no cracks occur at the workpiece surface after grinding (zmin = 2.873 mm) (Ref 67).

Using the proposed analytical model, the researchers (Ref 67) evaluated the magnitude of stresses in steel after surface hardening and grinding. They found a mutual dependence of the hardened-layer profile and the resistance against material cracking during grinding. Cracks will appear at the surface of those bearings in which short overheating of steel occurred and the hardness in a narrow zone decreased strongly during grinding. In such cases, through hardenability of a material reaches a depth of 150 mm, where tempered martensite will form. The thickness of the surface-hardened layer with all the cracked bearings was approximately 2.0 mm, which indicated that an appropriate model was chosen.

Fatigue Strength of Materials

Fig. 90

Hardness characteristic of neck cross section without cracks as a function of depth, z. Source: Ref 67

A heat treatable AISI 4140 steel was used for manufacturing crankshafts. This steel is very appropriate for statically and dynamically loaded parts of car engines and machines because of its high hardness achieved after hardening (57 HRC). The steel is characterized by good hardenability and is thus suitable for manufacturing machine parts with large cross sections in which a very high strength can be obtained after refinement. After tempering, the steel does not show a tendency to brittleness, and therefore, no special heat treatment procedures are required. This steel is also suitable for surface hardening (flame surface hardening, induction surface hardening) and displays a very good resistance to wear. However, special attention must be paid to the part design phase, and great care should be given to the design of radius and transition areas to prevent notch effects under dynamic loads. The steel is adapted for use in a wide range of Table 3 Minimum case depth required for no cracks at workpiece surface

Fig. 91

Thermal cracks in surface-hardened neck section. Source: Ref 67

Excentricity of hardened layer (e), mm

Depth (z), mm

Temperature prior to quenching used in Hook’s model (TK), °C

0.594 0.709 0.877 1.146

3.156 3.0412 2.873 2.6045

343 336 325 307

Source: Ref 67

N/mm2

Resultant stress after hardening and grinding (sR(T0)), N/mm2

1123 1144 1177 1232

1306 1327 1360 1414

Yield stress of martensite , RK 0:2, (T K )


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temperatures, and it preserves high toughness even at low temperatures (Ref 44, 68, 69). Impact tests provide useful quantitative comparative data with relatively simple test specimens and equipment. However, these tests do not provide property data for design purposes for material sections containing cracks or defects. Data of this type are obtained from fracture mechanics, in which theoretical and/or experimental ways determine the fracture of structural materials containing pre-existing cracks and defects. The number of cycles to failure (N ) obtained by the S-N curve under load or a stress-control condition is related to the total life of the part up to failure. Fatigue cycles throughout the life of the machine part consist of crack initiation and its propagation. Crack propagation refers to stable crack growth up to the crack instability. The S-N curve approach does not separate the crack initiation phase from propagation. Industrial experts assume that the crack has already initiated in the machine part, and only the total number of cycles associated with the propagation are of interest. Existing or initiated cracks assumed in the machine part are the result of dynamic load cycles induced during manufacturing, prior to its use. The size of the preexisting crack can be assumed based on the capability of suitable inspection. The expert may assume an initial surface crack after testing and makes a decision about the application of the part. Using the available initial defects in the material, the total life of the part can be assessed by an appropriate fracture mechanics method. Basic factors affecting the shape of the S-N curve are:

Materials selection and heat treating or cold working conditions Various types of loading on the specimen, such as tension, compression, torsion, or a combination Loading conditions described by medium stress, amplitude stress, and frequency Influences on the environment carried by temperature, corrosion, and other factors

Some major factors that affect the strength of a metal include:

Stress concentration: Fatigue strength is greatly reduced by the presence of stress raisers, such as notches, holes, keyways, or sharp changes in cross sections.

Surface roughness: In general, the smoother the surface finish on the metal sample, the higher the fatigue strength. Surface condition: Since most fatigue failures originate at the metal surface, any major change in the surface condition will affect the fatigue strength of the metal. Environment: If a corrosive environment is present during the cyclic stress of a metal, the chemical attack greatly accelerates the rate at which fatigue cracks propagate.

During machining processes, various defects occur on the surface, such as small scratches and grooves, and are introduced into the workpiece surface. Typical failures are moving machine parts, such as shafts, connecting rods, and gears. It is estimated that failures of machine parts in machines contribute approximately 80% of fatigue failures. These surface detects can limit the fatigue life. Improving the surface finish by polishing will increase fatigue life significantly. One of the most effective methods of increasing fatigue life is the existence of residual compressive stresses in a thin surface layer. Thus, applied surface tensile stress will be partially reduced in magnitude by the residual compressive stress. The net effect is the probability of crack formation and a consequent reduction in fatigue failure. According to the AISI standard, the heat treatable structural steel 4140 contains between 0.38 and 0.45% C, 0.90 and 1.2% Cr, and 0.15 and 0.30% Mo. It has very high hardenability, contributing to high strength values in products with high mass. Molybdenum yields a desirable fine microstructure after hot working as well as heat treatment, contributing to a good strengthto-toughness ratio. Due to its fine-grained microstructure, it also reaches a relatively high toughness in the heat treated condition. The strength of the steel as well as its surface hardness and wear resistance may be increased by heat treatment and thermochemical treatment. Mechanical properties of steel having a diameter of up to 40 mm and between 40 to 100 mm are given in Table 4. Tensile strength of the steel varies between 880 and 1080 N/mm2, and a minimum toughness value, r3, equals approximately 41 J. The steel is very sensitive to notch and transition on machine parts subjected to fatigue loading. Fatigue strength of the material is lowest under torsional load, sT, and varies for the diameters

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mentioned, that is, 16 to 40 mm, so that sT = 285 N/mm2, and for diameters from 40 to 100 mm, sT = 255 N/mm2. Fatigue strength under torsional load is three or four times lower than the tensile strength of steel under static load, according to the data in the table. In professional literature, data on fatigue strength of materials are usually presented for a prescribed specimen shape and size that has been adjusted to the testing device. Specimens for fatigue strength are usually of cylindrical shape, with a smaller diameter in the middle part and a rounded transition into the larger-diameter part. The latter is then usually clamped for testing. Modes of loading the specimens vary but are usually either torsion, bending, or tension/ compression. The highest fatigue strength is displayed by a specimen subjected to bending loads (Ref 27, 36). For other modes of loading, the relation with bending fatigue strength, swb, is expressed empirically, that is, fatigue strength in torsion, tw, is 0.58 swb, or fatigue strength in tension/compression, swz, is 0.70 swb. Table 4 Mechanical properties of heat treated structural steel 4140 Diameter (D), mm

16–540 40–5100 Source: Ref 15

Fig. 92

Tensile strength (Rm), N/mm2

Yield point (Rp0.2), N/mm2

980–1180 880–1080

769 635

Extension Toughness (r3), J (A5), %

11 12

41 41

Figure 92 presents curves for different modes of dynamically loaded specimens made from various steels that had been heat treated in assorted ways (Ref 15). From among six curves, four represent specimens made from heat treatable steel and two for cementation steel specimens. The heat treatable steel is Cr-Mo-Ni steel with 0.37% Ni, and the specimens were heat treated in two different ways:

Surface hardening applied to specimens with a smooth shape (curve 1) and specimens with a slot (curve 2) Quenching or tempering applied to smooth specimens (curve 4) and specimens with a slot (curve 6)

The cementation steel is a chromium-nickel steel with 0.15% C, where the specimens were smoothly shaped (curve 3) and slotted (curve 5). Axle shafts used in cars, trucks, and farm vehicles are, with few exceptions, surface hardened by induction. Although a portion of the hardened surface is used as a bearing in some axles, the primary purpose of induction hardening is to put the surface under a state of compressive residual stress (Ref 45). By this means, the bending and torsional fatigue life of an axle may be increased by as much as 200% over that for parts conventionally heat treated (Fig. 93) (Ref 45).

Graphs of the fatigue strength of surface-hardened and carburized specimens. Source: Ref 15

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Induction-hardened axles consist of a hard surface, high-strength core, and tough outer case with good torsional strength and a tough, ductile core. Many axles also have a region in which the case depth is kept very shallow, so that the part can be readily straightened following heat treatment. In addition to substantially improving strength, induction hardening is also very costeffective. This is because most shafts are made of inexpensive, unalloyed medium-carbon steel that is surface hardened to case depths of 2.5 to 8 mm, depending on the cross-sectional size. As with crankshafts, typical hardness (after tempering) is approximately 50 HRC. Such hard, deep cases improve yield strength considerably as well (Ref 45). Modern transmission shafts, particularly those for cars with automatic transmissions, are required to have excellent bending and torsional strength, as well as surface hardness for wear resistance. Under well-controlled conditions, induction-hardening processes are able to satisfy these needs, as shown by the data in Fig. 94, which compares the fatigue resistance of through-hardened axles (Ref 45). The inductionhardening methods employed are varied and

Fig. 93

include both single-shot and scanning techniques. The hardness achieved on the surface of the hardened specimens was 56 to 59 HRC and on the cemented specimens, 58 to 59 HRC. To test the effects of the slot on fatigue strength, a slot of equal size and shape was made on all the specimens, whether they were quenched or tempered, hardened or cemented. It was made in the middle of the cylindrical specimen to an equal depth of 0.4 mm. The depth on the surface-hardened and cementation specimens was 1.5 mm. The results of testing showed that there are significant differences in terms of heat treatment methods and that the highest fatigue strength was found in surface-hardened specimens. A comparison of the fatigue testing results showed that:

In surface-hardened specimens with a smooth cylindrical shape and with a slot, the difference in the achieved fatigue strength is minimal. This can be attributed to a very desirable distribution and size of compressive residual stresses throughout the hardened

Bending fatigue response of furnace-hardened and induction-hardened medium-carbon steel tractor axles. Shaft diameter: 70 mm. Fillet radius: 1.6 mm. Source: Ref 45

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layer. Since the depth of the slot reaches only one-quarter of the hardened surface layer, the size of compressive residual stresses at the slot is still very high, so that the weakening due to the slot and stress concentration along the slot does not cause any essential drop in fatigue strength. In quenched and tempered specimens with no earlier prestressing of the surface layer, the fatigue strength was considerably lower than in surface-hardened specimens. The fatigue strength of quenched and tempered specimens with a slot was also remarkably lower. The results show that the difference in fatigue strength in slotted induction surfacehardened specimens and slotted quenched and tempered specimens is 5 to 1. Smooth cemented specimens displayed 25% lower fatigue strength than the surfacehardened specimens of the same shape, whereas the cemented specimens with a slot displayed 50% lower fatigue strength than the same surface-hardened specimens. The question arises about what the fatigue strength is in those specimens where the slot reaches deeper than the hardened layer. A regular problem in these cases is crack occurrence and propagation of cracks starting from the slot. Due to the shape of the specimen and the slot, stresses start concentrating at these places, depending on the type and size of external loads. It should not be forgotten that there are no compressive residual stresses along the slot, and the size of tensile stresses along the slot plays a decisive role in crack occurrence.

Figure 95 shows bending stress in a tooth root subjected to dynamic load versus the number of

Fig. 94

oscillations (Ref 36). Figure 95(a) shows the bending stress for induction surface hardening of adjacent flanks of two teeth with a coil reaching into the tooth gap. In the process, the tooth flank as well as the tooth root are hardened (Ref 36). This kind of heat treatment of gears from steels for induction surface hardening provides a fatigue bending strength in the range of 320 to 490 N/mm2. Figure 95(b) shows the same relationship when the induction coil encircles an individual gear tooth. In this process, the tooth flank is hardened, and the microstructure and hardness in the tooth root are preserved (Ref 36). A result of this method of hardening is that the fatigue strength is drastically lowered to values ranging from 200 to 300 N/mm2 for the entire range of steels suitable for induction surface hardening. This is a considerable drop in fatigue strength for the material in the tooth root (Ref 45, 70).

Stress Profiles in Machine Parts in the Loaded State Heat treatment engineers must be very careful in choosing the conditions of induction surface hardening in order to benefit from the distribution of residual stresses achieved in dynamically loaded parts. In industrial practice, induction surface hardening should satisfy the requirement of fatigue resistance of machine components. The main reason for this worsening of the properties of the machine part is attributed to tensile residual stresses in the hardened layer and undesirable hardness distribution in the transition zone from the hardened into the unhardened part of the subsurface. These effects

Comparison of fatigue life of induction surface-hardened transmission shafts with that of through-hardened and carburized shafts. Arrow in lower bar (induction-hardened shafts) indicates that one shaft had not failed after testing for the maximum number of cycles shown. Source: Ref 45

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are quite natural, and in the first phase are a result of very rapid local heating of the thin surface layer, while in the second phase this is accompanied by forced quenching, which ensures a critical cooling rate and the occurrence of a martensitic microstructure. Both phases in induction surface hardening can increase the risk

Fig. 95

of fatigue, especially if the latter is assessed only from the point of view of surface hardness. To successfully estimate the quality of the hardened layer, one must select the optimal synergetic effects between the input electric energy and the interdependence between the induction coil and the workpiece surface, connected with the

Bending fatigue strength of gear teeth at (a) tooth gap hardening and (b) flank hardening for various steels. Broken lines denote confidence limit according to DIN 3990. Source: Ref 36

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occurrence of eddy currents in the workpiece surface layer that leads to heating. Due to complex synergetic effects in induction heating or hardening, it is necessary to carefully study each influence on the properties of the hardened surface layer (Ref 27, 45). Fatigue strength in machine components that have been induction surface hardened is increased if the total sum of the load tensions and residual stresses in the surface layer is of a compressive nature. To ensure the highest fatigue strength of a component, it is necessary to provide the following:

be treated from the point of view of heating, overheating, and cooling/quenching as well as the internal stresses created at a certain point during the treatment. During heat treatment, internal stresses are created by the temperature differences and phase transformations between the core and the surface, which are a result of the volume differences between the core and the surface. The created volume differences between the core and the surface then give rise to internal stresses. During the process of heating and cooling, internal stresses may produce the following effects:

In dynamically loaded components, the surface is prone to fatigue occurrence, so the surface must have the highest compressive stresses. If the total sum of stresses, that is, load tension plus residual stresses on the surface, is always of a compressive nature, then there is no chance for the occurrence of cracks and crack growth. To ensure good behavior of the surface and the hardened surface layer in the loaded condition, it is necessary to induce a suitable prestressing in the surface layer. This can be achieved by a carefully selected heat treatment method that would create the highest compressive residual stresses on the surface and a desirable profile of the latter in the hardened subsurface layer. Induction surface hardening offers opportunities to ensure a considerable amount of compressive stresses in the machine component surface and to ensure restrained transformation of compressive surface stresses into tensile residual stresses in the subsurface layer.

The endurance of machine components subjected to bending and torsion loads can be successfully increased by ensuring sufficiently high compressive residual stresses. A manufacturing goal is to create a sufficient amount of compressive residual stresses with a favorable distribution, since this is the only way to increase the reliability of components in operation. An early failure of a component in operation may cause catastrophic damage on a machine and thus a loss in profit. A decisive role in the occurrence of residual stresses is played by the synergetic effects between the heat treatment method, the type of material, and the shape of the workpiece. For this reason, heat treatment must

When internal stresses are lower than that of the yield point, higher residual stresses are induced by heat treatment in the workpiece, but these would not cause distortions, cracks, or failure. During a certain moment in heat treatment, internal stresses exceed the yield point, which leads to distortions and lower residual stresses in the workpiece. During very detrimental conditions in heat treatment, internal stresses are higher than the tensile strength of the material, causing the workpiece to crack and creating larger distortions and high residual stresses.

Numerous changes that take place in the hardened surface layer of the workpiece are always a result of the heating and quenching conditions. Therefore, it is necessary to study the events taking place in the workpiece directly after the hardening temperature is reached. Three zones are distinguishable in a workpiece heated to the hardening temperature (Fig. 96a): the first zone, where the outer layer is heated to the hardening temperature; the second zone, which is heated below the hardening temperature between the temperatures TA1 and TA3 for rapid heating; and the third zone, where the temperature is lower than TA1 (Ref 15). Heating to the hardening temperature at a certain depth is followed by quenching. Quenching results in the occurrence of compressive residual stresses (Fig. 96b), when the familiar transformations in the hardened layer take place. The second layer does not suffer the same distortions as the surface layer, although the heating there has been sufficient enough to improve the properties of the material. In the second layer, hardening is incomplete, which, in comparison with the first layer, results in lower hardness and strength of the material.

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When a machine component that has been surface induction hardened in this way is subjected to an external load (Fig. 96c), additional tensile residual stresses can be noted in the first and second layers (Ref 15).

Fig. 96

Since fatigue of a machine component is a very delicate problem connected with the total sum of tensile stresses in the second zone, there is a danger that the effects of fatigue are transferred to the surface, due to the size, shape, and

Stress profile in a round bar in the loaded state, where residual stresses after induction surface hardening and loading stresses add up. Source: Ref 15

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location of the hardened trace. When a loaded layer is exposed to tensile external stress up to the surface, the surface becomes very sensitive to the occurrence of cracks (Fig. 96d) (Ref 15). This propensity to cracking is further increased if defects are present in the surface of the workpiece material. A critical state in the second zone of the hardened surface layer occurs on locally hardened workpieces, where the surface was overheated or nonuniformly heated. The conditions are quite the opposite in surface heating of workpieces, such as in nitriding and cementation or intensive but shallow quenching of steel (Fig. 97) (Ref 15). In these surface heat treatment procedures, heating is carried out throughout the whole machine component, followed by slow cooling, for example, nitriding or quenching such as in cementation. Residual stresses in nitrided and cemented surface layers are compressive, while tensile residual stresses occur in the central part with the refined pearlite microstructure. The distribution of residual stresses in heat treatment procedures where only the surface of

Fig. 97

Stress profile in a round bar in the loaded state, where residual stresses after carburizing or nitriding and loading stresses add up. Source: Ref 15

the workpiece is heated (induction hardening, flame hardening) differs greatly from the procedures where heating is performed throughout the entire volume (nitriding, cementation). In nitriding and cementation, the aforementioned second layer in the subsurface does not appear at all, because the direction of the heat flow is opposite to the direction of the heat flow in induction and flame hardening. The resultant operating tensile stresses on the surface or in the surface layer can thus be considerably smaller. Due to the hardness of the surface, induction and flame hardening lowers the fatigue strength of machine components. Therefore, care should be taken to diminish all detrimental effects in the surface layer. A typical example of induction surface hardening is surface hardening of gears that are heated with a low heating rate and relatively low current frequency. The outer hardened zone includes almost the entire height of the gear teeth, whereas the second zone is in the tooth root area. A gear heat treated in this way will meet the wear resistance requirements expected of the gear tooth, while the strength of the other part of the tooth is of minor importance. The fatigue strength will be relatively low due to high tensile residual stresses in the tooth root, that is, in the second zone where the operating or load tensions and the tensile residual stresses are summed up. The energy input in heating a gear tooth or the whole gear was such that the second zone has not appeared. A similar heat treatment can be applied to the spline inside the gear. A gear heat treated in this way is more resistant to wear and corrosion and should have high resistance to fatigue in bending because of a smaller thickness of the layer in the second zone. Residual Stresses in Carburized Machine Parts. Carburizing is a process in which an austenitized ferrous alloy is brought into contact with an environment of sufficient carbon potential to cause absorption of carbon at the surface and, by diffusion, to create a carbonconcentration gradient in the thin surface layer. As this definition clearly indicates, two factors may control carburizing. Either the carbonabsorption reaction at the surface or the diffusion of carbon in the steel will determine the rate of carburizing. Carburizing is done at elevated temperature, generally in the range of 850 to 950 C, although occasionally at temperatures as low as 790 C and as high as 1095 C.

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With the hardened steel, that is, hardened machine part, the following was studied:

Effective case depth Depth at which martensitic microstructure is present Hardness variation in the transition zone between martensite and matrix Microstructural composition in the transition zone with bainite-pearlite-ferrite

The characteristics obtained are essentially affected by:

The grade, chemical composition, and microstructure of steel in the soft state The size, that is, mass, of the machine part

Considering that certain properties described by the previously mentioned characteristics are specified for the machine part, adequate conditions for quenching from the austenitizing temperature should be specified too. The quenching conditions are described with reference to the mass of the machine part. Thus, in steel quenching, the following are studied:

Volume, that is, dimensional changes due to the differences in microstructure between the initial material condition and the hardened condition of the same material Distortion of the machine part due to incorrect shaping of the machine part and/or an improper quenching procedure Residual stresses due to volume changes occurring between the hardened zone and unhardened zone of the machine part

Residual stresses may result from a deformation of the machine part during the quenching process. Internal forces due to temperature stresses during quenching may exceed the yield stress of the material, which results in plastic deformation during the quenching process and residual stresses after cooling. The magnitude of residual stresses is related to the yield stress of the material at the temperature at which the deformation occurred. Thus, in the case of through hardening in which a complete and homogeneous austenitic microstructure turns into a martensitic one, a 4% volume change of the machine part or a 1.3% linear increase in the machine-part size is obtained: DV 100=4% V

In incomplete hardening, the volume changes at the surface are greater than in the core, which produces transformation strains. The residualstress variation thus depends on:

Quenching/cooling conditions Temperature difference between the surface and the core during the quenching process Temperature interval between the beginning and end of martensitic transformation and the cooling rate in this zone

An image of the magnitude of internal stresses during the quenching process can be provided by a test of quenching cylindrical specimens from a temperature lower than the transformation temperature. During the quenching process, temperatures are measured at the specimens, that is, at their surface and in their core. Maximum thermal stresses can thus be calculated (Ref 1, 20, 71, 72). The maximum thermal stresses found in the cylindrical specimens with the smallest diameter, 25 mm, and further specimen diameters in a geometrical ratio with a factor of 2 are given in Table 5. As an approximate orientation to the cooling rates obtained, a calculation was made for a temperature of 500 C. Comparative data on maximum thermal stresses refer to air cooling and oil quenching. With air cooling, the stresses obtained in the specimen with the smallest diameter, 25 mm, equal only 7 MPa, whereas for the specimen with the largest diameter, 800 mm, the stresses are as much as 200 MPa. The maximum thermal stresses, however, are considerably higher with oil quenching; they range between 230 MPa with the 25 mm diameter specimen and 620 MPa with the 800 mm diameter specimen. The latter represents an extremely high internal thermal stress that may result in material plastification. The data exclude stresses due to phase changes and stresses due to inhomogeneity, which additionally increase their value. Results of the measured residual stresses with surface-hardened steels are given in Table 6. The stresses were measured by the x-ray diffraction technique just below the surface, that is, 0.05 mm. The results refer to steel grades 832M13, 805A20, 805A17, 897M39, 905M39, and cold rolled steel after induction hardening, with the longitudinal residual stresses being measured. The values of longitudinal stresses after carburizing to 1.0 to 1.5 mm case with 0.8% surface carbon and direct quenching without tempering range between 190 and

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400 MPa, which depends on the steel grade selected. With steel 832M13, testing of specimens after quenching and by undercooling to 80 and 90 C with tempering and without tempering was performed. Tempering contributes to obtaining tempered martensite and possible transformation of retained austenite, which produces a reduction, that is, a compensation, of the existing residual-stress profile. Thus, the lowest residual stress is obtained in steel quenching by undercooling and subsequent tempering. In nitriding of steel 897M39 to a depth of 0.5 mm, however, a residual stress ranging between 400 and 600 MPa is obtained. In steel 905M39, however, a considerably higher residual stress, ranging between 800 and Table 5 Cooling rate thermal stresses in simple rounds with no transformation Air cooled Diameter of round (D), mm

25 50 100 200 300 400 800

Oil quenched

Cooling rate at 500 °C (v500), °C/s

Maximum thermal TH stress smax , MPa

Cooling rate at 500 °C (v500), °C/s

Maximum thermal TH stress smax , MPa

0.662 0.312 0.146 0.070 0.0445 0.0326 0.0158

7 15 28 54 73 100 200

20.0 6.12 1.88 0.59 0.29 ... ...

230 290 370 450 510 (540) (620)

Source: Ref 1, 71

Table 6 Residual stresses measured in surface heat treated steels Steel

832M13

805A20 805A20 805A17 805A17

897M39 905M39 Cold rolled steel

Heat treatment

Carburized at 970 C to 1.0 mm case with 0.8% surface carbon Direct quenched, 80 C subzero treatment, no temper Direct quenched, 90 C subzero treatment, tempered Carburized and quenched Carburized to 1.1–1.5 mm case at 920 C, direct oil quench, no temper Carburized to 1.1–1.5 mm case at 920 C, direct oil quench, tempered at 150 C Nitrided to case depth of approximately 0.5 mm Induction hardened, untempered Induction hardened, tempered at 200 C Induction hardened, tempered at 300 C Induction hardened, tempered at 400 C

(a) Immediately subsurface, that is, 0.05 mm. Source: Ref 71

Residual stress (longitudinal), MPa

1000 MPa, is obtained in nitriding to the same depth. The highest residual stresses were measured in cold-rolled steel subjected to additional induction surface hardening without tempering. A maximum value obtained was 1000 MPa. With additional tempering at 200 C, martensite will become tempered and the residual stresses reduced to 650 MPa. By additional tempering at an even higher temperature, at which martensite disintegrates, the maximum measured stresses decrease. For example, at a tempering temperature of 400 C, they decrease to 170 MPa or a factor of 6 lower than those obtained immediately after steel quenching. Figure 98 shows the results of carbonitriding SAE 1118 steel with reference to the contents of carbon and nitrogen at the surface and the gradient of the two in the subsurface (Ref 71). The austenite content, that is, the austenite gradient of the two in the subsurface, and the way a complementary residual-stress variation proceeds in the subsurface are very important. With a gradual reduction of austenite content to a depth of 0.5 mm, the residual stresses decrease as well, and the maximum pressure achieved is 200 MPa. To a depth of 45 mm, the stresses gradually increase to tensile values, and at a depth of 1.5 mm, they are as much as +100 MPa.

Input and Output Control of Steel for Induction Surface Hardening of Gears Production automation necessitates an everincreasing demand for greater uniformity of the

280 340 200 240–340(a) 190–230 400 150–200

400–600 800–1000 1000 650 350 170

Fig. 98

Residual stress, carbon, nitrogen, and retained austenite through a carbonitrided case on SAE 1118 steel. Source: Ref 71

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properties of heat treated steel parts (Ref 73). Therefore, a general market demand is for a certain prescribed quality that is usually the result of various properties (Ref 41, 74–82). One of the most important properties is hardenability, which should be a uniform and constant property of steel regardless of the treatment conditions in the ironworks. In the experimental part of the study, the starting point was the analysis of steel hardenability according to the Jominy test (Ref 40, 74–76). On this basis, the conditions for induction hardening were prescribed for a worm gear tooth. The success of the induction hardening was verified by the microhardness measurement along the tooth profile centerline (Ref 83, 84). Figure 99 shows Jominy curves for the hardness profiles of the specimen front of AISI 1050, 4150, and 4340 steels (Ref 40, 41). It generally applies that Jominy curves for common testspecimen heating differ strongly from those for the specimens induction heated and quenched in the same way. Hardness differences in the Jominy test specimen are minimal at the surface and increase through the depth. It is important that the hardness profiles of the furnace-heated test specimens or short-term induction-heated specimens differ less if carbon steels have a smaller carbon content (AISI 1050) and differ more strongly with alloyed steel, as was the case. Figure 100 shows Jominy curves for furnace heating and induction heating. Both figures indicate that the through-height hardness profile of a Jominy specimen depends on the maximum temperature obtained at the surface (Fig. 100b) and on heating time (Fig. 100a), the power density being given (Ref 40, 41). Induction heating was performed under the conditions given and with different heating times. Furnace heating of Jominy specimens at 870 C provides the highest hardness values for the hardened front. In induction heating at the same temperature, 870 C, a quite different hardness profile of the Jominy specimen hardened front is obtained (Fig. 100a). With an increasing surface austenitizing temperature, the through-height hardness profile of the specimen increases so that at a temperature of 1040 C, there is negligible difference in the hardness profile of the Jominy specimen, regardless whether it was furnace heated or induction heated. The prescription for the verification of the effects of induction hardening was carried out for the upper- and lower-limit microhardness

from the tooth top and to a depth of 9/6 of the tooth height. The experiments on heat treatable carbon steel C35 with a different history have

Fig. 99

Jominy curves for end-quenched bars of (a) AISI 1050, (b) 4150, and (c) 4340 steels, austenitized conventionally and by short-time induction heating. Source: Ref 40, 41

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shown that with the existing heat treatment conditions, it is not possible to achieve the microhardness profile required by the user (Ref 79). First, it was necessary to find a Slovenian substitute for the French steel CC35. This was done with the help of a cooperative industry. Steel CC35 is a heat treatable steel that has been mechanically machined and induction hardened to ensure a uniform quality. Steel C35, chosen as a substitute, is classified among the high-grade, unalloyed, heat treatable carbon steels with a maximum phosphorus and sulfur content lower than 0.035%. It is intended for parts smaller than 100 mm, subjected to lower loads, and with high requirements on homogeneity of material after heat treatment and mechanical machining.

Effect of (a) time at an 870 C austenitizing temperature and (b) maximum surface temperature on the Jominy curves for induction-hardened AISI 4150 steel. The curve for conventional furnace-heated 4150 is also shown in (b). Source: Ref 40, 41

Fig. 100

For the selected heat treatable carbon steel C35, a suitable acceptance control procedure had to be prescribed, with special emphasis on hardenability. Based on the corresponding criteria, the acceptance conditions should enable the classification of steels into quality classes in accordance with the different heat treatment prescriptions. The chemical composition of the steel is presented according to International Organization for Standardization (ISO) standards and that of the French according to the Association Francaise de Normalisation (AFNOR). The chemical composition is given by quoting the upper and lower content limit for each particular part. The limits of the Slovenian steel are much more constrained than that of the French. In the French steel, there is a higher content of carbon, manganese, and traces of chromium, which leads to somewhat better hardenability compared to the Slovenian steel C35. Figure 101 presents the hardness curves of the Jominy specimen versus the distance from the face (Ref 83). The starting point for the analysis was the upper and lower confidence limit for Slovenian steel C35 as prescribed by the producer, Ravne Ironworks. The given confidence limits are presented by the hatched area on the extreme right of the figure. Since testing enables a successful selection of steel as well as optimal heat treatment conditions, it was decided to test the upper and lower confidence limits of a small-sized, hardened worm gear (Fig. 102a) (Ref 83). This gear must have hard, wear-resistant surfaces and an increased strength due to loading conditions or expected load-carrying capacity of the teeth. Both of these properties are achieved by induction hardening. The success of the heat treatment was verified by Vickers microhardness measurements at a loading of 0.3 daN acting along the centerline of the tooth profile. The verification of the effects of surface hardening and simultaneous heat treatment of the teeth was carried out in ten measurements: the first on the tooth top and then each subsequent measurement at a distance of 1/6 h or 0.367 mm. Thus, in ten measurements the microhardness was measured down to the depth of 9/6 of the tooth height, as shown in Fig. 102(b) (Ref 83). When ordering the gear, the customer specified the prescribed allowable hardness for forming the upper and lower confidence limits along the height of the gear tooth.

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The microhardness limit values were:

On the tooth top at h = 0, the microhardness is 680 to 800 HV0.3. In the depth of 4/6 h or 1.45 mm, the microhardness is 520 HV0.3 maximum. In the depth of 8/6 h or 2.90 mm, the microhardness is 520 HV0.3 maximum. In the depth of 9/6 h or 3.30 mm, the microhardness is 220 to 275 HV0.3.

On the basis of hardenability testing, the following conclusions can be drawn:

Due to a higher content of carbon, manganese, and some traces of chromium, the

French steel CC35 displays a very favorable hardenability according to the Jominy tests. Even to a depth of 7 mm, the hardenability curve shows 40 HRC, then hardness begins to decrease rapidly and, at a depth of 10 mm, is only 22 HRC. The Slovenian steel C35 produced by the Ravne Ironworks has a slightly lower content of carbon and a considerably lower content of manganese. For example, for steel of charge “B”, the carbon content is lower by 0.06% and the manganese content by 0.35%, according to the French steel. On the basis of the data in the literature, it was determined that this considerably lower manganese

Fig. 101

Hardenability of the analyzed steels and determination of the upper and lower confidence limit. Source: Ref 83

Fig. 102

Worm characteristics after induction surface hardening. Source: Ref 83

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content reduces the effect of hardenability by a factor of 1.5. The relative effect of silicon on hardenability is practically negligible, although the content may vary within 0.1 to 0.4%. The results of the hardenability testing of the discussed steels are presented in Fig. 101 (Ref 83). Based on the mean values and their deviations, it was determined that the curve of hardenability or hardness shows solid progress into the depth in the French steel CC35 and steel C35 of charge “C” and “D”. On the other side, the hardenability of Slovenian steel C35 of charge “C” and “D” is much lower, so that at a depth of 4 to 5 mm, the lower confidence limit is exceeded. The input control of steel should be focused mainly on the lower confidence limit of hardenability, because the progress of hardness has a decisive influence on the strength properties of the gear teeth. It was determined that the upper confidence limit corresponds to the maximum hardness achieved, and that it cannot be exceeded in any way. It was therefore decided that the place of the corrected lower confidence limit would be defined using Jominy hardenability tests, determining the actual progress of hardness. As a criterion for the correction of the lower confidence limit, the microhardness along the tooth height (9/6 h), limited by the maximum and minimum microhardness, was used. This helped to establish the primary criteria conditioning the microhardness distribution on the

Fig. 103

product, which was also limited by the upper and lower confidence limit. The answer to which charges of the steel are more suitable for use can be obtained by associating the corresponding microhardness confidence limits on the product and also on the Jominy specimen. The desired product characteristics were defined on the basis of the following conditions of induction hardening:

Generator power—22.5, 24.0 and 25.5 kW Heating time—3.9, 4.4, and 4.9 s Time/pause between the end of heating and the beginning of quenching—0.1 s Quenching time—4.0 s

The results of induction hardening on the gear are presented in Fig. 103 for the French steel CC35 (Ref 83), and in Fig. 104 for the Slovenian steel C35 at the given power values of a highfrequency generator and heating times (Ref 83). The diagrams show the upper and the lower confidence limits for microhardness along the tooth height. The results of microhardness distribution along the tooth height must fall within the mentioned limits. However, it can be seen that it is possible to reach only a lower microhardness on the tooth that results from a lower hardenability of the steel or from unsuitable heat treatment conditions. The lower confidence limit is defined only by the microhardness of the basic material at a depth of 4/6 to 9/6 h and represents only a theoretical limit. The data on

Hardness distribution along the tooth symmetry line after heat treatment of the French steel CC35. Source: Ref 83

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the steel properties and heat treatment conditions provide various actual measured microhardness distributions, where the microhardness of the basic material is reached in the depth from h to 9/6 h in the French steel CC35 and from 11/12 h to h in the Slovenian steel C35. From all these data, it can be stated that it is possible, with adequate heating and quenching conditions and adequate hardenability, to reach the desired microhardness distribution in the tooth. From Fig. 103, it can be noted that in the case of the French steel CC35 only three out of nine heat treatment conditions fall out of the confidence limit at the minimum power P = 22.5 kW (Ref 83). Quite the contrary can be found from Fig. 104 presenting Slovenian steel C35. Here, a desirable distribution of microhardness along the tooth height was achieved only in the conditions of maximum power (Ref 83). This means that for an equal efficiency of inductionhardened gears from the French steel CC35, a significantly shorter heating time is necessary than with steel C35. This time is estimated to be for even a second shorter and represents a 25% shorter heat treatment cycle, contributing to lower costs of manufacturing. Induction surface hardening of machine components and especially gears is a very complex process involving a whole range of possible heat treatment methods, which are all reflected in either good or bad serviceability of machine components. The heat treatment engineer must be aware of the different effects of particular design shapes of induction coils, be

Fig. 104

familiar with electromagnetic phenomena and eddy currents, and have some experience in the right choice of energy inputs necessary for heating. The energy needed for heating can be provided by changing the generator power as well as the frequency of the current. In progressive hardening, to achieve a suitable energy input, the workpiece feed rate or the rate at which the coil is moved must be adjusted, whereas in single-shot hardening, suitable energy input is achieved by adjusting the heating time with a high-frequency current. An essential advantage of induction surface hardening is that it is possible to achieve a sufficient repeatability of the hardened layer thickness on the workpiece as well as a desirable or even prescribed hardened-layer profile, ensuring sufficient hardness and favorable distribution of residual stresses in the hardened layer. A variety of steels and a whole range of induction-hardening methods provide the possibilities for very accurate planning of the size and distribution of residual stresses. This is of growing importance, since manufacturers are frequently required to produce machine components that, among other surface properties, must possess quite specific residual-stress distribution along the depth of the hardened layer. It has become a proven fact that high compressive stresses ensure high fatigue strength of machine components and reduce the danger of the occurrence and growth of cracks on the surface of components. As far as induction surface hardening is concerned, it is also quite important to choose the right quenching

Hardness distribution along the tooth symmetry line after heat treatment of the Slovenian steel C35. Source: Ref 83

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medium and method of quenching. For this reason, engineers must direct their attention not only to the method of heating and possible overheating of the surface layer but also to the methods of quenching and the right choice of the medium for quenching. With increasing surface heating power in the final phase, an increase in microhardness along the tooth symmetry line is also reached, especially in the depth from 0.5 to 2.5 mm, whereas by making the heating time longer, a more pronounced effect directed into the depth can be achieved. In this way, an adequate heat treatment can ensure the required surface hardness (52 HRC) as well as an increased strength of the gear in the tooth-root part, which is of extreme importance for certain operating conditions. Tests have shown that it is necessary to move the lower hardenability limit at least on the level indicated by the scattered microhardness values of the French steel CC35. The new lower confidence limit is named the corrected limit, ensuring the prescribed microhardness distribution in the subsurface at a power P = 24.0 kW and P = 25.5 kW. On the basis of the tests, it can be concluded that the criterion of hardenability can be very successfully applied in the input control of steels. Hardness decreases with the distance from the face of the Jominy specimen and must fall within the confidence limits. In some cases, that is, with Slovenian steel C35 of charge “A” and “B”, when the hardenability exceeds the lower confidence limit, an adequate selection of induction-hardening conditions (power in kilowatts and heating time in seconds) can ensure the desired microhardness along the tooth symmetry line. The results confirm that heat treatment conditions can be successfully determined by relatively simple experiments that also make the procedure more economical.

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69. V. Nemkov, Role of Computer Simulation in Induction Heating Technique, Proc. of the Int. Induction Heating Seminar (Padua), 1998, p 301–308 70. P.K. Braisch, The Influence of Tempering and Surface Conditions on the Fatigue Behaviour of Surface Induction Hardened Parts, Mater. Sci. Forum, Vol 102–104, 1992, p 319–334 71. H.C. Child, Residual Stress in Heat-Treated Components, Heat Treat. Met., Vol 4, 1981, p 89–94 72. H.C. Child, The Heat Treatment of Tools and Dies—A Review of Present Status and Future Trends, Tools and Dies for Industry Book, The Metals Society, p 185, 291–321 73. Electromagnetic Induction and Electric Conduction in Industry, Chapter 9, Heat Treatments by Induction, Centre Francais de I’Electricite, 1997 74. R. Jo¨nsson, Toleranzen bei der Wa¨rmebehandlung von Sta¨hlen in Ha¨rterein, TZ Met. bearb., Vol 80 (No. 5), 1986 (in German) 75. K.E. Thelning, Chapter 4: Hardenability, Steel and Its Heat Treatment, Bofors Handbook, Butterworths, London and Boston, 1975, p 127–181 76. H. Arend and W. Neuhaus, Die Ha¨rtbarkeit des Stahles, Verlag W. Girardet, Essen, 1955 (in German) 77. D.H. Breen and G.H. Walter, ComputerBased System Selects Optimum Cost Steel—I, Source Book on Materials Selection, Vol I, American Society for Metals, 1977, p 182–185 78. D.H. Breen, G.H. Walter, and C.J. Keith, Computer-Based System Selected Optimum Cost Steels—II, Source Book on Materials Selection, Vol II, American Society for Metals, 1977, p 186–192 79. D.H. Breen, G.H. Walter, and C.J. Keith, Computer-Based System Selected Optimum Cost Steels—III, Source Book on Materials Selection, Vol III, American Society for Metals, 1977, p 193–197 80. D.H. Breen, G.H. Walter, and J.T. Sponzilli, Computer-Based System Selected Optimum Cost Steels—IV, Source Book on Materials Selection, Vol IV, American Society for Metals, 1977, p 198–201 81. D.H. Breen, G.H. Walter, and J.T. Sponzilli, Computer-Based System Selected Optimum Cost Steels—V, Source Book on Materials Selection, Vol V, American Society for Metals, 1977, p 202–206

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82. H. Matthes, Novel Process of Quality Control during Inductive Hardening Process, 17th ASM Heat Treating Society Conference Proceedings, Including the First International Induction Heat Treating Symposium, D.L. Milan, D.A. Poteet, G.D. Pfaffmann, V. Rudnev, A. Muehlbauer, and W.A. Albert, Ed., ASM International, 1989, p 727–733 83. J. Grum, Input and Output Control of Steel Intended for Surface Hardening, 17th ASM

Heat Treating Society Conference Proceedings, Including the First International Induction Heat Treating Symposium, D.L. Milan, D.A. Poteet, G.D. Pfaffmann, V. Rudnev, A. Muehlbauer, and W.A. Albert, Ed., ASM International, 1989, p 763–768 84. P. Birk, Ha¨rtenscnutzmassen bei der Einstazha¨rtung, Ha¨rt.-Tech. Mitt., Herausgegeben von Reibensahm, Vol 10 (No. 3), 1957, p 9–19 (in German)

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Failure Analysis of Steel Welds J.H. Devletian, Portland State University D. Van Dyke, MEI-Charlton, Inc.

FAILURE ANALYSIS OF STEEL WELDS may be divided into three categories. They include failures due to: design deficiencies, weldrelated defects usually found during inspection, and failures in field service. Failures due to design deficiencies usually result in a ductile overload due to overstressing the component or exposing the component to service conditions well beyond expected performance requirements. Weld-related defects, which fall outside the quality-acceptance criteria specified by the applicable welding code, contribute to premature failures upon testing as well as in service. Failures in the field may be insidious, because an unanticipated fracture mechanism has been causing problems without the designer’s knowledge. In this chapter, the failures due to various discontinuities in the steel weldment are emphasized. These include poor workmanship, a variety of hydrogen-assisted cracking (HAC) failures, stress-corrosion cracking, fatigue, and solidification cracking in steel welds.

butt welds are listed in Table 1 and illustrated in Fig. 1. In this figure, discontinuities 1 through 7 are caused by poor workmanship. These include (1) porosity, (2) inclusions, (3) lack of fusion, (4) lack of penetration, (5) undercut, (6) underfill, and (7) overlap. Discontinuities 8, 9, and 10 are associated with the rolling operation in the steel mill. These include (8) laminations, (9) delaminations, and (10) seams and laps in the steel plate. The most insidious discontinuities are those that cause brittle cracking due to metallurgical origins, such as HAC, solidification cracking, or stress-corrosion cracking. Such brittle failures are characterized by fracture surfaces exhibiting intergranular, interdendritic, or cleavage modes of fracture, which occur at stress levels well below the yield stress. Additionally, even failures by quasi-cleavage and/or

Table 1 Common types of weld discontinuities illustrated in Fig. 1 No.

Discontinuities in Steel Welds

1

By far, the most common steel welding problems are associated with poor workmanship, such as lack of fusion, lack of penetration, porosity, undercut, arc strikes, and others. The number and size of such discontinuities may be cause for rejection by most codes and can act as stressconcentration sites to reduce weld joint strength and promote cracking, such as fatigue failures. A description of the wide variety of discontinuities that can occur in steel weld metal and base metal is provided by many welding codes, one of which is the American Association of State Highway Transportation Officials/American Welding Society (AASHTO/AWS) D1.5 Bridge Welding Code (Ref 1). From this code, rejectable discontinuities that can occur in common

2 3 4 5 6 7 8 9 10 11 12

Discontinuity

Porosity(b) Cluster(1b in Fig. 1) Piping(1d in Fig. 1) Inclusions Non-metallic slag(2b) Lack of fusion(b) Lack of joint penetration(b) Undercut(b) Underfill(b) Overlap(b) Laminations(b) Delaminations(b) Seams and laps(b) Lamellar tears Cracks(b) a. Longitudinal(12a) b. Transverse(12b) c. Crater(12c) d. Throat(12d) e. Toe(12e) f. Root(12f) g. Underbead and HAZ(12g)

Location(a)

W

W W W BM W W BM BM BM BM W, HAZ, BM W, HAZ, BM W W HAZ

(a) W, weld; HAZ, heat-affected zone; BM, base metal. (b) See Fig. 1. Source: Ref 1


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microvoid coalescence can occur at reduced levels of ductility. Such cracking is illustrated in discontinuities 11 and 12a through 12g in Fig. 1. These types of brittle failures are especially insidious, because the fracture occurs at stress levels and ductility levels well below the design parameters. Two examples of brittle failure in steel welds are hydrogen-assisted cracking and solidification cracking. Among the discontinuities illustrated in Fig. 1, HAC can commonly take the form of lamellar tears, longitudinal cracks in the heat-affected zone (HAZ) (12a), transverse cracks (12b), toe cracks in the HAZ (12e), and underbead cracks in the HAZ (12g). Solidification cracking can commonly take the form of longitudinal cracks along the weld centerline (12a), crater cracks (12c), and fissures and microfissures in the weld metal. Example 1: Porosity in Weld Metal. Twopass submerged arc welding was used to butt weld two 19 mm (3/4 in.) thick A709-grade 250 plates. The joint design was a simple square groove butt joint with 1.6 mm (1/16 in.) root opening. It was designed to be welded in two passes, with a single pass on each side of the plate. After the first side was welded, the weld bead exhibited an unusually large amount of reinforcement. Subsequent magnetic particle

Fig. 1

inspection indicated extensive discontinuities in the weld bead. The first pass of the weld was sectioned, polished, and etched in 5% nital. The macrograph of the welded plate is shown in Fig. 2. Extensive rejectable porosity (in accordance with AWS D1.1 Structural Welding Code) developed within the weld metal but was not readily visible on the surface. The welding system was inspected for possible sources of porosity. Inspection revealed that the mating flanges were covered with thick mill scale and were neither ground clean nor preheated prior to welding. The welding electrode and flux combination was F7A2-EM12K. Normally, this flux/wire combination is designed to be tolerant of light rust. However, in this case, the porous mill scale was so thick that it acted as a sink for moisture, oil, and other volatile contaminants. Excessive porosity was caused by welding on heavily rusted steel. Simply rough grinding the faying surfaces to bare metal prior to welding completely eliminated the porosity problem. If rough grinding is not possible, the welding engineer can choose a more active flux that is designed for welding rusted plate. Active fluxes contain strong deoxidizers, such as titanium and aluminum, but should only be used in a single pass.

Schematic illustration of common types of discontinuities in welds. Reproduced with permission of the American Welding Society. See also Table 1. Source: Ref 1

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Failure Analysis of Steel Welds / 505

Example 2: Lack of Penetration. A36 structural steel was butt welded by automatic flux-cored arc welding. A double V-groove joint preparation was used for a two-pass welding operation on 6.4 mm (1/4 in.) thick plate. By visual examination, the weld appeared satisfactory. Full penetration of the butt joint was required. Transverse-to-weld tensile testing revealed inadequate yield strength. The tensile test failure occurred in the weld metal. Normally, previous tensile tests always failed in the unaffected base metal. Radiographic examination of the weld joint and the mandatory transverse-to-weld tensile test showed the weld to have inadequate penetration. Inspection of the fracture surface of the tensile specimen revealed unfused metal at the plate midthickness. The fractured steel exhibited a ductile slant fracture. By measuring the approximate area of fused cross section and area of unfused cross section, it was determined that approximately 80% of the joint thickness carried the load. Even though the weld fracture was ductile, the reduced cross-sectional area of fused metal (due to lack of penetration) did not provide enough sound metal to meet the required joint strength. Welding parameters were changed to ensure full penetration. Example 3: Lack of Side-Wall Fusion. DH36 structural steel for shipbuilding was butt welded by robotic gas metal arc welding. A double V-groove joint preparation was used for a two-pass welding operation. By visual examination, the weld appeared satisfactory. Dyepenetrant testing revealed a discontinuity along one edge of the weld. Face bend testing revealed

Fig. 2

Wormhole or piping porosity in weld metal deposited by submerged arc welding. Plate is 19 mm thick.

a rejectable discontinuity along the same edge of the weld. Visual inspection of the discontinuity in the bend test showed that the base metal and weld metal were not fused at the toe. Subsequent metallographic examination showed that the lack of side-wall fusion existed throughout the length of the weld. Clearly, the robot and fixturing were slightly misaligned. The robot welder directed the welding heat too far on one side of the joint and not enough on the other side. The robot was reprogrammed to provide adequate weave to melt both faying surfaces sufficiently to prevent a recurrence of the problem.

Fatigue of Welded Joints Fatigue cracking is a result of repetitive fluctuating stress causing fracture well below the yield strength of the steel. Factors required for fatigue fracture include a sufficiently high tensile stress, a sufficiently large variation in applied stress, and a sufficient number of cycles of applied stress. Cracks resulting from fatigue account for more than half of all failures. Fatigue fractures are insidious because there is no visible plastic deformation preceding fracture. Fatigue fractures are promoted by stress concentrators such as notches, sharp fillets, corners, holes, threads, splines, keyways, dents, gouges, laps, folds, flakes, and delaminations in plates, sheets, and forgings. Tensile residual stresses arise from punched holes, heat treatment, welding, and so on. Visible beach marks indicate the location of the fracture initiation site. Beach marks are produced when oxidation at the crack tip propagating under an alternating stress is interrupted from time to time, for example, constantspeed machinery that is turned on and off. Also, visible beach marks are generated when non-steady-state loading occurs or with rapid changes in loading, for example, auto/truck components and road signs. Usually, fatigue cracks have multiple initiation sites, particularly in rotating shafts subject to bending loads with stress concentrations. Multiple initiation sites are easily recognized because of the presence of ratchet marks. A ratchet mark ridge or ligament occurs at the merger of two adjacent fatigue cracks that propagate on different planes. When these two fatigue cracks meet, they are connected by a ridge of deformed

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metal called a ratchet mark. For example, three adjacent propagating fatigue cracks will have two ratchet marks connecting them. When viewing a fatigue fracture surface at high magnification of approximately 500 to 5000 · under the scanning electron microscope (SEM), striations are usually visible. In ductile structural materials, crack propagation is by microvoid coalescence, and the crack path becomes perpendicular to the tensile axis. The correlation between striation spacing and tensile load has been well established, where one striation represents one fatigue load cycle. Striation spacing increases with increasing maximum tensile loading. Secondary cracks sometimes occur alongside striations, which better resolves striation spacing. For brittle materials, crack propagation occurs in the cleavage or intergranular mode. Since brittle cracks can grow unstably, the crack path provides less clear correlation between striation spacing and loading. Although striations are less well defined, secondary cracks are helpful in resolving some striations. An absence of observable striations is possible when metals with complex metallurgical microstructures obscure striations, or when the fatigue crack closes and plastically deforms the striations. The latter case usually produces a very shiny fracture surface. Example 4: Fatigue Cracking of Welded Pipe Flange. A short drain pipe from a heatrecovery steam generator failed in service by developing leaks around a circumferential weld near a flange. The pipe was a 25 mm (1 in.) schedule 80 pipe (25 mm, or 1 in., internal diameter by 4.8 mm, or 3/16 in., wall thickness) section 127 mm (5 in.) long with flanges welded onto both ends. The pipe had been in service for 5 years with an operating pressure of 8.3 MPa (1200 psi). Upon failure, the crack completely separated the flange from the pipe. The 127 mm (5 in.) section of pipe with flanges was cut from the boiler and submitted for examination, as shown in Fig. 3. The fracture surface of the pipe is shown in Fig. 4. There are clearly defined ratchet marks in both the top and bottom edges of the fracture surface. Beach marks are present on both the top and bottom of the fracture surface that run toward the center, where there is a brightly colored band. Beach marks are not present in the brightly colored band. Metallographic sectioning of the weld profile showed that the crack initiated at the toe of the weld and proceeded in between the weld filler

metal and HAZ. There was no evidence of weld defects such as undercut or porosity present in the section. Scanning electron microscopy showed fine striations within the beach mark areas of the pipe. The SEM showed a ductile dimpling structure within the brightly colored band of the fracture surface. Energy-dispersive spectroscopy of several areas of the fracture surface revealed only iron, manganese, and oxygen. The pipe fractured due to fatigue cracks that initiated at several locations at the toe of the weld. The fatigue cracks propagated until there was insufficient cross section to support the service loads. The morphology of the fracture suggests the cracking propagated under reversed bending stresses.

Hydrogen-Assisted Cracking Theory In addition to discontinuities resulting from poor workmanship, a very common cause of

Fig. 3

Overview of pipe section. Cracking is visible on right end of the pipe at the toe of the weld. Courtesy of MEICharlton, Inc.

Fig. 4

Overview of pipe end. Ratchet marks and beach marks are clearly visible. Courtesy of MEI-Charlton, Inc.

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failure in welded steel plates, castings, and structural components is HAC. Although HAC of steel weldments has been studied for over a half-century, the mechanism of cracking is still uncertain. Fortunately, judicious “hydrogen management” has permitted the crack-free welding of high-strength steels primarily by preheating the weld to a temperature that must be increased with increasing:

Diffusible hydrogen content Yield strength Hardness Carbon equivalent Thickness Restraint Susceptible microstructure

Cracking can take place in either the weld metal or the HAZ, depending on which is more susceptible. Generally, the peak hardness, strength, and carbon equivalent values of the HAZ are greater than those of the weld metal. So, cracking usually takes place in the HAZ. This is the most common form of HAC and is called underbead cracking. Because HAC is dependent on time for hydrogen atoms to diffuse at room temperature to crack nucleation sites, it is also called delayed cracking. Such cracking can take place anytime, from minutes to several days after the weld has cooled. Several other forms of hydrogen cracking are also possible and are discussed. After over a half-century of research on the deleterious effects of hydrogen in steel, the fundamental mechanism of HAC is still uncertain. Preheating, however, is the most reliable method to prevent weld cracking. Unfortunately, preheating is labor-intensive and costly. More recent studies have shown that preheating temperatures can be either reduced or eliminated if the composition of the steel base metal and filler metal is low in crack-promoting elements. (Ref 2–19). For example, an empirical carbon equivalent formula (CEN), developed by Yurioka et al. (Ref 2), clearly reflects the fact that alloying elements such as carbon and boron produce high CEN values and have a strong detrimental effect on hydrogen crack resistance: CEN= C+ A( C) fSi=24+ Mn=6+ Cu=15 + Ni=20+( Cr+ Mo+ Nb+ V)=5+ 5Bg

where A(C) = 0.75+0.25 tanh{20(C0.12)}, and all elements are in weight percent.

Thus, modern structural steel plate and filler metals contain reduced carbon for increased resistance to hydrogen cracking and, coincidently, for increased fracture toughness. Many low-carbon steels, such as the U.S. Navy’s HSLA-65, HSLA-80, and HSLA-100, are designed to be welded without preheating (Ref 20–22). The presence of hydrogen in welds can cause failure at levels of strength and plastic strain well below design yield strength and ductility. The current models of HAC involve pre-existing defect sites in the weld metal. These initiation sites include inclusions, minor phase particles, microscopic cracks, and other discontinuities. Many theories attempt to explain the brittleness caused by the presence of as little as a few parts per million (ppm) of hydrogen in steel. The decohesion theory of Oriani (Ref 23–29) states that dissolved hydrogen migrates to regions of triaxial tensile stress, and that the cohesive forces between atoms in the iron lattice are reduced in proportion to the interstitial hydrogen concentration. This theory provides for the observed increase in hydrogen solubility at the tip of a crack (Ref 30). Diffusible hydrogen has the additional effect of supplying internal mechanical stresses that produce an apparent softening of the steel. Hydrogen also causes growth of microvoids in undeformed specimens and the initiation of microcracks in specimens deformed to the critical strain (Ref 29). During cooling and shrinkage of a weld, these defect sites become localized regions of triaxial tensile stress concentration. Atomic hydrogen diffuses interstitially to these regions of the expanded lattice. As the concentration of hydrogen increases, the cohesive strength between iron atoms decreases below the local intensified stress level, at which point cracking occurs. A definitive reason why the cohesive strength between iron atoms decreases in the presence of hydrogen atoms has not been fully explained. In 1960, Troiano (Ref 31) suggested that certain sites act to concentrate stress to promote cracking, and that hydrogen reduces the cohesive strength between iron atoms. In a classic 1972 paper (Ref 32), C.D. Beachem presented a model suggesting that the presence of sufficiently concentrated hydrogen dissolved in the lattice just ahead of the crack tip aids whatever deformation processes the microstructures will allow. Intergranular, quasicleavage, or microvoid coalescence fracture modes are dependent on the microstructure, the

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crack-tip stress intensity, and the concentration of hydrogen. That is, hydrogen reduces the stress intensity needed for cracking regardless of the mode of fracture, as shown in Fig. 5. Even the stress intensity needed for HAC by microvoid coalescence is reduced in the presence of hydrogen. Beachem’s model unifies several theories but shows how the stress-sorption and lattice embrittlement models are unnecessary. The model shows that the planar pressure effects are necessary at low stress intensities and are necessary only to augment the driving force from the applied loads. The basic hydrogen-steel interaction appears to be an easing of dislocation motion or dislocation generation, or both. In 2003, Albert et al. (Ref 33) found that hydrogen traps increased with alloying elements. Measurements of diffusible hydrogen in weld metal decreased in steel welds containing increasing amounts of chromium and molybdenum alloying elements. The HAC in steel is dependent on hydrogen trapping, which is the interaction between diffusible hydrogen and lattice defects. Several investigators have examined hydrogen trapping in steel (Ref 34– 40). These traps include point defects (vacancies, solute atoms), dislocations, interfaces and surfaces (grain boundaries, particle-matrix interfaces, cracks, external surfaces), and volume defects (voids, second-phase particles). Gibala and DeMiglio (Ref 34) experimentally measured binding energies (Table 2) and saturabilities of hydrogen traps in AISI 4340 steel, which led them to predict probable susceptibilities to HAC. Both primary and secondary crack paths were explained in terms of the types of microstructural traps that predominate in a

given material. Management of hydrogen traps can have a profound effect on the susceptibility to HAC. In theory, decreasing the amount of diffusible hydrogen by introducing permanent or irreversible traps has the beneficial possibility of reducing the susceptibility to HAC. Olson et al. (Ref 37–40) have shown that effective hydrogen traps can include microscopic particles of Ce2O3, TiC, Y2O3 VC, and NbC. Although these traps would decrease the amount of diffusible hydrogen in the weld, large quantities of such particles would have a deleterious effect on ductility and fracture toughness. Steel can be embrittled by hydrogen if it is stressed at temperatures high enough to allow hydrogen to diffuse to potential embrittlement sites, but also low enough that the hydrogen is not depleted from embrittlement-producing traps. With increasing temperature, steel is expected to exhibit mechanical behavior determined largely by the nature of hydrogen transport among the populations of various hydrogen traps (Ref 34). Increasing the tensile strain or residual tensile stress also increases the ability of the steel to retain hydrogen in body-centered cubic (bcc) iron. Normally, the solubility (s) of hydrogen in bcc iron is given by: s=47:66 p1=2 exp (72:72 · 107=RT)

where s is in ppm by mass, R is the gas constant, T is the absolute temperature, and p is the pressure in atmospheres. For example, at room temperature and pressure, the solid solubility of hydrogen in bcc iron is only 6.7 · 10 4 ppm. Solubility is increased by the presence of traps. For example, by increasing tensile strain, the solid solubility of iron (Ref 30) increased by a

Table 2 Hydrogen trap interactions in steel Hydrogen trap type

Fig. 5

Effect of hydrogen content on hydrogen-assisted cracking (HAC) for microvoid coalescence (MVC), quasi-cleavage (QC), and intergranular (IG) fracture modes. Adapted from Beachem. Source: Ref 32

Interstitial solutes Ti atom Vacancy Screw dislocation core Mixed dislocation core Hydrogen vapor/void Grain boundary Free surface AlN interface Fe3C interface TiC interface Source: Ref 34

Binding energy (Ed), kJ/mol

3–15 26 46 20–30 59 29 59 70–95 65 84 96


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factor given by: c=cn =exp(Asys 1=2 Bsys ) A={2(1+n)V=3RT}{2E=p}1=2 B={2(1+n)V=RT}{2=p}

where c is the hydrogen concentration at the crack tip, cn is the hydrogen concentration in the bulk steel, n is Poisson’s ratio, V is the partial molar volume of hydrogen in iron, E is the elastic modulus, and sys is the yield strength. For a high-yield-strength steel having sys = 1000 MPa, the solubility of hydrogen in steel increases by a c/cn factor of 6.5. The tip of a crack is an excellent site for hydrogen accumulation, because the strain is greater than that associated with sys. Although the mechanism for HAC is still debated, equations to prevent HAC have been developed empirically. For example, the preheat parameter (Pw) and the hydrogen accumulation parameter (PHA), were developed from experimental work in Japan by Yurioka et al. (Ref 15, 41) as practical tools to empirically predict a preheating temperature to prevent HAC: Pw =Pcm +HD =60+R=40,000

where Pcm is the Ito-Bessyo (Ref 42) carbon equivalent given as: P cm =C+Si=30+(Mn+Cu+Cr)=20+Ni=60 +Mo=15+V=10+5B

where HD is the diffusible hydrogen content in mL/100 g, elements are in weight percent, and R is the restraint intensity in MPa. The HAC can be avoided if the value of Pw50.3. In the Pw equation, the effects of alloy composition, hydrogen content, and restraint are taken into account empirically. Consumables manufacturers have recognized that filler metals containing less carbon than the steel plate would produce not only enhanced weld metal toughness at required strength levels but also greater resistance to HAC. In fact, virtually all of the older algorithms for determining preheating temperatures were based on the composition of the base metal and not the weld metal. Cracking of the HAZ by HAC was so common in structural steels that it was assumed

that the coarse-grainced HAZ would always be the location of maximum susceptibility to cracking. Using the advantages of modern lowcarbon steels and consumables, Nippon Steel (Ref 4) designed a series of commercial lowcarbon steels for line pipes that could be welded while maintaining high strength and toughness with equally low-carbon filler metals. These pipeline steels included the X-65, X-70, and X-80 grades, which contain very low carbon (50.03%), high manganese for strength and to control the bainite transformation, and 0.001% boron to suppress the proeutectoid ferrite nucleation at austenite grain boundaries. It is desirable to have a large amount of acicular ferrite in the weld metal for optimal strength and toughness as well as good resistance to HAC (Ref 43, 44). Microconstituents detrimental to weld metal toughness and possibly increased susceptibility to HAC include grain-boundary ferrite, martensite, and side-plate ferrite, because these structures provide a continuous path for cleavage crack propagation.

Types of Hydrogen-Assisted Cracking Hydrogen-assisted cracking can appear in four common forms:

Underbead or delayed cracking Weld metal fisheyes Ferrite vein cracking Hydrogen-assisted reduced ductility

As mentioned earlier, the mechanism of HAC is not clear, but management of hydrogen and the prevention of HAC are well established. Preheating the weld area prior to and during welding provides the most reliable resistance to HAC. There are many empirically-derived methods to calculate preheat temperatures to prevent HAC. All of the various types of HAC can be avoided by good welding practice. The forms of HAC are discussed in the following sections. Underbead or Delayed Cracking By far, the most common form of HAC is underbead or delayed cracking, schematically illustrated as discontinuity 12g in Fig. 1 and described in Table 1. Typically, this form of cracking occurs in the coarse-grained HAZ up to 72 h after the weld has cooled. This is because the HAZ typically has higher carbon content and

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generally a higher carbon equivalent level than the weld metal. Even though the source of hydrogen is virtually always from the welding consumable, atomic hydrogen rapidly diffuses to crack nucleation sites in the HAZ. The diffusion of hydrogen explains the time-dependent nature of HAC, thus the name delayed cracking. The coarse-grained HAZ is the zone adjacent to the weld and represents a region that has been heated to nearly the melting temperature, followed by rapid cooling. Because of its higher carbon content and large austenite grain size, the coarse-grained HAZ develops a significantly higher hardness than the weld metal. The HAZ will transform to martensite upon cooling if the carbon equivalent is high enough. Since the harder HAZ is more susceptible to HAC than either the weld metal or the unaffected base metal, the cracking in a butt weld is typically confined to a narrow strip of metal immediately adjacent to the weld bead. Cracking in fillet welds occurs at the toe of the weld because that is the location of highest stress concentration. Toe cracking is schematically illustrated as discontinuity 12e in Fig. 1 and described in Table 1. Example 5: Underbead Cracking. Crosscountry line pipe is welded continuously for long distances. At regular intervals, a flange needs to be welded onto the pipe for coupling to a valve or other device. Recently, a section of pipe was removed from service because of cracking that had occurred in the toes of the fillet welds joining the flange to the pipe. The pipe was 203 mm (8 in.) outside diameter by 6.4 mm (0.25 in.) wall thickness, and the flange was 205 mm (8.1 in.) inside diameter by 305 mm (12 in.) outside diameter by 18 mm (0.71 in.) thick. Since the pipe was only 6.4 mm thick, the weld was not preheated. Cracks measuring approximately 10 cm long developed at the toes of the fillet welds on the flange side, as shown in Fig. 6. From the illustration in Fig. 6, the cracking occurred only in the toes of the two fillet welds on the flange side. No cracking was observed at the toes of the two fillet welds on the pipe side. A metallographic section of the crack, shown in Fig. 7, clearly reveals that the crack was confined to the brittle martensitic HAZ on the flange side. Scanning electron microscopy of the cracked area clearly shows an intergranular mode of fracture (Fig. 8). Chemical analysis of the pipe and flange in Table 3 revealed that an incorrect steel was used for the flange. The

Fig. 6

Underbead cracking at the toe of the fillet weld on the flange

Fig. 7

Toe cracking on the flange side of the flange-to-pipe fillet weld, showing the weld metal, heat-affected zone, and unaffected base metal. Cracking occurred in the martensitic (white) heat-affected zone of the flange.

Fig. 8

Fracture surface of flange failure in the as-received condition. Intergranular fracture is shown as well as debris retained from the field.

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correct flange steel was supposed to be a lowcarbon steel. Instead, the flange was a 0.8% (high) carbon steel, which was very susceptible to HAC in the HAZ when welded without preheating. Cracking in the HAZ on the flange side of the fillet weld was due to the mistaken use of a high 0.8% C steel instead of the specified low-carbon steel. Quality control measures need to be followed to prevent mixed steels from being used. Weld Metal HAC and Fisheyes Welding of modern low-carbon steels often results in HAZs with greater resistance to HAC. Thus, the weld metal composition is now as susceptible as the HAZ. If the weld metal contains sufficient diffusible hydrogen content, has high yield strength, and is in a highly stressed condition, the susceptibility of such weld metal to HAC is very possible. For example, in the line-pipe industry, new thermomechanicalcontrolled processing steels achieve high yield strength through thermal processing in the rolling mill, so that the carbon content and carbon equivalent levels for a given yield strength have dropped substantially. With this reduction in carbon equivalent, the susceptibility to HAZ cracking has also declined significantly. Since the as-deposited weld metal achieves strength primarily through alloying, the weld metal is now very susceptible to HAC. Often, field welding of X-65 and X-70 line pipe is performed with high-hydrogen E8010G cellulosic electrodes. In this case, the weld metal yield strength is greater than both the HAZ and unaffected base metal. Thus, the weld metal has become the weak link and is most susceptible to HAC. The strong influence of hydrogen on weld metal cracking can be observed in tensile testing and bend testing as well as in failures of welds subject to slowly applied tensile stress. Fisheyes occur typically on the fracture surface of steel

all-weld-metal tensile specimens that fail due to HAC. In tensile testing, fisheyes reduce the weld metal ductility measurements, such as percent elongation and percent reduction of area. Fisheyes are local areas within the weld that are more hardenable due to solute banding, cellular, or dendritic segregation of alloying elements (Ref 45). These initiation areas may also be richer in localized hydrogen due to their proximity to hydrogen traps such as inclusions. Since these alloy-rich segregated areas are more susceptible to brittle HAC, small, localized brittle-fracture zones appear visually on a tensile test fracture surface as bright round spots surrounded by gray ductile fracture. The bright round spot may consist of a local region of typically intergranular or possibly cleavage fracture surrounded by ductile dimpled failure. Both intergranular and cleavage failures are brittle fracture modes that appear much more brightly than the surrounding material, which is ductile dimpled and gray-appearing. Example 6: Fisheyes on Fracture Surface. A high-strength steel, HSLA-100, was butt-welded with a matching-strength filler metal using gas metal arc welding (GMAW) and argon-5%CO2 shielding gas at a heat input of 1.1 kJ/mm (28 kJ/in.) without preheating. The filler metal contained Fe-0.03%C-1.4%Mn3%Ni-0.7%Mo. Because of the very low carbon content, the weld metal hardness did not exceed 24 HRC. Tensile test results showed inadequate ductility of only 8% elongation. Upon examining the fracture surface of the tensile specimen, multiple fisheyes were observed, as shown in Fig. 9. An SEM image of

Table 3 Chemical analysis of the flange and pipe Chemical element

Flange

Pipe

C Mn Si Ni Cr S P Al Nb

0.80 0.67 0.25 0.01 0.23 0.13 0.018 ... ...

0.07 1.10 0.24 0.01 0.01 0.005 0.017 0.038 0.02

Fig. 9

Brittle fisheyes appear as bright spots in a gray ductile matrix.

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the center of a bright fisheye exhibited brittle intergranular fracture, as shown in Fig. 10. The SEM images outside of the fisheyes showed ductile dimpled microvoid coalescence. Welds were repeated to determine the amount of diffusible hydrogen in similar welds in accordance with AWS A4.3 (a standard welding test for diffusible hydrogen). Despite the fact that GMAW is known as a very low-hydrogen process, the value obtained for diffusible hydrogen was 9 mL/100 g, which was far greater than expected. It was found that the filler metal manufacturer used excess hydrocarbon-base lubricant during the wire-drawing operation because of the high strength of the wire. Fisheyes were caused by excessive amounts of diffusible hydrogen in the weld metal due to the lubricant residue on the filler metal. Ferrite Vein Cracking A very unexpected form of HAC is ferrite vein cracking, which can occur in slowly cooled electroslag welds. In recent studies of electroslag welding of 50 and 75 mm thick low-carbon steel at Portland State University (Ref 46), ferrite vein cracking of A709-grade 245 steel occurred only in welds that were made with flux and/or filler metal known to have high moisture content. Although the mechanism is not certain, diffusible hydrogen causes the ferrite at prioraustenite grain boundaries to crack under the residual tensile stress produced by contraction during weld cooling. This is very unusual, because typical HAC is associated with hard martensitic microstructures. Ferrite was always thought to be immune to HAC because of its low

Fig. 10

SEM image of center of fisheye showing intergranular fracture

strength and low hardness. It was also found that nickel alloying additions tended to promote HAC in the form of ferrite vein cracking, while an equivalent amount of molybdenum resisted cracking. Both nickel and molybdenum are essential alloying elements for enhancing fracture toughness in both the weld metal and base metal. The mechanism by which nickel and molybdenum appear to have virtually opposite effects on susceptibility to HAC is not known. Example 7: Ferrite Vein Cracking in High-Heat-Input Welds. Single-pass fullpenetration electroslag welds were deposited on 50 mm (2 in.) thick ASTM A588 steel using a heat input of 42 kJ/mm (1070 kJ/in.) for bridge applications. The ASHTO/AWS D1.5 Bridge Welding Code required both radiographic and ultrasonic testing (UT) of the completed welds. The UT revealed possible indications of cracking around the weld center. The weld metal was sectioned for metallographic examination, and ferrite vein cracking was found, as shown in Fig. 11. Clearly, cracking was confined to the grainboundary ferrite, which was nucleated at the prior-austenitic grain boundaries.

Fig. 11 grade 50W

Ferrite vein crack occurring in the prior-austenite grain boundaries of weld metal deposited on A709-

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Review of the welding procedure revealed that the flux and tubular wire used in this process were not baked prior to welding. The unbaked flux was a source of moisture. Despite its extremely high weld heat input and the use of mild steel electrodes, electroslag weld metal has been shown to be susceptible to HAC. The fabricated metal-cored tubular wire was also unbaked. Under the welding heat, moisture in the flux and filler metal produces atomic oxygen and hydrogen. Since electroslag welds are large single-pass deposits, the weld center is under substantial tensile stress. The combination of ample diffusible hydrogen and high-shrinkage tensile residual stress at the weld center provides the necessary ingredients for ferrite vein cracking. Subsequent welds were made with flux that was baked to 204 C (400 F) and a new metalcored wire that was baked at an elevated temperature prior to shipment. The resulting welds have since been free of ferrite vein cracking. Although the mechanism of HAC of grainboundary ferrite is unknown, elimination of cracking was achieved by reducing the sources of moisture or hydrogen.

exhibited an intergranular mode of fracture. Scanning electron microscopy at 50 · revealed that the crack propagated intergranularly along the prior-austenite grain boundaries, as shown in Fig. 13. However, at approximately 4000 · , the intergranular fracture surface exhibited shallow dimples, as shown in Fig. 14. This dimpled intergranular mode of fracture is due to the weakness of the ferrite envelopes surrounding each prior-austenite grain, as shown in Fig. 15. The presence of diffusible hydrogen caused a reduction of ductility of grain-boundary ferrite sufficient to fail the bend test. The Beachem diagram in Fig. 5 shows that microvoid coalescence can also be adversely affected by diffusible hydrogen. Using low-hydrogen practices, such as baking the flux prior to use, eliminated the cracking problem during bend testing.

Stress-Corrosion Cracking of Steel Stress-corrosion cracking of steels is possible when the steel is subject to both adequate tensile

Hydrogen-Assisted Reduced Ductility This form of HAC occurs when the damage due to diffusible hydrogen is not sufficient to cause cracking in the weldment but is sufficient to cause reduced ductility in subsequent tensile and bend tests. This is a clear illustration of the principles reported by Beachem (Ref 32) and shown schematically in Fig. 5. All fracture modes become more severe with increasing diffusible hydrogen. Even the ductility associated with ductile microvoid coalescence is substantially reduced in the presence of diffusible hydrogen. Example 8: Failure to Pass Bend Tests due to Hydrogen. Multipass submerged arc welds deposited on 50 mm thick A588 steel were subject to inspection in accordance with the AWS D1.1 Structural Welding Code. The following tests were performed for the procedure qualification welds: tensile testing, bend testing, Charpy V-notch impact toughness testing, as well as both ultrasonic and radiographic testing. All of these tests were passed successfully except the guided bend test. As shown in Fig. 12, the side-bend specimen cracked well before the prescribed bend radius could be achieved. Visual inspection and low-magnification optical microscopy of the cracked bend specimen

Fig. 12

Side-bend test failure of weld

Fig. 13

Scanning electron micrograph of fracture surface of bend failure

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stress (but below the yield strength) and an aggressive environment. The absence of one of these elements will eliminate stress-corrosion cracking. Welds are particularly good sites for stress-corrosion cracking, because substantial tensile residual stresses are always present. The shrinkage stresses associated with the solidification and cooling of welds produce near-yield tensile residual stresses in and around the weld. Relatively mild chemical environments can activate the stress-corrosion cracking process. Environments known to cause stress-corrosion cracking of plain carbon and alloy steels include liquefied ammonia, hydrogen sulfide, molybdenum disulfide, sodium hydroxide sour gas, high pH values, nitrate solutions, and many other corrosive environments. Stress-corrosion cracking of carbon steel can even take place in pure water under high temperature and pressure.

Fig. 14

Higher-magnification image of fracture surface in Fig. 13 showing dimpled intergranular fracture

Fig. 15

Optical microscopy of grain structure of electroslag weld metal. Original magnification: 50 ·

Although many theories for stress-corrosion cracking have been suggested, only two appear to be the basis for such cracking. These are the stress-sorption theory and the electrochemical theory. The stress-sorption theory states that damaging substances in the environment are chemically absorbed onto the surface of the steel, causing a reduction in the cohesive bonding force between iron atoms. Only an atom-thick surface layer is needed to seriously affect the bonding forces between surface atoms. There is a threshold stress necessary to initiate stresscorrosion cracking. In some ways, this mechanism resembles HAC in steels (discussed earlier), where only a few ppm of hydrogen in concert with tensile stress are needed to reduce the cohesive bonds between iron atoms and cause cracking. In stress-corrosion studies, acoustic emission sensors have been attached to the cracking sample to monitor the propagation of the crack. Strong acoustic emissions were emitted and recorded each time the advancing crack jumped or burst. Acoustic emission sensors used to monitor HAC of steel displayed a similar jump or burst behavior. The electrochemical theory involves the setting up of galvanic cells within the microstructure of the steel. Anodic dissolution paths are produced along concentration gradients in the metal or in grain boundaries. When the grain boundaries are anodic to the bulk of the metal, tensile stresses (although below yield) are necessary to continue the cracking process in order to open up dissolved pathways for further penetration by the corrosive environment. As evidence of the electrochemical nature of this cracking process, stress-corrosion cracking can be stopped by applying cathodic protection. As soon as cathodic protection is removed, stresscorrosion cracking continues. Example 9: Stress-Corrosion Cracking of a Weld. After 30 years in service, a low-pressure steam supply line developed a reoccurring cracking problem in a circumferential butt weld. The weld was on a 25 cm (10 in.) supply line that carried 0.4 MPa (55 psig), 205 C (400 F) steam to a paper machine. This line was 30 m (100 ft) downstream from a spray attemperator that cooled higher-temperature steam by spraying boiler feedwater into the line. When cracks were initially discovered, they were ground out and rewelded. The repair welds reportedly cracked after only a few days in service.

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A visual examination revealed that neither the original welds nor the repair welds penetrated through to the inside of the pipe. The cracking consisted primarily of circumferential cracks down the center of the weld. In addition, there were a few locations where short longitudinal cracks emanated from the primary circumferential crack. Macroscopically, the cracks were generally irregular and branching. Sections of the cracked region were prepared for metallographic examination of the weld geometry and crack morphology, as shown in Fig. 16. A microhardness scan of the base metal revealed a hardness of 79 HRB, with a hardness of 87 to 94 HRB in the original weld. The repair weld was significantly harder, with a hardness of 22 to 34 HRC. Metallographic investigations revealed an intergranular branching morphology, as shown in Fig. 17. Scanning electron microscopy also showed the presence of several secondary branching cracks emanating from the primary crack. These cracks also proceeded in an intergranular fashion. Energy-dispersive spectroscopy did not reveal the presence of foreign materials at the fracture surface in a measurable quantity. The pipe fractured due to stress-corrosion cracking (SCC) precipitated by the presence of geometric stress concentrations and high residual tensile stresses in the weld. The SCC is an environmentally induced cracking mechanism that can occur in a susceptible material in the presence of tensile stress and an aggressive chemical agent. Although the particular agent involved could not be identified, this is not unusual, because small amounts of a caustic agent can often cause cracking in the proper conditions. Due to the location of the cracking, it is likely that the agent entered the line in the spray

Fig. 16.

Cross section of weld at butt joint. Etchant: 2% nital. Courtesy of MEI-Charlton, Inc.

attemperator process. It is likely that an upset in the chemistry of the boiler feedwater at some time in the past contaminated the downstream lines and led to the SCC in this instance. The consequences of additional steam line failures need to be evaluated. Given the nature of the cracking, inspection methods capable of detecting the cracks in early stages are limited. It may be most cost-effective to replace the steam lines downstream of the attemperator and reevaluate the methods used for ensuring the proper chemistry of the boiler feedwater. Future repairs should ensure complete crack removal and full penetration welds using proper preheat and interpass temperatures to minimize hardness gradients.

Solidification Cracking of Steel Solidification cracking is one of several forms of hot cracking. Solidification cracking in steel and steel alloys occurs near the end of the solidification process and is caused by two dominant factors: tensile stress acting on the weld during solidification, and a large temperature range between the solidus and liquidus temperatures or the presence of low-melting impurities such as sulfer and phosphorus. The tensile stress acting on the weld can arise from either shrinkage tensile stresses produced during solidification and cooldown, or externally applied tensile stress or tensile restraint stress. The effect of the liquidus-to-solidus temperature range has been

Fig. 17

Micrograph of the crack near the weld root. Original magnification: 100 · . Etchant: 2% nital. Courtesy of MEI-Charlton, Inc.

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dealt with generally by several empirical solidification cracking equations. For example, Matsuda (Ref 47) developed a very popular parameter for solidification cracking of steels called Lt, where increasing Lt increased susceptibility to cracking: Lt =70(C Si=12 Mn=9+3P+4S+Ni=23 +Cr=35+Mo=70)

where all elements are in weight percent. Clearly the effects of carbon and alloying elements such as manganese, molybdenum, chromium, and nickel were assumed to be linear, as shown in the Lt equation. As the Lt equation suggests, decreasing carbon content has always been assumed to decrease solidification cracking susceptibility in steel weld metal. Prediction formulas for solidification cracking show the effect of carbon on cracking to be linear.

Fig. 18

Maximum crack length (MCL) as a function of carbon content in iron weld metal obtained in transvarestraint tests at 4% augmented strain. Source: Ref 53, 54

Fig. 19

Solidification cracking in weld metal

However, recent research has shown that the effect of carbon on solidification cracking of low-carbon steels is far more complex and nonlinear than predicted by the Lt formula. For example, Masumoto (Ref 48) showed that solidification cracking was enhanced for carbon contents 40.1%. Conversely, Ohshita et al. (Ref 41) reported that cracking was enhanced for carbon 50.1% and that nickel additions were beneficial in reducing the cracking effect of carbon, Karjalainen et al. (Ref 49) surveyed the technical literature and reported that there was a least-susceptible range of carbon contents between 0.1 and 0.17%. Within this range, the cracking susceptibility was minimized. Ichikawa et al. (Ref 50) reported peak solidification cracking susceptibility at 0.035% C, followed by enhanced cracking when the carbon content exceeded 0.1%. Most recently, Kim et al. (Ref 51) and Won et al. (Ref 52) showed a peak in solidification cracking susceptibility at approximately 0.10% C. It was apparent from the literature that the effect of carbon on solidification cracking susceptibility of steel weld metal required further study. In the most recent work by Shankar and Devletian (Ref 53, 54), the effect of carbon on solidification cracking was nonlinear, with a peak in cracking susceptibility at 0.1% C, as shown in Fig. 18. In this study, testing was performed on high-purity iron-carbon alloy castings using the varestraint and transvarestraint tests. Maximum crack distance and maximum crack length were measured at a 4% augmented strain.

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The data (schematically plotted in Fig. 18) clarify some of the conflicting issues in the literature. Solidification cracking susceptibility is controlled by the brittle temperature range (BTR), the d?c transformation stress, and the iron-carbon peritectic reaction. According to Shankar and Devletian (Ref 53, 54), there are four distinct %C ranges that produced characteristic solidification behavior. These include:

Region 1: less than ~0.09% C (maximum solid solubility of carbon in d-iron) Region 2: ~0.09 to 0.11% C (maximum solidification cracking) Region 3: ~0.11 to 0.16% C Region 4: greater than 0.16% C (iron-carbon peritectic point)

In region 1, there was no cracking below 0.01% C in transvarestraint tests. This is because the solidus/liquidus temperature range was negligible. However, as the carbon content increased, the cracking susceptibility increased rapidly up to approximately 0.06% C. The cracking dropped slightly at 0.075% and then continued to increase with carbon content up to 0.09%, due to the increasing solidus/liquidus temperature range. In region 2, a large peak in solidification cracking was observed. This critical cracking peak, centered at approximately 0.1% C, was found to be due to the simultaneous action of three factors: the maximum solidus/ liquidus temperature range, the d?c transformation stresses, and the occurrence of the BTR. At 0.1% C, cracking occurred with a minimum critical strain and low fracture stress. In region 3, solidification cracking decreased

with increasing carbon content because of the decreasing solidus/liquidus temperature range. In region 4, the solidification cracking susceptibility increased due to the increasing solidus/ liquidus temperature range. Example 10: Solidification Cracking of Steel Weld. Welds were deposited by fluxcored arc welding on 12 mm thick AISI/SAE 1020 steel plate at high travel speeds for maximum cost-effectiveness. Weld joints were highly restrained during welding to prevent distortion. Visible longitudinal centerline cracks were observed, as shown in Fig. 19. The portions of the weld seam that were not visibly cracked failed the root bend test. The cracked specimen was broken open and observed under the SEM. Rejected bend-test specimens were also broken open for examination by SEM. The crack surface clearly showed that the mode of fracture was solidification cracking, as shown in Fig. 20. Spectrographic analysis of the weld metal admixture revealed a carbon content of 0.1%. Subsequent welds were deposited with reduced welding travel speed in order to reduce the length of the teardrop shape of the weld puddle. Subsequent welds deposited at the reduced welding speed were crack-free. Centerline cracking failure was caused by solidification cracking. Decreasing welding speed reduced susceptibility to solidification cracking in the weld metal. Reducing restraint during welding and reducing the weld metal carbon content (or Matsuda’s Lt factor, mentioned previously) would have also decreased the occurrence of solidification cracking.

REFERENCES

Fig. 20

Scanning electron micrograph of fracture surface showing solidification cracking

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APPENDIX 1

Metric Conversion Guide THIS APPENDIX is intended as a guide for expressing weights and measures in the Systeme International d’Unite´s (SI). The purpose of SI units, developed and maintained by the General Conference of Weights and Measures, is to provide a basis for worldwide standardization of units and measure (Table A1.1). For more information on metric conversions (Table A1.2), the reader should consult the following references:

“Standard for Metric Practice,” E380, Annual Book of ASTM Standards, American Society for Testing and Materials, 1916 Race Street, Philadelphia, PA 19103 “Metric Practice,” ANSI/IEEE 268– 1982, American National Standards

Institute, 1430 Broadway, New York, NY 10018 The International System of Units, SP 330, 1986, National Institute of Standards and Technology. Order from Superintendent of Documents, U.S. Government Printing Office, Washington, DC 20402-9325 Metric Editorial Guide, 4th ed. (revised), 1985, American National Metric Council, 1010 Vermont Avenue NW, Suite 1000, Washington, DC 2005–4960 ASME Orientation and Guide for Use of SI (Metric) Units, ASME Guide SI 1, 9th ed., 1982, The American Society of Mechanical Engineers, 345 East 47th Street, New York, NY 10017

Table A1.1 Base, supplementary, and derived SI units Measure

Unit

Symbol

Base units Amount of substance Electric current Length Luminous intensity Mass Thermodynamic temperature Time

mole ampere meter candela kilogram kelvin

mol A m cd kg K

second

s

radian steradian

rad sr

gray meter per second squared becquerel

Gy m/s2 Bq

radian per second squared radian per second square meter farad mole per cubic meter

rad/s2 rad/s m2 F mol/m3

siemens ampere per square meter kilogram per cubic meter

S A/m2 kg/m3

Supplementary units Plane angle Solid angle Derived units Absorbed dose Acceleration Activity (of radionuclides) Angular acceleration Angular velocity Area Capacitance Concentration (of amount of substance) Conductance Current density Density, mass

Measure

Electric charge density Electric field strength Electric flux density Electric potential, potential difference, electromotive force Electric resistance Energy, work, quantity of heat Energy density Entropy Force Frequency Heat capacity Heat flux density llluminance Inductance Irradiance Luminance Luminous flux Magnetic field strength Magnetic flux Magnetic flux density Molar energy Molar entropy Molar heat capacity Moment of force Permeability Permittivity

(continued)

Unit

Symbol

coulomb per cubic meter volt per meter coulomb per square meter volt

C/m3 V/m C/m2 V

ohm joule

V J

joule per cubic meter joule per kelvin newton hertz joule per kelvin watt per square meter lux henry watt per square meter candela per square meter lumen ampere per meter weber tesla joul per mole joule per mole kelvin joule per mole kelvin newton meter henry per meter farad per meter

J/m3 J/K N Hz J/K W/m2 lx H W/m2 cd/m2 lm A/m Wb T J/mol J/mol K J/mol K Nm H/m F/m


18/8/2008 4:24PM Plate # 0

pg 522


Table A1.1 (continued) Measure

Power, radiant flux Pressure, stress Quantity of electricity, electric charge Radiance Radiant intensity Specific heat capacity Specific energy

Unit

Symbol

watt pascal coulomb

W Pa C

watt per square meter steradian watt per steradian joule per kilogram kelvin joule per kilogram

W/m2 sr W/sr J/kg K J/kg

Measure

Specific entropy Specific volume Surface tension Thermal conductivity Velocity Viscosity, dynamic Viscosity, kinematic Volume Wavenumber

Unit

Symbol

joule per kilogram kelvin cubic meter per kilogram newton per meter watt per meter kelvin meter per second pascal second square meter per second cubic meter 1 per meter

J/kg K m3/kg N/m W/m K m/s Pa s m2/s m3 l/m

Table A1.2 Conversion factors To convert from

to

multiply by

Angle degree

rad

1.745 329 E02

mm2 cm2 m2 m2

6.451 6.451 6.451 9.290

600 E+02 600 E+00 600 E04 304 E02

Nm Nm Nm Nm

1.129 1.355 9.806 7.061

848 E01 818 E+00 650 E+00 552 E03

Bending moment or torque per unit length lbf in./in. lbf ft/in.

N m/m N m/m

4.448.222 E+00 5.337 866 E+01

A/cm2 A/mm2 A/m2

1.550 003 E01 1.550 003 E03 1.076 400 E01

Current density A/in.2 A/in.2 A/ft2

T mWb S A/m Vm mV m

1.000 1.000 1.000 7.957 1.000 1.662

000 E04 000 E02 000 E+00 700 E+01 000 E02 426 E 03

Energy (impact, other) ft lbf Btu (thermochemical) cal (thermochemical) kW h Wh

J J J J J

1.355 1.054 4.184 3.600 3.600

818 E+00 350 E+03 000 E+00 000 E+06 000 E+03

L/min L/min L/min L/min

4.719 2.831 6.309 3.785

475 E01 000E+01 020 E+02 412 E+00

N N kN N

4.448 4.448 8.896 9.806

222 E+00 222 E+03 443 E+00 650 E+00

Flow rate ft3/h ft3/min gal/h gal/min

N/m N/m

Fracture toughness pffiffiffiffiffiffi ksi in:

MPa

1.459 390 E+01 1.751 268 E+02 pffiffiffiffi m

1.098 800 E+00

Heat content kJ/kg kJ/kg

2.326 000 E+00 4.186 800 E+00

J/in. kJ/in.

J/m kJ/m

3.937 008 E+01 3.937 008 E+01

Length ˚ A min. mil in. in. ft yd mile

nm mm mm mm cm m m km

1.000 2.540 2.540 2.540 2.540 3.048 9.144 1.609

oz lb ton (short, 2000 lb) ton (short, 2000 lb) ton (long, 2240 lb)

kg kg kg kg · 103(a) kg

2.834 952 E02 4.535.924 E01 9.071 847 E+02 9.071 847 E01 1.016 047 E+03

Mass per unit area oz/in.2 oz/ft2 oz/yd2 lb/ft2

kg/m2 kg/m2 kg/m2 kg/m2

4.395 3.051 4.882 4.882

kg/m kg/m

1.488 164 E+00 1.785 797 E+01

kg/s kg/s kg/s

1.259 979 E04 7.559 873 E03 4.535 924 E01

Heat input

000 E01 000 E02 000 E+01 000 E+01 000 E+00 000 E01 000 E01 300 E+00

000 E+01 517 E01 428 E+00 428 E+00

Mass per unit length lb/ft lb/in. Mass per unit time

Force lbf kip (100 lbf ) tonf kgf

multiply by

Mass

Electricity and magnetism gauss maxwell mho Oersted V cm V circular-mil/ft

lbf/ft lbf/in.

Btu/lb cal/g

Bending moment or torque lbf in. lbf ft kgf m ozf in.

to

Force per unit length

Area in.2 in.2 in.2 ft2

To convert from

lb/h lb/min lb/s

Mass per unit volume (includes density) g/cm3 lb/ft3 lb/ft3 lb/in.3 lb/in.3 (continued)

(a) kg · 103 = 1 metric ton or 1 megagram (Mg)

kg/m3 g/cm3 kg/m3 g/cm3 kg/m3

1.00 000 E+03 1.601 846 E02 2.767 846 E+01 2.767 846 E+04 2.767 990 E+04


18/8/2008 4:25PM Plate # 0

pg 523

Appendix 1: Metric Conversion Guide / 523

Table A1.2 (continued) To convert from

to

multiply by

Btu/s Btu/min Btu/h erg/s ft lbf/s ft lbf/min ft lbf/h hp (550 ft lbf/s) hp (electric)

kW kW W W W W W kW kW

1.055 1.758 2.928 1.000 1.355 2.259 3.766 7.456 7.460

056 E+00 426 E02 751 E01 000 E07 818 E+00 697 E02 161 E04 999 E01 000 E01

to

multiply by

2

W/in.

W/m

2

1.550 003 E+03

Press capacity See Force Pressure (fluid) atm (standard) bar in. Hg (32 F) in. Hg (60 F) lb/in.2 (psi) torr (mm Hg, 0 C)

Pa Pa Pa Pa Pa Pa

1.013 1.000 3.386 3.376 6.894 1.333

250 E+05 000 E+05 380 E+03 850 E+03 757 E+03 220 E+02

J/kg K J/kg K

4.186 800 E+03 4.186 800 E+03

Specific heat Btu/lb F cal/g C tonf/in. (tsi) kgf/mm2 ksi lbf/in.2 (psi) MN/m2

5.192 1.730 1.442 4.184

m/m k m/m k

1.000 000 E+00 1.800 000 E+00

m/s m/s m/s m/s m/s mm/h

8.466 5.080 3.048 2.540 2.777 1.609

rad/s rad/s

1.047 164 E01 6.283 185 E+00

pa s m2/s m2/s mm2/s

1.000 1.000 9.290 6.451

000 E01 000 E04 304 E02 600 E+02

m3 m3 m3 m3

1.638 2.831 2.957 3.785

706 E05 685 E02 353 E05 412 E03

204 E+02 735 E+00 279 E01 000 E+02

Thermal expansion in./in. C in./in. F

ft/h ft/min ft/s in./s km/h mph

667 E05 000 E03 000 E01 000 E02 778 E01 344 E+00

Velocity of rotation rev/min (rpm) rev/s Viscosity poise stokes ft2/s in.2/s

951 E+01 650 E+00 757 E+00 757 E03 000 E+00

in.3 ft3 fluid oz gal (U.S. liquid)

MPa MPa MPa MPa MPa

1.378 9.806 6.894 6.894 1.000

C K

5/9 ( F32) 5/9

ft3/min ft3/s in.3/min

m3/s m3/s m3/s

4.719 474 E04 2.831 685 E02 2.731 177 E07

C

5/9

Wavelength ˚ A

nm

1.000 000 E01

Temperature F R

W/m K W/m K W/m K W/m K

Volume

Stress (force per unit area) 2

Btu in./s ft2 F Btu/ft h F Btu in./h ft2 F cal/cm s C

Velocity

Power density

Temperature interval F

To convert from

Thermal conductivity

Power

(a) kg · 103 = 1 metric ton or 1 megagram (Mg)

Volume per unit time


18/8/2008 4:26PM Plate # 0


pg 525


APPENDIX 2

Temperature Conversion Table Temperature Conversions The general argument of this conversion table was devised by Sauveur and Boylston. The middle columns of numbers (in boldface type) contain the temperature readings ( F or C) to be converted. When converting from degrees Fahrenheit to degrees Celsius, read the Celsius equivalent in the column headed C. When converting from Celsius to Fahrenheit, read the Fahrenheit equivalent in the column headed F. °F

°C

°F

°C

°F

°C

°F

°C

... ... ... ... ...

458 456 454 452 450

272.22 271.11 270.00 268.89 267.78

... ... ... ... ...

368 366 364 362 360

222.22 221.11 220.00 218.89 217.73

... ... ... 457.6 454.0

278 276 274 272 270

172.22 171.11 170.00 168.89 167.78

306.4 302.8 299.2 295.6 292.0

188 186 184 182 180

122.22 121.11 120.00 118.89 117.78

... ... ... ... ...

448 446 444 442 440

266.67 265.56 264.44 263.33 262.22

... ... ... ... ...

358 356 354 352 350

216.67 215.56 214.44 213.33 212.22

450.4 446.8 443.2 439.6 436.0

268 266 264 262 260

166.67 165.56 164.44 163.33 162.22

288.4 284.8 281.2 277.6 274.0

178 176 174 172 170

116.67 115.56 114.44 113.33 112.22

... ... ... ... ...

433 436 434 432 430

261.11 260.00 258.89 257.78 256.67

... ... ... ... ...

348 346 344 342 340

211.11 210.00 208.89 207.78 206.67

432.4 428.8 425.2 421.6 418.0

258 256 254 252 250

161.11 160.00 158.89 157.78 156.67

270.4 266.8 263.2 259.6 256.0

168 166 164 162 160

111.11 110.00 108.89 107.78 106.67

... ... ... ... ...

428 426 424 422 420

255.56 254.44 253.33 252.22 251.11

... ... ... ... ...

338 336 334 332 330

205.56 204.44 203.33 202.22 201.11

414.4 410.8 407.2 403.6 400.0

248 246 244 242 240

155.56 154.44 153.33 152.22 151.11

252.4 248.8 245.2 241.6 238.0

158 156 154 152 150

105.56 104.44 103.33 102.22 101.11

... ... ... ... ...

418 416 414 412 410

250.00 248.89 247.78 246.67 245.56

... ... ... ... ...

328 326 324 322 320

200.00 198.89 197.78 196.67 195.56

396.4 392.8 389.2 385.6 382.0

238 236 234 232 230

150.00 148.89 147.78 146.67 145.56

234.4 230.8 227.2 223.6 220.0

148 146 144 142 140

100.00 98.89 97.78 96.67 95.56

... ... ... ... ...

408 406 404 402 400

244.44 243.33 242.22 241.11 240.00

... ... ... ... ...

318 316 314 312 310

194.44 193.33 192.22 191.11 190.00

378.4 374.8 371.2 367.6 364.0

228 226 224 222 220

144.44 143.33 142.22 141.11 140.00

216.4 212.8 209.2 205.6 202.0

138 136 134 132 130

94.44 93.33 92.22 91.11 90.00

... ... ... ... ...

398 396 394 392 390

238.89 237.78 236.67 235.56 234.44

... ... ... ... ...

308 306 304 302 300

188.89 187.78 186.67 185.56 184.44

360.4 356.8 353.2 349.6 346.0

218 216 214 212 210

138.89 137.78 136.67 135.56 134.44

198.4 194.8 191.2 187.6 184.0

128 126 124 122 120

88.89 87.78 86.67 85.56 84.44

... ... ... ... ...

388 386 384 382 380

233.33 232.22 231.11 230.00 228.89

... ... ... ... ...

298 296 294 292 290

183.33 182.22 181.11 180.00 178.89

342.4 338.8 335.2 331.6 328.0

208 206 204 202 200

133.33 132.22 131.11 130.00 128.89

180.4 176.8 173.2 169.6 166.0

113 116 114 112 110

83.33 82.22 81.11 80.00 78.89

... ... ... ... ...

378 376 374 372 370

227.78 226.67 225.56 224.44 223.33

... ... ... ... ...

288 286 284 282 280

177.78 176.67 175.56 174.44 173.33

324.4 320.8 317.2 313.6 310.0

198 196 194 192 190

127.78 126.67 125.56 124.44 123.33

162.4 158.8 155.2 151.6 148.0

108 106 104 102 100

77.78 76.67 75.56 74.44 73.33

(continued)


18/8/2008 4:27PM Plate # 0

pg 526


°C

°F

144.4 140.8 137.2 133.6 130.0

°F

98 96 94 92 90

°C

72.22 71.11 70.00 68.89 67.78

+89.6 +93.2 +96.8 +100.4 +104.0

°F

+32 +34 +36 +38 +40

+0.00 +1.11 +2.22 +3.33 +4.44

°C

323.6 327.2 330.8 334.4 338.0

°F

162 164 166 168 170

72.22 73.33 74.44 75.66 76.67

557.6 561.2 564.8 568.4 572.0

292 294 296 298 300

144.44 145.56 146.67 147.78 148.89

126.4 122.8 119.2 115.6 112.0

88 86 84 82 80

66.67 65.56 64.44 63.33 62.22

+107.6 +111.2 +114.8 +118.4 +122.0

+42 +44 +46 +48 +50

+5.56 +6.67 +7.78 +8.89 +10.00

341.6 345.2 348.8 352.4 356.0

172 174 176 178 180

77.78 78.89 80.00 81.11 82.22

575.6 579.2 582.8 586.4 590.0

302 304 306 308 310

150.00 151.11 152.22 153.33 154.44

108.4 104.8 101.2 97.6 94.0

78 76 74 72 70

61.11 60.00 58.89 57.78 56.67

+125.6 +129.2 +132.8 +136.4 +140.0

+52 +54 +56 +58 +60

+11.11 +12.22 +13.33 +14.44 +15.56

359.6 363.2 366.8 370.4 374.0

182 184 186 188 190

83.33 84.44 85.56 86.67 87.78

593.6 597.2 600.8 604.4 608.0

312 314 316 318 320

155.56 156.67 157.78 158.89 160.00

90.4 86.8 83.2 79.6 76.0

68 66 64 62 60

55.56 54.44 53.33 52.22 51.11

143.6 147.2 150.8 154.4 158.0

62 64 66 68 70

16.67 17.78 18.89 20.00 21.11

377.6 381.2 384.8 388.4 392.0

192 194 196 198 200

88.89 90.00 91.11 92.22 93.33

611.6 615.2 618.8 622.4 626.0

322 324 326 328 330

161.11 162.22 163.33 164.44 165.56

72.4 68.8 65.2 61.6 58.0

58 56 54 52 50

50.00 48.89 47.78 46.67 45.56

161.6 165.2 168.8 172.4 176.0

72 74 76 78 80

22.22 23.33 24.44 25.56 26.67

395.6 399.2 402.8 406.4 410.0

202 204 206 208 210

94.44 95.56 96.67 97.73 98.89

629.6 633.2 636.8 640.4 644.0

332 334 336 338 340

166.67 167.78 168.89 170.00 171.11

54.4 50.8 47.2 43.6 40.0

48 46 44 42 40

44.44 43.33 42.22 41.11 40.00

179.6 183.2 186.8 190.4 194.0

82 84 86 88 90

27.78 28.89 30.00 31.11 32.22

413.6 417.2 420.8 424.4 428.0

212 214 216 218 220

100.00 101.11 102.22 103.33 104.44

647.6 651.2 654.8 658.4 662.0

342 344 346 348 350

172.22 173.33 174.44 175.56 176.67

36.4 32.8 29.2 25.6 22.0

38 36 34 32 30

38.89 37.78 36.67 35.56 34.44

197.6 201.2 204.8 208.4 212.0

92 94 96 98 100

33.33 34.44 35.56 36.67 37.78

431.6 435.2 438.8 442.4 446.0

222 224 226 228 230

105.56 106.67 107.78 108.89 110.00

665.6 669.2 672.8 676.4 680.0

352 354 356 358 360

177.78 178.89 180.00 181.11 182.22

18.4 14.8 11.2 7.6 4.0

28 26 24 22 20

33.33 32.22 31.11 30.00 28.89

215.6 219.2 222.8 226.4 230.0

102 104 106 108 110

38.89 40.00 41.11 42.22 43.33

449.6 453.2 456.8 460.4 464.0

232 234 236 238 240

111.11 112.22 113.33 114.44 115.56

683.6 687.2 690.8 694.4 698.0

362 364 366 368 370

183.33 184.44 185.56 186.67 187.78

0.4 +3.2 +6.8 +10.4 +14.0

18 16 14 12 10

27.78 26.67 25.56 24.44 23.33

233.6 237.2 240.8 244.4 248.0

112 114 116 118 120

44.44 45.56 46.67 47.78 48.89

467.6 471.2 474.8 478.4 482.0

242 244 246 248 250

116.67 117.78 118.89 120.00 121.11

701.6 705.2 708.8 712.4 716.0

372 374 376 378 380

188.89 190.00 191.11 192.22 193.33

+17.6 +21.2 +24.8 +28.4 +32.0

8 6 4 2 +0

22.22 21.11 20.00 18.89 17.78

251.6 255.2 258.8 262.4 266.0

122 124 126 128 130

50.00 51.11 52.22 53.33 54.44

485.6 489.2 492.8 496.4 500.0

252 254 256 258 260

122.22 123.33 124.44 125.56 126.67

719.6 723.2 726.8 730.4 734.0

382 384 386 388 390

194.44 195.56 196.67 197.78 198.89

+35.6 +39.2 +42.8 +46.4 +50.0

+2 +4 +6 +8 +10

16.67 15.56 14.44 13.33 12.22

269.6 273.2 276.8 280.4 284.0

132 134 136 138 140

55.56 56.67 57.78 58.89 60.00

503.6 507.2 510.8 514.4 518.0

262 264 266 268 270

127.78 128.89 130.00 131.11 132.22

737.6 741.2 744.8 748.4 752.0

392 394 396 398 400

200.00 201.11 202.22 203.33 204.44

+53.6 +57.2 +60.8 +64.4 +68.0

+12 +14 +16 +18 +20

11.11 10.00 8.89 7.78 6.67

287.6 291.2 294.8 298.4 302.0

142 144 146 148 150

61.11 62.22 63.33 64.44 65.56

521.6 525.2 528.8 532.4 536.0

272 274 276 278 280

133.33 134.44 135.56 136.67 137.78

755.6 759.2 762.8 766.4 770.0

402 404 406 408 410

205.56 206.67 207.78 208.89 210.00

+71.6 +75.2 +78.8 +82.4 +86.0

+22 +24 +26 +28 +30

5.56 4.44 3.33 2.22 1.11

305.6 309.2 312.8 316.4 320.0

152 154 156 158 160

66.67 67.73 68.83 70.00 71.11

539.6 543.2 546.8 550.4 554.0

282 284 286 288 290

138.89 140.00 141.11 142.22 143.33

773.6 777.2 780.8 784.4 788.0

412 414 416 418 420

211.11 212.22 213.33 214.44 215.56

(continued)

°C


18/8/2008 4:27PM Plate # 0

Appendix 2:

°F

°C

°F

°C

pg 527

Temperature Conversion Table / 527

°F

°C

°F

°C

791.6 795.2 798.8 802.4 806.0

422 424 426 428 430

216.67 217.78 218.89 220.00 221.11

1040.0 1058.0 1076.0 1094.0 1112.0

560 570 580 590 600

293.33 298.89 304.44 310.00 315.56

2210.0 2228.0 2246.0 2264.0 2282.0

1210 1220 1230 1240 1250

654.44 660.00 665.56 671.11 676.67

3380.0 3398.0 3416.0 3434.0 3452.0

1860 1870 1880 1890 1900

1015.6 1021.1 1026.7 1032.2 1037.8

809.6 813.2 816.8 820.4 824.0

432 434 436 438 440

222.22 223.33 224.44 225.56 226.67

1130.0 1148.0 1166.0 1184.0 1202.0

610 620 630 640 650

321.11 326.67 332.22 337.78 343.33

2300.0 2318.0 2336.0 2354.0 2372.0

1260 1270 1280 1290 1300

682.22 687.78 693.33 698.89 704.44

3470.0 3488.0 3506.0 3524.0 3542.0

1910 1920 1930 1940 1950

1043.3 1048.9 1054.4 1060.0 1065.6

827.6 831.2 834.8 838.4 842.0

442 444 446 448 450

227.78 228.89 230.00 231.11 232.22

1220.0 1238.0 1256.0 1274.0 1292.0

660 670 680 690 700

348.89 354.44 360.00 365.56 371.11

2390.0 2408.0 2426.0 2444.0 2462.0

1310 1320 1330 1340 1350

710.00 715.56 721.11 726.67 732.22

3560.0 3578.0 3596.0 3614.0 3632.0

1960 1970 1980 1990 2000

1071.1 1076.7 1082.2 1087.8 1093.3

845.6 849.2 852.8 856.4 860.0

452 454 456 458 460

233.33 234.44 235.56 236.67 237.78

1310.0 1328.0 1346.0 1364.0 1382.0

710 720 730 740 750

376.67 382.22 387.78 393.33 398.89

2480.0 2498.0 2516.0 2534.0 2552.0

1360 1370 1380 1390 1400

737.78 743.33 748.89 754.44 760.00

3650.0 3668.0 3686.0 3704.0 3722.0

2010 2020 2030 2040 2050

1098.9 1104.4 1110.0 1115.6 1121.1

863.6 867.2 870.8 874.4 878.0

462 464 466 468 470

238.89 240.00 241.11 242.22 243.33

1400.0 1418.0 1436.0 1454.0 1472.0

760 770 780 790 800

404.44 410.00 415.56 421.11 426.67

2570.0 2588.0 2606.0 2624.0 2642.0

1410 1420 1430 1440 1450

765.56 771.11 776.67 782.22 787.78

3740.0 3758.0 3776.0 3794.0 3812.0

2060 2070 2080 2090 2100

1126.7 1132.2 1137.8 1143.3 1148.9

881.6 885.2 888.8 892.4 896.0

472 474 476 478 480

244.44 245.56 246.67 247.78 248.89

1490.0 1508.0 1526.0 1544.0 1562.0

810 820 830 840 850

432.22 437.78 443.33 448.89 454.44

2660.0 2678.0 2696.0 2714.0 2732.0

1460 1470 1480 1490 1500

793.33 798.89 804.44 810.00 815.56

3830.0 3848.0 3866.0 3884.0 3902.0

2110 2120 2130 2140 2150

1154.4 1160.0 1165.6 1171.1 1176.7

899.6 903.2 906.8 910.4 914.0

482 484 486 488 490

250.00 251.11 252.22 253.33 254.44

1580.0 1598.0 1616.0 1634.0 1652.0

860 870 880 890 900

460.00 465.56 471.11 476.67 482.22

2750.0 2768.0 2786.0 2804.0 2822.0

1510 1520 1530 1540 1550

821.11 826.67 832.22 837.78 843.33

3920.0 3938.0 3956.0 3974.0 3992.0

2160 2170 2180 2190 2200

1182.2 1187.8 1193.3 1198.9 1204.4

917.6 921.2 924.8 928.4 932.0

492 494 496 498 500

255.56 256.67 257.78 258.89 260.00

1670.0 1688.0 1706.0 1724.0 1742.0

910 920 930 940 950

487.78 493.33 498.89 504.44 510.00

2840.0 2858.0 2876.0 2894.0 2912.0

1560 1570 1580 1590 1600

848.89 854.44 860.00 865.56 871.11

4010.0 4028.0 4046.0 4064.0 4082.0

2210 2220 2230 2240 2250

1210.0 1215.6 1221.1 1226.7 1232.2

935.6 939.2 942.8 946.4 950.0

502 504 506 508 510

261.11 262.22 263.33 264.44 265.56

1760.0 1778.0 1796.0 1814.0 1832.0

960 970 980 990 1000

515.56 521.11 526.67 532.22 537.78

2930.0 2948.0 2966.0 2984.0 3002.0

1610 1620 1630 1640 1650

876.67 882.22 887.78 893.33 898.89

4100.0 4118.0 4136.0 4154.0 4172.0

2260 2270 2280 2290 2300

1237.8 1243.3 1248.9 1254.4 1260.0

953.6 957.2 960.8 964.4 968.0

512 514 516 518 520

266.67 267.78 268.89 270.00 271.11

1850.0 1868.0 1886.0 1904.0 1922.0

1010 1020 1030 1040 1050

543.33 548.89 554.44 560.00 565.56

3020.0 3038.0 3056.0 3074.0 3092.0

1660 1670 1680 1690 1700

904.44 910.00 915.56 921.11 926.67

4190.0 4208.0 4226.0 4244.0 4262.0

2310 2320 2330 2340 2350

1265.6 1271.1 1276.7 1282.2 1287.8

971.6 975.2 978.8 982.4 986.0

522 524 526 528 530

272.22 273.33 274.44 275.56 276.67

1940.0 1958.0 1976.0 1994.0 2012.0

1060 1070 1080 1090 1100

571.11 576.67 582.22 587.78 593.33

3110.0 3128.0 3146.0 3164.0 3182.0

1710 1720 1730 1740 1750

932.22 937.78 943.33 948.89 954.44

4280.0 4298.0 4316.0 4334.0 4352.0

2360 2370 2380 2390 2400

1293.3 1298.9 1304.4 1310.1 1315.6

989.6 993.2 996.8 1000.4 1004.0

532 534 536 538 540

277.78 278.89 280.00 281.11 282.22

2030.0 2048.0 2066.0 2084.0 2102.0

1100 1120 1130 1140 1150

598.89 604.44 610.00 615.56 621.11

3200.0 3218.0 3236.0 3254.0 3272.0

1760 1770 1780 1790 1800

960.00 965.56 971.11 976.67 982.22

4370.0 4388.0 4406.6 4424.0 4442.0

2410 2420 2430 2440 2450

1321.1 1326.7 1332.2 1337.8 1343.3

1007.6 1011.2 1014.8 1018.4 1022.0

542 544 546 548 550

283.22 284.44 285.56 286.67 287.78

2120.0 2138.0 2156.0 2174.0 2192.0

1160 1170 1180 1190 1200

626.67 632.22 637.78 643.33 648.89

3290.0 3308.0 3326.0 3344.0 3362.0

1810 1820 1830 1840 1850

987.78 993.33 998.89 1004.4 1010.0

4460.0 4478.0 4496.0 4514.0 4532.0

2460 2470 2480 2490 2500

1348.9 1354.4 1360.0 1365.6 1371.1

(continued)


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pg 528


°F

°C

°F

°C

°F

°C

°F

°C

4550.0 4568.0 4586.0 4604.0 4622.0

2510 2520 2530 2540 2550

1376.7 1382.2 1387.8 1393.3 1398.9

5090.0 5108.0 5126.0 5144.0 5162.0

2810 2820 2830 2840 2850

1543.3 1548.9 1554.4 1560.0 1565.6

5702.0 5792.0 5882.0 5972.0 6062.0

3150 3200 3250 3300 3350

1732.2 1760.0 1787.7 1815.5 1843.3

8402.0 8492.0 8582.0 8672.0 8762.0

4650 4700 4750 4800 4850

2565.5 2593.3 2621.1 2648.8 2676.6

4640.0 4658.0 4676.0 4694.0 4712.0

2560 2570 2580 2590 2600

1404.4 1410.0 1415.6 1421.1 1426.7

5180.0 5198.0 5216.0 5234.0 5252.0

2860 2870 2880 2890 2900

1571.1 1576.7 1582.2 1587.8 1593.3

6152.0 6242.0 6332.0 6422.0 6512.0

3400 3450 3500 3550 3600

1871.1 1898.8 1926.6 1954.4 1982.2

8852.0 8942.0 9032.0 9122.0 9212.0

4900 4950 5000 5050 5100

2704.4 2732.2 2760.0 2787.7 2815.5

4730.0 4748.0 4766.0 4784.0 4802.0

2610 2620 2630 2640 2650

1432.2 1437.8 1443.3 1448.9 1454.4

5270.0 5288.0 5306.0 5324.0 5342.0

2910 2920 2930 2940 2950

1598.9 1604.4 1610.0 1615.6 1621.1

6602.0 6692.0 6782.0 6872.0 6962.0

3650 3700 3750 3800 3850

2010.0 2037.7 2065.5 2093.3 2121.1

9302.0 9392.0 9482.0 9572.0 9662.0

5150 5200 5250 5300 5350

2843.3 2871.1 2898.8 2926.6 2954.4

4820.0 4838.0 4856.0 4874.0 4892.0

2660 2670 2680 2690 2700

1460.0 1465.6 1471.1 1476.7 1482.2

5360.0 5378.0 5396.0 5414.0 5432.0

2960 2970 2980 2990 3000

1626.7 1632.2 1637.8 1643.3 1648.9

7052.0 7142.0 7232.0 7322.0 7412.0

3900 3950 4000 4050 4100

2148.8 2176.6 2204.4 2232.2 2260.0

9752.0 9842.0 9932.0 10022.0 10112.0

5400 5450 5500 5550 5600

2982.2 3010.0 3037.7 3065.5 3093.3

4910.0 4928.0 4946.0 4964.0 4982.0

2710 2720 2730 2740 2750

1487.8 1493.3 1498.9 1504.4 1510.0

5450.0 5468.0 5486.0 5504.0 5522.0

3010 3020 3030 3040 3050

1654.4 1660.0 1665.5 1671.1 1676.7

7502.0 7592.0 7682.0 7772.0 7862.0

4150 4200 4250 4300 4350

2287.7 2315.5 2343.3 2371.1 2398.8

5000.0 5018.0 5036.0 5054.0 5072.0

2760 2770 2780 2790 2800

1515.6 1521.1 1526.7 1532.2 1537.8

5540.0 5558.0 5576.0 5594.0 5612.0

3060 3070 3080 3090 3100

1682.2 1687.8 1693.3 1698.9 1704.4

7952.0 8042.0 8132.0 8222.0 8312.7

4400 4450 4500 4550 4600

2426.6 2454.4 2482.2 2510.0 2537.7


18/8/2008 4:28PM Plate # 0



APPENDIX 3

Steel Hardness Conversions FROM A PRACTICAL STANDPOINT, it is important to be able to convert the results of one type of hardness test into those of a different test. Because a hardness test does not measure a well-defined property of a material and because all the tests in common use are not based on the same type of measurements, it is not surprising that universal hardness conversion relationships have not been developed. Hardness conversions instead are empirical relationships that are defined by conversion tables limited to specific categories of materials. That is, different conversion tables are required for materials with greatly different elastic moduli or with different strain-hardening capacity. The most reliable hardness-conversion data exist for steel that is harder than 240 HB. The indentation hardness of soft metals depends on the strain-hardening behavior of the material during the test, which in turn depends on the previous degree of strain hardening of the material before the test. The modulus of elasticity also has been shown to influence conversions at high hardness levels. At low hardness levels, conversions between hardness scales measuring depth and those measuring diameter are likewise influenced by differences in the modulus of elasticity. Hardness conversions are covered in standards such as SAE J417, “Hardness Tests and Hardness Conversions”; ISO 4964, “Hardness Conversions—Steel”; and ASTM E140, “Standard Hardness Conversion Tables for Metals.” Conversion tables for nickel and high-nickel alloys, cartridge brass, austenitic stainless steel plate and sheet, and copper can be found in ASTM E140. Recently, ASTM committee E-28 on indentation hardness has developed mathematical conversion formulas based on the conversion-table values fround in ASTM E140. Over 60 conversion formulas are listed in the appendix of ASTM E140, and these formulas can be used in place of the tables. A computer is helpful in performing the calculations quickly. Other hardness conversion formulas for various materials have also been published, and a list of some other conversion formulas is given in Table A3.1. The standard procedure for

Table A3.1 Examples of published hardness conversion equations Steels

HB=

7300 1307HRB

(40–100 HRB)

HB=

3710 1307HRE

(30–100 HRE)

HB=

1,520,00074500 HRC (1007HRC)2

(540 HRC)

25,000710(577HRC)2 1007HRC 6700 HRB=1347 HB 1=2 2:43 · 106 HRC=119:07 HV 1=2 6:85 · 105 HRA=112:37 HV 1=2 5:53 · 105 HR15N=117:947 HV 1=2 1:88 · 106 HR30N=129:527 HV 1=2 3:132 · 106 HR45N=133:517 HV HB=

HB = 0.951 HV HB = 0.941 HV

(40–70 HRC) (+7 HRB, 95% CL) (240–1040 HV) (240–1040 HV) (240–1040 HV) (240–1040 HV) (240–1040 HV) (steel ball, 200–400 HV) (tungsten-carbide ball, 200–700 HV)

Cemented carbides

1=2 2:43 · 106 HRC=117:357 HV 1=2 2:437106 2117 HV HRA= 1:885

(900–1800 HV)

(900–1800 HV)

Rockwell from Knoop for steels

HRC = 64.934 log HK HRC = 67.353 log HK HRC = 71.983 log HK HRC = 76.572 log HK HRC = 79.758 log HK HRC = 82.283 log HK HRC = 83.58 log HK HRC = 85.848 log HK

140.38 144.32 154.28 163.89 170.92 176.92 179.30 184.55

(15 gf) (25 gf) (50 gf) (100 gf) (200 gf) (300 gf) (500 gf) (1000 gf)

White cast irons

HB = 0.363 (HRC)2 22.515 (HRC)+717.8 HV = 0.343 (HRC)2 18.132 (HRC)+595.3 HV = 1.136 (HB)2 26.0 Austenitic stainless steel

1 =0:0001304(1307HRB) HB Stable alpha-beta titanium alloys

HRC = 0.078 HV+8.1

(60–90 HRB, 110–192 HB)

pg 529


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from ASTM E140, are for conversion among Rockwell, Brinell, and Vickers hardness for heat treated carbon and alloy steels, almost all constructional alloy steels, and tool steels in the asforged, annealed, normalized, and quenched and tempered conditions. The tables are also summarized in graphical form in Fig. A3.1.

reporting converted hardness numbers indicates the measured hardness and test scale in parentheses—for example, 451 HB (48 HRC). The method of conversion (table, formula, or other method) should also be defined. When making hardness correlations, it is best to consult ASTM E140. Tables A3.2 to A3.5,

Table A3.2 Approximate Rockwell B hardness conversion numbers for nonaustenitic steels Rockwell B, 100 kgf, 1/16 in. ball

100 99 98 97 96 95 94 93 92 91 90 89 88 87 86 85 84 83 82 81 80 79 78 77 76 75 74 73 72 71 70 69 68 67 66 65 64 63 62 61 60 59 58 57 56 55 54 53 52 51 50 49 48 47 46 45 44

A, 60 kgf, diamond

61.5 60.9 60.2 59.5 58.9 58.3 57.6 57.0 56.4 55.8 55.2 54.6 54.0 53.4 52.8 52.3 51.7 51.1 50.6 50.0 49.5 48.9 48.4 47.9 47.3 46.8 46.3 45.8 45.3 44.8 44.3 43.8 43.3 42.8 42.3 41.8 41.4 40.9 40.4 40.0 39.5 39.0 38.6 38.1 37.7 37.2 36.8 36.3 35.9 35.5 35.0 34.6 34.1 33.7 33.3 32.9 32.4

Superficial Rockwell E, 100 kgf, 15T, 15 kgf, 1/8 in. ball 1/16 in. ball

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 100 99.5 99.0 98.0 97.5 97.0 96.0 95.5 95.0 94.5 93.5 93.0 92.5 92.0 91.0 90.5 90.0 89.5 89.0 88.0 87.5 87.0 86.5 85.5 85.0 84.5 84.0 83.5

93.1 92.8 92.5 92.1 91.8 91.5 91.2 90.8 90.5 90.2 89.9 89.5 89.2 88.9 88.6 88.2 87.9 87.6 87.3 86.9 86.6 86.3 86.0 85.6 85.3 85.0 84.7 84.3 84.0 83.7 83.4 83.0 82.7 82.4 82.1 81.8 81.4 81.1 80.8 80.5 80.1 79.8 79.5 79.2 78.8 78.5 78.2 77.9 77.5 77.2 76.9 76.6 76.2 75.9 75.6 75.3 74.9

30T, 30 kgf, 1/16 in. ball

83.1 82.5 81.8 81.1 80.4 79.8 79.1 78.4 77.8 77.1 76.4 75.8 75.1 74.4 73.8 73.1 72.4 71.8 71.1 70.4 69.7 69.1 68.4 67.7 67.1 66.4 65.7 65.1 64.4 63.7 63.1 62.4 61.7 61.0 60.4 59.7 59.0 58.4 57.7 57.0 56.4 55.7 55.0 54.4 53.7 53.0 52.4 51.7 51.0 50.3 49.7 49.0 48.3 47.7 47.0 46.3 45.7

45T, 45 kgf, 1/16 in. ball Vickers

72.9 71.9 70.9 69.9 68.9 67.9 66.9 65.9 64.8 63.8 62.8 61.8 60.8 59.8 58.8 57.8 56.8 55.8 54.8 53.8 52.8 51.8 50.8 49.8 48.8 47.8 46.8 45.8 44.8 43.8 42.8 41.8 40.8 39.8 38.7 37.7 36.7 35.7 34.7 33.7 32.7 31.7 30.7 29.7 28.7 27.7 26.7 25.7 24.7 23.7 22.7 21.7 20.7 19.7 18.7 17.7 16.7

240 234 228 222 216 210 205 200 195 190 185 180 176 172 169 165 162 159 156 153 150 147 144 141 139 137 135 132 130 127 125 123 121 119 117 116 114 112 110 108 107 106 104 103 101 100 ... ... ... ... ... ... ... ... ... ... ...

Knoop, 500 gf and over

Brinell, 3000 kgf, 10 mm ball

Tensile strength MPa (ksi)


251 246 241 236 231 226 221 216 211 206 201 196 192 188 184 180 176 173 170 167 164 161 158 155 152 150 147 145 143 141 139 137 135 133 131 129 127 125 124 122 120 118 117 115 114 112 111 110 109 108 107 106 105 104 103 102 101

240 234 228 222 216 210 205 200 195 190 185 180 176 172 169 165 162 159 156 153 150 147 144 141 139 137 135 132 130 127 125 123 121 119 117 116 114 112 110 108 107 106 104 103 101 100 ... ... ... ... ... ... ... ... ... ... ...

800 (116) 787 (114) 752 (109) 724 (105) 704 (102) 690 (100) 676 (98) 648 (94) 634 (92) 620 (90) 614 (89) 607 (88) 593 (86) 579 (84) 572 (83) 565 (82) 558 (81) 552 (80) 524 (76) 503 (73) 496 (72) 482 (70) 475 (69) 469 (68) 462 (67) 455 (66) 448 (65) 441 (64) 434 (63) 427 (62) 421 (61) 414 (60) 407 (59) 400 (58) 393 (57) 386 (56) ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

201 195 189 184 179 175 171 167 163 160 157 154 151 148 145 142 140 137 135 133 130 128 126 124 122 120 118 116 114 112 110 109 108 106 104 102 100 99 98 96 95 94 92 91 90 89 87 86 85 84 83 82 81 80 80 79 78

Data are only approximate conversions for carbon and low-alloy steels in the annealed, normalized, and quenched-and-tempered conditions; less accurate for cold-worked condition and for austenitic steels. Source: ASTM E140, except for values for E scale and tensile strength, which are not from standards


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Appendix 3:

pg 531

Steel Hardness Conversions / 531

Table A3.3 Approximate Rockwell C hardness conversion numbers for nonaustenitic steels, according to ASTM E140 C, 150 kgf, diamond

68 67 66 65 64 63 62 61 60 59 58 57 56 55 54 53 52 51 50 49 48 47 46 45 44 43 42 41 40 39 38 37 36 35 34 33 32 31 30 29 28 27 26 25 24 23 22 21 20

A, 60 kgf, diamond

D, 100 kgf, diamond

15 N, 15 kgf, diamond



Vickers

Knoop, 500 gf and over


Tensile strength, MPa (ksi)

85.6 85.0 84.5 83.9 83.4 82.8 82.3 81.8 81.2 80.7 80.1 79.6 79.0 78.5 78.0 77.4 76.8 76.3 75.9 75.2 74.7 74.1 73.6 73.1 72.5 72.0 71.5 70.9 70.4 69.9 69.4 68.9 68.4 67.9 67.4 66.8 66.3 65.8 65.3 64.8 64.3 63.8 63.3 62.8 62.4 62.0 61.5 61.0 60.5

76.9 76.1 75.4 74.5 73.8 73.0 72.2 71.5 70.7 69.9 69.2 68.5 67.7 66.9 66.1 65.4 64.6 63.8 63.1 62.1 61.4 60.8 60.0 59.2 58.5 57.7 56.9 56.2 55.4 54.6 53.8 53.1 52.3 51.5 50.8 50.0 49.2 48.4 47.7 47.0 46.1 45.2 44.6 43.8 43.1 42.1 41.6 40.9 40.1

93.2 92.9 92.5 92.2 91.8 91.4 91.1 90.7 90.2 89.8 89.3 88.9 88.3 87.9 87.4 86.9 86.4 85.9 85.5 85.0 84.5 83.9 83.5 83.0 82.5 82.0 81.5 80.9 80.4 79.9 79.4 78.8 78.3 77.7 77.2 76.6 76.1 75.6 75.0 74.5 73.9 73.3 72.8 72.2 71.6 71.0 70.5 69.9 69.4

84.4 83.6 82.8 81.9 81.1 80.1 79.3 78.4 77.5 76.6 75.7 74.8 73.9 73.0 72.0 71.2 70.2 69.4 68.5 67.6 66.7 65.8 64.8 64.0 63.1 62.2 61.3 60.4 59.5 58.6 57.7 56.8 55.9 55.0 54.2 53.3 52.1 51.3 50.4 49.5 48.6 47.7 46.8 45.9 45.0 44.0 43.2 42.3 41.5

75.4 74.2 73.3 72.0 71.0 69.9 68.8 67.7 66.6 65.5 64.3 63.2 62.0 60.9 59.8 58.6 57.4 56.1 55.0 53.8 52.5 51.4 50.3 49.0 47.8 46.7 45.5 44.3 43.1 41.9 40.8 39.6 38.4 37.2 36.1 34.9 33.7 32.5 31.3 30.1 28.9 27.8 26.7 25.5 24.3 23.1 22.0 20.7 19.6

940 900 865 832 800 772 746 720 697 674 653 633 613 595 577 560 544 528 513 498 484 471 458 446 434 423 412 402 392 382 372 363 354 345 336 327 318 310 302 294 286 279 272 266 260 254 248 243 238

920 895 870 846 822 799 776 754 732 710 690 670 650 630 612 594 576 558 542 526 510 495 480 466 452 438 426 414 402 391 380 370 360 351 342 334 326 318 311 304 297 290 284 278 272 266 261 256 251

... ... ... 739(a) 722(a) 705(a) 688(a) 670(a) 654(a) 634(a) 615 595 577 560 543 525 512 496 481 469 455 443 432 421 409 400 390 381 371 362 353 344 336 327 319 311 301 294 286 279 271 264 258 253 247 243 237 231 226

... ... ... ... ... ... ... ... ... 2420 (351) 2330 (338) 2240 (325) 2158 (313) 2075 (301) 2013 (292) 1951 (283) 1882 (273) 1820 (264) 1758 (255) 1696 (246) 1634 (237) 1579 (229) 1524 (221) 1482 (215) 1434 (208) 1386 (201) 1344 (195) 1296 (188) 1254 (182) 1220 (177) 1179 (171) 1137 (166) 1110 (161) 1075 (156) 1048 (152) 1027 (149) 1006 (146) 972 (141) 951 (138) 930 (135) 903 (131) 882 (128) 861 (125) 848 (123) 820 (119) 806 (117) 792 (115) 772 (112) 758 (110)

Data are only approximate conversions for carbon and low-alloy steels in the annealed, normalized, and quenched-and-tempered conditions; less accurate for cold-worked condition and for austenitic steels. (a) Hardness values outside the recommended range for Brinell testing per ASTM E10. Source: ASTM E140, except for values for tensile strength, which are not from standards


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pg 532


Table A3.4 Approximate equivalent hardness numbers for Brinell hardness numbers for steel Brinell hardness number(a) 3000 kgf load, 10 mm ball(a) Brinell indentation diam, mm

2.25 2.30 2.35 2.40 2.45 2.50 2.55 2.60 2.65 2.70 2.75 2.80 2.85 2.90 2.95 3.00 3.05 3.10 3.15 3.20 3.25 3.30 3.35 3.40 3.45 3.50 3.55 3.60 3.65 3.70 3.75 3.80 3.85 3.90 3.95 4.00 4.05 4.10 4.15 4.20 4.25 4.30 4.35 4.40 4.45 4.50 4.55 4.60 4.65 4.70 4.75 4.80

Standard ball

Tungstencarbide ball

... ... ... ... ... ... ... ... ... ... (495) ... (477) ... (461) ... 444 ... 429 415 401 388 375 363 352 341 331 321 311 302 293 285 277 269 262 255 248 241 235 229 223 217 212 207 201 197 192 187 183 179 174 170 167 163 159 156

(745) (712) (682) (653) 627 601 578 555 534 514 ... 495 ... 477 ... 461 ... 444 429 415 401 388 375 363 352 341 331 321 311 302 293 285 277 269 262 255 248 241 235 229 223 217 212 207 201 197 192 187 183 179 174 170 167 163 159 156

Rockwell hardness No. C scale, D scale, A scale, B scale, 150 kgf 100 kgf 60 kgf 100 kgf Vickers load, load, load, load, 1 /16 in. hardness diamond diamond diamond No. indenter diam ball indenter indenter

840 783 737 697 667 640 615 591 569 547 539 528 516 508 495 491 474 472 455 440 425 410 396 383 372 360 350 339 328 319 309 301 292 284 276 269 261 253 247 241 234 228 222 218 212 207 202 196 192 188 182 178 175 171 167 163

84.1 83.1 82.2 81.2 80.5 79.8 79.1 78.4 77.8 76.9 76.7 76.3 75.9 75.6 75.1 74.9 74.3 74.2 73.4 72.8 72.0 71.4 70.6 70.0 69.3 68.7 68.1 67.5 66.9 66.3 65.7 65.3 64.6 64.1 63.6 63.0 62.5 61.8 61.4 60.8 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... (110.0) (109.0) (108.5) (108.0) (107.5) (107.0) (106.0) (105.5) (104.5) (104.0) (103.0) (102.0) (101.0) 100.0 99.0 98.2 97.3 96.4 95.5 94.6 93.7 92.8 91.9 90.9 90.0 89.0 88.0 87.0 86.0 85.0 83.9 82.9

65.3 63.4 61.7 60.0 58.7 57.3 56.0 54.7 53.5 52.1 51.6 51.0 50.3 49.6 48.8 48.5 47.2 47.1 45.7 44.5 43.1 41.8 40.4 39.1 37.9 36.6 35.5 34.3 33.1 32.1 30.9 29.9 28.8 27.6 26.6 25.4 24.2 22.8 21.7 20.5 (19.0) (17.7) (16.4) (15.2) (13.8) (12.7) (11.5) (10.2) (9.0) (8.0) (6.7) (5.4) (4.4) (3.3) (2.0) (0.9)

74.8 73.4 72.0 70.7 69.7 68.7 67.7 66.7 65.8 64.7 64.3 63.8 63.2 62.7 61.9 61.7 61.0 60.8 59.7 58.8 57.8 56.8 55.7 54.6 53.8 52.8 51.9 51.0 50.0 49.3 48.3 47.6 46.7 45.9 45.0 44.2 43.2 42.0 41.4 40.5 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

Rockwell superficial hardness No., diamond indenter Knoop 15 N 30 N 45 N scale, scale, scale, hardness No., Scleroscope 15 kgf 30 kgf 45 kgf 500 gf load and greater hardness No. load load load

92.3 91.6 91.0 90.2 89.6 89.0 88.4 87.8 87.2 86.5 86.3 85.9 85.6 85.3 84.9 84.7 84.1 84.0 83.4 82.8 82.0 81.4 80.6 80.0 79.3 78.6 78.0 77.3 76.7 76.1 75.5 75.0 74.4 73.7 73.1 72.5 71.7 70.9 70.3 69.7 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

82.2 80.5 79.0 77.5 76.3 75.1 73.9 72.7 71.6 70.3 69.9 69.4 68.7 68.2 67.4 67.2 66.0 65.8 64.6 63.5 62.3 61.1 59.9 58.7 57.6 56.4 55.4 54.3 53.3 52.2 51.2 50.3 49.3 48.3 47.3 46.2 45.1 43.9 42.9 41.9 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

72.2 70.4 68.5 66.5 65.1 63.5 62.1 60.6 59.2 57.6 56.9 56.1 55.2 54.5 53.5 53.2 51.7 51.5 49.9 48.4 46.9 45.3 43.6 42.0 40.5 39.1 37.8 36.4 34.4 33.8 32.4 31.2 29.9 28.5 27.3 26.0 24.5 22.8 21.5 20.1 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

852 808 768 732 703 677 652 626 604 579 571 558 545 537 523 518 499 496 476 459 441 423 407 392 379 367 356 345 336 327 318 310 302 294 286 279 272 265 259 253 247 242 237 232 227 222 217 212 207 202 198 194 190 186 182 178

91 ... 84 81 79 77 75 73 71 70 ... 68 ... 66 ... 65 ... 63 61 59 58 56 54 52 51 50 48 47 46 45 43 42 41 40 39 38 37 36 35 34 ... 33 32 31 ... 30 29 ... 28 27 ... 26 ... 25 ... 24

(continued) Note: Values in parentheses are beyond normal range and are given for information only. Data are for carbon and alloy steels in the annealed, normalized, and quenchedand-tempered conditions; less accurate for cold-worked condition and for austenitic steels (a) Brinell numbers are based on the diameter of impressed indentation. If the ball distorts (flattens) during test, Brinell numbers will vary in accordance with the degree of such distortion when related to hardnesses determined with a Vickers diamond pyramid. Rockwell diamond indenter, or other indenter that does not sensibly distort. At high hardnesses, therefore, the relationship between Brinell and Vickers or Rockwell scales is affected by the type of ball used. Standard steel balls tend to flatten slightly moe than tungsten-carbide balls, resulting in a larger indentation and a lower Brinell number than shown by a tungsten carbide ball. Thus, on a specimen of about 539–547 HV, a standard ball will leave a 2.75 mm indentation (495 HB), and a tungsten carbide ball a 2.70 mm indentation (514 HB). Conversely, identical indentation diameters for both types of ball will correspond to different Vickers and Rockwell values. Thus, if indentation in two different specimens both are 2.75 mm diameter (495 HB), the specimen tested with a standard ball has a Vickers hardness of 539, whereas the specimen tested with a tungsten-carbide ball has a Vickers hardness of 528. Source: ASTM E140


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pg 533


Table A3.4 (continued) Brinell hardness number(a) 3000 kgf load, 10 mm ball(a) Brinell indentation diam, mm

4.85 4.90 4.95 5.00 5.10 5.20 5.30 5.40 5.50 5.60

Standard ball

Tungstencarbide ball

152 149 146 143 137 131 126 121 116 111

152 149 146 143 137 131 126 121 116 111

Rockwell hardness No. C scale, D scale, A scale, B scale, 150 kgf 100 kgf 60 kgf 100 kgf Vickers load, load, load, load, 1 /16 in. hardness diamond diamond diamond No. indenter diam ball indenter indenter

... ... ... ... ... ... ... ... ... ...

159 156 153 150 143 137 132 127 122 117

81.9 80.8 79.7 78.6 76.4 74.2 72.0 69.8 67.6 65.4

... ... ... ... ... ... ... ... ... ...

... ... ... ... ... ... ... ... ... ...

Rockwell superficial hardness No., diamond indenter Knoop 15 N 30 N 45 N scale, scale, scale, hardness No., Scleroscope 15 kgf 30 kgf 45 kgf 500 gf load and greater hardness No. load load load

... ... ... ... ... ... ... ... ... ...

... ... ... ... ... ... ... ... ... ...

... ... ... ... ... ... ... ... ... ...

... 23 ... 22 21 ... 20 19 18 17

174 170 166 163 157 151 145 140 135 131

Note: Values in parentheses are beyond normal range and are given for information only. Data are for carbon and alloy steels in the annealed, normalized, and quenchedand-tempered conditions; less accurate for cold-worked condition and for austenitic steels (a) Brinell numbers are based on the diameter of impressed indentation. If the ball distorts (flattens) during test, Brinell numbers will vary in accordance with the degree of such distortion when related to hardnesses determined with a Vickers diamond pyramid. Rockwell diamond indenter, or other indenter that does not sensibly distort. At high hardnesses, therefore, the relationship between Brinell and Vickers or Rockwell scales is affected by the type of ball used. Standard steel balls tend to flatten slightly moe than tungsten-carbide balls, resulting in a larger indentation and a lower Brinell number than shown by a tungsten carbide ball. Thus, on a specimen of about 539–547 HV, a standard ball will leave a 2.75 mm indentation (495 HB), and a tungsten carbide ball a 2.70 mm indentation (514 HB). Conversely, identical indentation diameters for both types of ball will correspond to different Vickers and Rockwell values. Thus, if indentation in two different specimens both are 2.75 mm diameter (495 HB), the specimen tested with a standard ball has a Vickers hardness of 539, whereas the specimen tested with a tungsten-carbide ball has a Vickers hardness of 528. Source: ASTM E140

Table A3.5 Approximate equivalent hardness numbers for Vickers (diamond pyramid) hardness numbers for steel

Vickers hardness No.

940 920 900 880 860 840 820 800 780 760 740 720 700 690 680 670 660 650 640 630 620 610 600 590 580 570

Standard ball

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

A scale, 60 kgf Tungstenload, carbide diamond ball indenter

... ... ... (767) (757) (745) (733) (722) (710) (698) (684) (670) (656) (647) (638) (630) 620 611 601 591 582 573 564 554 545 535

Rockwell superficial (diamond pyramid) hardness No., diamond indenter

Rockwell hardness No.

Brinell hardness No., 3000 kg load, 10 mm ball

85.6 85.3 85.0 84.7 84.4 84.1 83.8 83.4 83.0 82.6 82.2 81.8 81.3 81.1 80.8 80.6 80.3 80.0 79.8 79.5 79.2 78.9 78.6 78.4 78.0 77.8

B scale, 100 kgf load, 1/16 in. diam ball

C scale, 150 kgf load, diamond indenter

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

68.0 67.5 67.0 66.4 65.9 65.3 64.7 64.0 63.3 62.5 61.8 61.0 60.1 59.7 59.2 58.8 58.3 57.8 57.3 56.8 56.3 55.7 55.2 54.7 54.1 53.6

D scale, 100 kgf load, 15 N scale, diamond 15 kgf indenter load

76.9 76.5 76.1 75.7 75.3 74.8 74.3 73.8 73.3 72.6 72.1 71.5 70.8 70.5 70.1 69.8 69.4 69.0 68.7 68.3 67.9 67.5 67.0 66.7 66.2 65.8

93.2 93.0 92.9 92.7 92.5 92.3 92.1 91.8 91.5 91.2 91.0 90.7 90.3 90.1 89.8 89.7 89.5 89.2 89.0 88.8 88.5 88.2 88.0 87.8 87.5 87.2

30 N scale, 30 kgf load

84.4 84.0 83.6 83.1 82.7 82.2 81.7 81.1 80.4 79.7 79.1 78.4 77.6 77.2 76.8 76.4 75.9 75.5 75.1 74.6 74.2 73.6 73.2 72.7 72.1 71.7

Knoop 45 N scale, hardness No., 500 gf load 45 kgf and greater load

75.4 74.8 74.2 73.6 73.1 72.2 71.8 71.0 70.2 69.4 68.6 67.7 66.7 66.2 65.7 65.3 64.7 64.1 63.5 63.0 62.4 61.7 61.2 60.5 59.9 59.3

920 908 895 882 867 852 837 822 806 788 772 754 735 725 716 706 697 687 677 667 657 646 636 625 615 604

Scleroscope hardness No.

97 96 95 93 92 91 90 88 87 86 84 83 81 ... 80 ... 79 78 77 76 75 ... 74 73 72 ...

(continued) Note: Values in parentheses are beyond normal range and are given for information only. Data are for carbon and alloy steels in the annealed, normalized, and quenchedand-tempered conditions; less accurate for cold-worked condition and for austenitic steels. Source: ASTM E140


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Table A3.5

(continued)

Vickers hardness No.

560 550 540 530 520 510 500 490 480 470 460 450 440 430 420 410 400 390 380 370 360 350 340 330 320 310 300 295 290 285 280 275 270 265 260 255 250 245 240 230 220 210 200 190 180 170 160 150 140 130 120 110 100 95 90 85

Standard ball

... (505) (496) (488) (480) (473) (465) (456) (448) 441 433 425 415 405 397 388 379 369 360 350 341 331 322 313 303 294 284 280 275 270 265 261 256 252 247 243 238 233 228 219 209 200 190 181 171 162 152 143 133 124 114 105 95 90 86 81

A scale, 60 kgf Tungstenload, carbide diamond ball indenter

525 517 507 497 488 479 471 460 452 442 433 425 415 405 397 388 379 369 360 350 341 331 322 313 303 294 284 280 275 270 265 261 256 252 247 243 238 233 228 219 209 200 190 181 171 162 152 143 133 124 114 105 95 90 86 81

Rockwell superficial (diamond pyramid) hardness No., diamond indenter

Rockwell hardness No.

Brinell hardness No., 3000 kg load, 10 mm ball

77.4 77.0 76.7 76.4 76.1 75.7 75.3 74.9 74.5 74.1 73.6 73.3 72.8 72.3 71.8 71.4 70.8 70.3 69.8 69.2 68.7 68.1 67.6 67.0 66.4 65.8 65.2 64.8 64.5 64.2 63.8 63.5 63.1 62.7 62.4 62.0 61.6 61.2 60.7 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

B scale, 100 kgf load, 1/16 in. diam ball

C scale, 150 kgf load, diamond indenter

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... (110.0) ... (109.0) ... (108.0) ... (107.0) ... (105.5) ... (104.5) ... (103.5) ... (102.0) ... (101.0) ... 99.5 ... 98.1 96.7 95.0 93.4 91.5 89.5 87.1 85.0 81.7 78.7 75.0 71.2 66.7 62.3 56.2 52.0 48.0 41.0

53.0 52.3 51.7 51.1 50.5 49.8 49.1 48.4 47.7 46.9 46.1 45.3 44.5 43.6 42.7 41.8 40.8 39.8 38.8 37.7 36.6 35.5 34.4 33.3 32.2 31.0 29.8 29.2 28.5 27.8 27.1 26.4 25.6 24.8 24.0 23.1 22.2 21.3 20.3 (18.0) (15.7) (13.4) (11.0) ( 8.5) (6.0) (3.0) (0.0) ... ... ... ... ... ... ... ... ...

D scale, 100 kgf load, 15 N scale, diamond 15 kgf indenter load

65.4 64.8 64.4 63.9 63.5 62.9 62.2 61.6 61.3 60.7 60.1 59.4 58.8 58.2 57.5 56.8 56.0 55.2 54.4 53.6 52.8 51.9 51.1 50.2 49.4 48.4 47.5 47.1 46.5 46.0 45.3 44.9 44.3 43.7 43.1 42.2 41.7 41.1 40.3 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

86.9 86.6 86.3 86.0 85.7 85.4 85.0 84.7 84.3 83.9 83.6 83.2 82.8 82.3 81.8 81.4 80.8 80.3 79.8 79.2 78.6 78.0 77.4 76.8 76.2 75.6 74.9 74.6 74.2 73.8 73.4 73.0 72.6 72.1 71.6 71.1 70.6 70.1 69.6 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

30 N scale, 30 kgf load

71.2 70.5 70.0 69.5 69.0 68.3 67.7 67.1 66.4 65.7 64.9 64.3 63.5 62.7 61.9 61.1 60.2 59.3 58.4 57.4 56.4 55.4 54.4 53.6 52.3 51.3 50.2 49.7 49.0 48.4 47.8 47.2 46.4 45.7 45.0 44.2 43.4 42.5 41.7 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

Knoop 45 N scale, hardness No., 45 kgf 500 gf load load and greater

58.6 57.8 57.0 56.2 55.6 54.7 53.9 53.1 52.2 51.3 50.4 49.4 48.4 47.4 46.4 45.3 44.1 42.9 41.7 40.4 39.1 37.8 36.5 35.2 33.9 32.5 31.1 30.4 29.5 28.7 27.9 27.1 26.2 25.2 24.3 23.2 22.2 21.1 19.9 ... ... . . .. ... ... ... ... ... ... ... ... ... ... ... ... ... ...

594 583 572 561 550 539 528 517 505 494 482 471 459 447 435 423 412 400 389 378 367 356 346 337 328 318 309 305 300 296 291 286 282 277 272 267 262 258 253 243 234 226 216 206 196 185 175 164 154 143 133 123 112 107 102 97

Scleroscope hardness No.

71 70 69 68 67 ... 66 65 64 ... 62 ... 59 58 57 56 55 ... 52 51 50 48 47 46 45 ... 42 ... 41 ... 40 39 38 ... 37 ... 36 35 34 33 32 30 29 28 26 25 23 22 21 20 18 ... ... ... ... ...

Note: Values in parentheses are beyond normal range and are given for information only. Data are for carbon and alloy steels in the annealed, normalized, and quenchedand-tempered conditions; less accurate for cold-worked condition and for austenitic steels. Source: ASTM E140


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1000

1000

900

900

800

800 Vickers hardness

Vickers hardness

Appendix 3:

700 600 500

45 N

600 500 400

300

300 200 0

10

20

30

40

50

60

70

80

90

100

0

10

700

700 Brinell hardness number

Brinell hardness number

800

600 500 400 300

40

50

60

70

80

90

100

45 N

30 N

500

15 N

400 300

0

10

20

30

40

50

60

70

80

90

200

100

0

10

1000

70

800

60

600 400 200

0

100 200 300 400 500 600 700 800 900 1000 Brinell hardness number

20

30

40

50

60

70

80

90

100

Rockwell superficial hardness number

Rockwell C hardness

Vickers hardness

30

600

Rockwell C hardness

Fig. A3.1

20


800

0

15 N

30 N

Rockwell C hardness

200


700

400

200

pg 535

45 N

50

30 N

40

15 N

30 20 0

10

20

30

40

50

60

70

80

90

100


Approximate equivalent hardness numbers for steel. Points represent data from the hardness conversion tables.


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APPENDIX 4

Austenitizing Temperatures for Steels TEMPERATURES RECOMMENDED for austenitizing carbon and low-alloy steels prior to hardening are given in Table A4.1 (for direct-hardening grades) and Table A4.2 (for carburized steels). Table A4.2 is applicable to carburized steels that have been cooled slowly from the carburizing temperature and are to be furnace hardened in a subsequent operation. For most applications, the rate of heating to the austenitizing temperature is less important than other factors in the hardening process, such as maximum temperature attained throughout the section, temperature uniformity, time at temperature, and rate of cooling. The thermal conductivity of the steel, the nature of the furnace atmosphere (scaling or nonscaling), thickness of section, method of loading (spaced or stacked), and the degree of circulation of

the furnace atmosphere all influence the rate of heating of the steel part to the required temperature selected from Tables A4.1 and A4.2. The difference in temperature rise within thick and thin sections of articles of varying cross section is a major problem in practical heat-treating operations. When temperature uniformity is the ultimate objective of the heating cycle, this is more safely attained by slowly heating than by rapidly heating. Furthermore, the maximum temperature in the austenite range should not exceed that required to achieve the necessary extent of solution of carbide. The temperatures listed in Tables A4.1 and A4.2 conform with this requirement. When heating with significant cross-section variations, provisions should be made for slower heating to minimize thermal stresses and distortions.

Table A4.1 Austenitizing temperatures for direct-hardening carbon and alloy steels (SAE) Temperature Steel

°C

Temperature °F

Carbon steels 1025 1030 1035 1037 1038(a) 1039(a) 1040(a) 1042 1043(a) 1045(a) 1046(a) 1050(a) 1055 1060 1065 1070 1074 1078 1080 1084 1085 1086 1090 1095

Steel

°C

°F

Free-cutting carbon steels 855–900 845–870 830–855 830–855 830–855 830–855 830–855 800–845 800–845 800–845 800–845 800–845 800–845 800–845 800–845 800–845 800–845 790–815 790–815 790–815 790–815 790–815 790–815 790–815(a)

1575–1650 1550–1600 1525–1575 1525–1575 1525–1575 1525–1575 1525–1575 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1450–1500 1450–1500 1450–1500 1450–1500 1450–1500 1450–1500 1450–1500(b)

1137 1138 1140 1141 1144 1145 1146 1151 1536 1541 1548 1552 1566

830–855 815–845 815–845 800–845 800–845 800–845 800–845 800–845 815–845 815–845 815–845 815–845 855–885

1525–1575 1500–1550 1500–1550 1475–1550 1475–1550 1475–1550 1475–1550 1475–1550 1500–1550 1500–1550 1500–1550 1500–1550 1575–1625

830–855 815–845 815–845 815–845 815–845 830–855 830–855 815–855

1525–1575 1500–1550 1500–1550 1500–1550 1500–1550 1525–1575 1525–1575 1500–1575

Alloy steels 1330 1335 1340 1345 3140 4037 4042 4047 (continued)

(a) Commonly used on parts where induction hardening is employed. All steels from SAE 1030 up may have induction hardening applications. (b) This temperature range may be employed for 1095 steel that is to be quenched in water, brine, or oil. For oil quenching, 1095 steel may alternatively be austenitized in the range 815 to 870 C (1500 to 1600 F). (c) This range is recommended for steel that is to be water quenched. For oil quenching, steel should be austenitized in the range 815 to 870 C (1500 to 1600 F).


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Table A4.1

(continued) Temperature

Temperature °C

Steel

4063 4130 4135 4137 4140 4142 4145 4147 4150 4161 4337 4340 50B40 50B44 5046 50B46 50B50 50B60 5130 5132 5135 5140 5145 5147 5150

800–845 815–870 845–870 845–870 845–870 845–870 815–845 815–845 815–845 815–845 815–845 815–845 815–845 815–845 815–845 815–845 800–845 800–845 830–855 830–855 815–845 815–845 815–845 800–845 800–845

°F

°C

°F

800–845 800–845 800–845 775–800(c) 775–800(c) 775–800(c) 845–885 815–855 830–870 830–855 830–855 815–855 815–855 815–855 815–855 800–845 800–845 830–855 830–855 815–900 815–900 815–900 845–885 845–885 830–855

1475–1550 1475–1550 1475–1550 1425–1475(c) 1425–1475(c) 1425–1475(c) 1550–1625 1500–1575 1525–1600 1525–1575 1525–1575 1500–1575 1500–1575 1500–1575 1500–1575 1475–1550 1475–1550 1525–1575 1525–1575 1500–1650 1500–1650 1500–1650 1550–1625 1550–1625 1525–1575

Steel

1475–1550 1500–1600 1550–1600 1550–1600 1550–1600 1550–1600 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1475–1550 1475–1550 1525–1575 1525–1575 1500–1550 1500–1550 1500–1550 1475–1550 1475–1550

5155 5160 51B60 50100 51100 52100 6150 81B45 8630 8637 8640 8642 8645 86B45 8650 8655 8660 8740 8742 9254 9255 9260 94B30 94B40 9840

(a) Commonly used on parts where induction hardening is employed. All steels from SAE 1030 up may have induction hardening applications. (b) This temperature range may be employed for 1095 steel that is to be quenched in water, brine, or oil. For oil quenching, 1095 steel may alternatively be austenitized in the range 815 to 870 C (1500 to 1600 F). (c) This range is recommended for steel that is to be water quenched. For oil quenching, steel should be austenitized in the range 815 to 870 C (1500 to 1600 F).

Table A4.2 Reheating (austenitizing) temperatures for hardening of carburized carbon and alloy steels (SAE)

Carburizing is commonly carried out at 900 to 925 C (1650 to 1700 F); slow cooled and reheated to given austenizing temperature. Temperature Steel

°C

Temperature °F

Carbon steels 1010 1012 1015 1016 1017 1018 1019 1020 1022 1513 1518 1522 1524 1525 1526 1527

°C

°F

790–830 830–845 815–845 815–845 815–845 815–845 815–845 815–845 815–845 800–830 800–830 800–830 845–870 845–870 845–870 845–870 845–870 845–870 845–870 845–870 845–870 790–830

1450–1525 1525–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1500–1550 1475–1525 1475–1525 1475–1525 1550–1600 1550–1600 1550–1600 1550–1600 1550–1600 1550–1600 1550–1600 1550–1600 1550–1600 1450–1525

Alloy steels 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790 760–790

1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450 1400–1450

760–790 760–790 760–790 760–790

1400–1450 1400–1450 1400–1450 1400–1450

Free-cutting carbon steels 1109 1115 1117 1118

Steel

3310 4320 4615 4617 4620 4621 4626 4718 4720 4815 4817 4820 8115 8615 8617 8620 8622 8625 8627 8720 8822 9310

Name ///sr-nova/Dclabs_wip/Failure_Analysis/5113_539-540.pdf/Appendix-5/

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pg 539


APPENDIX 5

Temper Colors for Steels IT IS COMMON and long-time knowledge that steel heated in contact with air at temperatures in the tempering range takes on various temper colors due to the formation of a thin oxide film. As an example, the 1948 ASM Metals Handbook gives the following table: Heating temperature, °F

400 440 475 520 540 590 640

Copper/Steel/Stainless Steel Test.

Place a piece of bright copper strip, a bright carbon steel strip and a bright stainless steel strip side by side on the belt and send them through the furnace. Use the following chart to determine the source of oxidation and the color chart to determine its location:

Color

Faint straw Straw Deep straw Bronze Peacock Full blue Light blue

Color Chart (back cover). Another example of temper colors is shown on the back cover of this book. This example was developed to illustrate the effect of oxidation with furnace leaks. Oxidation of steel in the front of a continuous furnace results in a frosted or matt surface finish with a loss of surface brightness. Oxidation in the high heat section is evidenced by a black oxide that is often scaled and flakes from the surface of the steel. If oxidation is occurring in the cooling section of a continuous furnace, use the following test to determine the cause:

Change the atmosphere flow in the furnace to 100% Nitrogen Make absolutely sure that every zone of the furnace is below 1900 F to prevent copper from melting

Sample surface appearance Copper

Oxidized Bright Bright

Carbon steel

Stainless steel

Conclusions

Oxidized Oxidized Bright

Oxidized Oxidized Oxidized

Air leak Water leak Very small air Or water leak

Time-Temperature Effect. A study at the Illinois Institute of Technology obtained detailed information on temper colors for plain carbon steel, especially on the effect of time and temperature. This was obtained using 16 mm (5/8 in.) diam hot-rolled bars of a standard SAE 1035 steel. Samples that were 50 mm (2 in.) long were cut from the bars and carefully machined and cleaned to give smooth, bright surfaces. They were then heated for various times at several temperatures in air-circulating furnaces controlled to within +3 C (+5 F) of the desired temperatures. The results of this study are shown in Fig. A5.1 for 1035 steel.


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Fig. A5.1

Temper colors after heating 1035 steel in circulating air (atmospheric pressure)

pg 540


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pg 541


APPENDIX 6

Physical Properties of Carbon and Low-Alloy Steels Thermal Expansion Table A6.1 Coefficients of linear thermal expansion for carbon and low-alloy steels Average coefficient of expansion, mm/m K, at °C (°F)(a) AISI-SAE grade

1008 1008 1010 1010 1010 1015 1015 1016 1017 1018 1019 1020 1020 1021 1022 1023 1025 1026 1029 1030 1035 1037 1039 1040 1043 1044 1045 1045 1046 1050 1052 1053 1055 1060

Treatment or condition

Annealed Annealed Annealed Unknown Unknown Rolled Annealed Annealed Unknown Annealed Unknown Annealed Unknown Unknown Annealed Unknown Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Unknown Annealed Annealed Unknown Annealed Annealed

20–100 (68–212)

12.6(b) 11.6 12.2(b) 11.9(c) ... 11.9(b) 12.2(b) 12.0(b) 12.2(b) 12.0(b) 12.2(b) 11.7 12.2(b) 12.0(b) 12.2(b) 12.2(b) 12.0(b) 12.0(b) 12.0(b) 11.7(b) 11.1 11.1(b) 11.1(b) 11.3 11.3(b) 11.1(b) 11.6(b) 11.2(e) 11.1(b) 11.1(b) 11.3(e) 11.1(b) 11.0 11.1(e)

20–200 (68–392)

20–300 (68–572)

20–400 (68–752)

20–500 (68–932)

20–600 (68–1112)

20–700 (68–1292)

13.1(b) 12.5 13.0(b) 12.6 ... 12.5(b) ... ... ... ... ... 12.1 ... ... 12.7(b) ... ... ... ... ... 11.9 ... ... 12.0 ... 12.0(b) 12.3(b) 11.9(e) ... 12.0(b) 11.8(e) ... 11.8 11.9(e)

13.5(b) 13.0 13.5(b) 13.3 ... 13.0(b) ... ... ... ... ... 12.8 ... ... 13.1(b) ... ... ... ... ... 12.7 ... ... 12.5 ... ... 13.1(b) 12.7(e) ... ... 12.7(e) ... 12.6 12.9(e)

13.8(b) 13.6 13.9(b) 13.8 15.1(d) 13.6(b) 13.4(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.4 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 13.4 13.5(b) 13.5(b) 13.3 13.5(b) 13.3(b) 13.7(b) 13.5(e) 13.5(b) 13.5(b) 13.7(e) 13.5(b) 13.4 13.5(e)

14.2(b) 14.2 14.3(b) 14.3 ... 14.2(b) ... ... ... ... ... 13.9 ... ... 13.9(b) ... ... ... ... ... 14.0 ... ... 13.9 ... ... 14.2(b) 14.1(e) ... ... 14.5(e) ... 14.0 14.1(e)

14.6(b) 14.6 14.7(b) 14.7 ... ... 14.2(b) 14.4(b) 14.5(b) 14.4(b) 14.7(b) 14.4 14.2(b) 14.3(b) 14.4(b) 14.4(b) 14.4(b) 14.4(b) 14.4(b) 14.4(b) 14.4 14.6(b) 14.6(b) 14.4 14.6(b) ... 14.7(b) 14.5(e) ... ... 14.7(e) ... 14.5 14.6(e)

15.0(b) 15.0 15.0(b) 14.9 ... ... ... ... ... ... ... 14.8 ... ... 14.9(b) ... ... ... ... ... 14.8 ... ... 14.8 ... ... 15.1(b) 14.8(e) ... ... 15.0(e) ... 14.8 14.9(e)

(continued) (a) To obtain coefficients in min./in. F, multiply stated values by 0.556. (b) Stated value represents average coefficient between 0 C (32 F) and indicated temperature. (c) 10.3 mm/m K from 100 to 20 C ( 148 to 68 F); 9.8 mm/m K from 150 to 20 C ( 238 to 68 F). (d) Stated value represents average coefficient between 20 and 650 C (68 and 1200 F). (e) Stated value represents average coefficient between 25 C (75 F) and indicated temperature. (f) Stated value represents average coefficient between 20 and 95 C (68 and 200 F). (g) Stated value represents average coefficient between 20 and 370 C (68 and 700 F). (h) 8.6 mm/m K from 195 to 20 C ( 320 to 68 F); 10.0 mm/m K from 130 to 20 C ( 200 to 68 F). (i) Stated value represents average coefficient between 20 and 260 C (68 and 500 F). (j) Stated value represents average coefficient between 20 and 540 C (68 and 1000 F). (k) Stated value represents average coefficient between 18 and 95 C (0 and 200 F). (l) Stated value represents average coefficient between 18 and 650 C (0 and 1200 F). (m) 11.2 mm/m K from 100 to 20 C ( 148 to 68 F) 10.4 mm/m K from 150 to 20 C ( 238 to 68 F). (n) Stated value represents average coefficient between 20 and 205 C (68 and 400 F). (o) Stated value represents average coefficient between 20 and 315 C (68 and 600 F). ( p) Stated value represents average coefficient between 25 and 270 C (77 and 518 F). (q) Stated value represents average coefficient between 20 and 275 C (68 and 525 F). (r) Stated value represents average coefficient between 18 and 260 C (0 and 500 F). (s) Stated value represents average coefficient between 18 and 540 C (0 and 1000 F).


18/8/2008 4:32PM Plate # 0

pg 542


Table A6.1

(continued) Average coefficient of expansion, mm/m K, at °C (°F)(a)

AISI-SAE grade

1064 1065 1070 1078 1080 1080 1085 1086 1095 1095 1095 1117 1118 1132 1137 1139 1140 1141 1144 1145 1145 1146 1151 1330 1335 1345 1522 1524 1524 1526 1541 1548 1551 1552 1561 1566 2330 2515 3120 3130 3140 3150 4023 4027 4028 4032 4042 4047 4130 4135 4137 4140 4142 4145 4147 4161 4320 4337


20–100 (68–212)

20–200 (68–392)

20–300 (68–572)

20–400 (68–752)

20–500 (68–932)

20–600 (68–1112)

20–700 (68–1292)

Unknown Unknown Unknown Unknown Annealed Unknown Annealed Unknown Unknown Annealed Hardened Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Annealed Annealed Unknown Unknown Unknown Unknown Unknown Annealed Unknown Annealed Annealed Annealed Unknown Annealed Unknown Annealed Annealed Annealed Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Oil hardened, tempered Unknown Oil hardened, tempered Unknown Unknown Unknown Unknown

11.1(b) 11.1(b) 11.5(b) 11.3(b) 11.0 11.7(b) 11.1(b) 11.1(b) ... 11.4 13.0(b) 12.2(f) 12.2(f) 12.6(f) 12.8 12.6(f) 12.6(f) ... 13.3 11.2(b) 11.6(b) 12.8 ... 12.0 12.2 12.0 12.0(b) 11.9 12.0(b) 12.0(b) 12.0(b) 11.9(b) 11.7(b) 11.1(b) 11.1(b) 11.5(b) 10.9(e) 10.9(f)(h) 11.3(k) 11.3(k) 11.3(k) 11.3(k) 11.7(k) 11.7(k) 11.9 11.9 11.9 11.9 12.2 11.7 11.7 12.3 11.7 11.7 11.7 11.5 11.3(k) 11.3(k)

... ... ... ... 11.6 12.2(b) 11.7(b) ... ... ... ... ... ... ... ... ... ... 12.6(b) ... 12.1(b) 12.3(b) ... 12.6(b) 12.8 12.8 12.6 ... 12.7 ... ... ... ... ... ... ... ... 11.2(e) ... ... ... ... ... ... ... 12.4 12.4 12.4 12.4 ... 12.2 12.2 12.7 12.2 12.2 12.2 12.2 ... ...

... ... ... ... 12.4 ... 12.5(b) ... ... ... ... ... ... ... ... ... ... ... ... 13.0(b) 13.1(b) ... ... 13.3 13.3 13.3 ... ... ... ... ... ... ... ... ... ... 12.1(e) 12.6(i) ... ... ... ... ... ... 12.9 12.9 12.9 12.9 ... 12.8 12.8 ... 12.8 12.8 12.8 12.9 ... ...

13.5(b) 13.5(b) 13.3(b) 13.3 13.2 ... 13.2(b) 13.1(b) ... ... ... 13.1(g) 13.3(g) ... ... ... ... ... ... 13.6(b) 13.7(b) ... ... ... ... ... 13.5(b) 13.9 13.5(b) 13.5(b) 13.5(b) 13.3(b) 13.5(b) 13.5(b) 13.5(b) 13.5(b) 12.9(e) ... ... ... ... ... ... ... ... ... ... ... 13.7 ... ... 13.7 ... ... ... ... ... ...

... ... ... ... 13.8 ... 13.6(b) ... ... ... ... ... ... ... ... ... ... ... ... 14.0(b) 14.2(b) ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 13.4(e) 13.5(j) ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

... ... ... ... 14.2 ... 14.2(b) ... ... ... ... ... ... ... ... ... ... ... ... 14.6(b) 14.7(b) ... ... ... ... ... 14.4(b) 14.7 14.4(b) 14.4(b) 14.4(b) 14.6(b) 14.6(b) ... 14.6(b) 14.7(b) 13.8(e) ... ... ... ... ... ... ... ... ... ... ... 14.6 ... ... 14.5 ... ... ... ... ... ...

... ... ... ... 14.7 ... 14.7 ... 14.6(d) ... ... ... ... ... ... ... ... ... ... 14.8(b) 15.1(b) ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 14.6(l) 14.6(l) 14.6(l) 14.6(l) ... ... ... ... ... ... ... ... ... ... ... ... ... ... 14.6(l) 14.6(l)

(continued) (a) To obtain coefficients in min./in. F, multiply stated values by 0.556. (b) Stated value represents average coefficient between 0 C (32 F) and indicated temperature. (c) 10.3 mm/m K from 100 to 20 C ( 148 to 68 F); 9.8 mm/m K from 150 to 20 C ( 238 to 68 F). (d) Stated value represents average coefficient between 20 and 650 C (68 and 1200 F). (e) Stated value represents average coefficient between 25 C (75 F) and indicated temperature. (f) Stated value represents average coefficient between 20 and 95 C (68 and 200 F). (g) Stated value represents average coefficient between 20 and 370 C (68 and 700 F). (h) 8.6 mm/m K from 195 to 20 C ( 320 to 68 F); 10.0 mm/m K from 130 to 20 C ( 200 to 68 F). (i) Stated value represents average coefficient between 20 and 260 C (68 and 500 F). (j) Stated value represents average coefficient between 20 and 540 C (68 and 1000 F). (k) Stated value represents average coefficient between 18 and 95 C (0 and 200 F). (l) Stated value represents average coefficient between 18 and 650 C (0 and 1200 F). (m) 11.2 mm/m K from 100 to 20 C ( 148 to 68 F) 10.4 mm/m K from 150 to 20 C ( 238 to 68 F). (n) Stated value represents average coefficient between 20 and 205 C (68 and 400 F). (o) Stated value represents average coefficient between 20 and 315 C (68 and 600 F). ( p) Stated value represents average coefficient between 25 and 270 C (77 and 518 F). (q) Stated value represents average coefficient between 20 and 275 C (68 and 525 F). (r) Stated value represents average coefficient between 18 and 260 C (0 and 500 F). (s) Stated value represents average coefficient between 18 and 540 C (0 and 1000 F).


Appendix 6:

18/8/2008 4:32PM Plate # 0

pg 543

Physical Properties of Carbon and Low-Alloy Steels / 543

Table A6.1 (continued) Average coefficient of expansion, mm/m K, at °C (°F)(a) AISI-SAE grade

4340 4340 4422 4427 4615 4617 4626 4718 4815 4820 5046 50B60 5117 5120 5130 5132 5135 5140 5145 5150 5155 52100 52100 6150 6150 615 0 8115 81B45 8617 8622 8625 8627 8630 8637 8645 8650 8655 8720 8822


Oil hardened, tempered 600 C(1110 F) Oil hardened, tempered 630 C (1170 F) Unknown Unknown Unknown Carburized and hardened Normalized and tempared Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Annealed Unknown Unknown Unknown Annealed Hardened Annealed Hardened, tempered 205 C (400 F) Annealed Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Oil hardened, tempered Oil hardened, tempered Oil hardened, tempered Unknown Unknown

20–100 (68–212)

12.3 . . .(m)

20–200 (68–392)

20–300 (68–572)

20–400 (68–752)

20–500 (68–932)

20–600 (68–1112)

20–700 (68–1292)

12.7

...

13.7

...

14.5

...

12.4

...

13.6

...

14.3

...

11.7(k) 12.6 11.5 12.5 11.70(f) 11.3 11.5(f) 11.3(f) 11.9 11.9 12.0 12.0 12.2 12.2 12.0 ... 12.2 12.8 12.2 11.9(b) 12.6(b) 12.2 12.0

... ... 12.1 13.1 ... 12.2 12.2(n) 12.2(n) 12.4 12.4 12.8 12.8 12.9 12.9 12.8 12.6 ... ... ... ... ... 12.7 12.5

... 13.8 12.7 ... 12.6(i) 13.1 13.1(o) 12.9(o) 12.9 12.9 13.5 13.5 13.5 13.5 13.5 13.4 13.1( p) 13.7( p) 13.1(q) ... ... 12.3 12.9

... ... 13.2 ... ... ... ... ... ... ... ... ... ... ... ... 13.9 ... ... ... ... ... 13.7 13.0

... ... 13.7 ... ... ... ... ... ... ... ... ... ... ... ... 14.3 ... ... ... ... ... 14.1 13.3

... 15.1 14.1 ... 13.8(j) ... ... ... ... ... ... ... ... ... ... 14.6 ... ... ... ... ... 14.4 13.7

... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 15.0 ... ... ... ... ... ... ...

12.4(e) 11.9 11.9 11.9(k) 11.1 11.1 11.3 11.3(k) 11.3 11.7 11.7 11.7 11.3(k) 11.3

12.6(e) 12.6 12.6 ... 12.2 12.2 12.2 ... 12.2 12.2 12.2 12.2 ... 12.2

13.3(e) 13.3 13.3 12.8(r) 12.9 12.9 12.9 ... 12.8 12.8 12.8 12.8 ... 12.9

13.8(e) ... ... ... ... ... ... ... ... ... ... ... ... ...

14.2(e) ... ... ... ... ... ... ... ... ... ... ... ... ...

14.5(e) ... ... 14.0(s) ... ... ... ... ... ... ... ... ... ...

14.7(e) ... ... ... ... ... ... 14.6(l) ... ... ... ... 14.6(l) ...

(a) To obtain coefficients in min./in. F, multiply stated values by 0.556. (b) Stated value represents average coefficient between 0 C (32 F) and indicated temperature. (c) 10.3 mm/m K from 100 to 20 C ( 148 to 68 F); 9.8 mm/m K from 150 to 20 C ( 238 to 68 F). (d) Stated value represents average coefficient between 20 and 650 C (68 and 1200 F). (e) Stated value represents average coefficient between 25 C (75 F) and indicated temperature. (f) Stated value represents average coefficient between 20 and 95 C (68 and 200 F). (g) Stated value represents average coefficient between 20 and 370 C (68 and 700 F). (h) 8.6 mm/m K from 195 to 20 C ( 320 to 68 F);10.0 mm/m K from 130 to 20 C ( 200 to 68 F). (i) Stated value represents average coefficient between 20 and 260 C (68 and 500 F). (j) Stated value represents average coefficient between 20 and 540 C (68 and 1000 F). (k) Stated value represents average coefficient between 18 and 95 C (0 and 200 F). (l) Stated value represents average coefficient between 18 and 650 C (0 and 1200 F). (m) 11.2 mm/m K from 100 to 20 C ( 148 to 68 F) 10.4 mm/m K from 150 to 20 C ( 238 to 68 F). (n) Stated value represents average coefficient between 20 and 205 C (68 and 400 F). (o) Stated value represents average coefficient between 20 and 315 C (68 and 600 F). ( p) Stated value represents average coefficient between 25 and 270 C (77 and 518 F). (q) Stated value represents average coefficient between 20 and 275 C (68 and 525 F). (r) Stated value represents average coefficient between 18 and 260 C (0 and 500 F). (s) Stated value represents average coefficient between 18 and 540 C (0 and 1000 F).


18/8/2008 4:32PM Plate # 0


Table A6.2 Summary of thermal expansion Coefficient of linear thermal expansion (CTE), approximate ranges at room temperature to 100 C (212 F), from lowest to highest CTE value CTE

CTE

10 6/K

10 6/°F

Material

10 6/K

10 6/°F

2.6–3.3 2.2–6.1 4.5–4.6 0.6–8.7 4.8–5.1 5.6 6.0 6.1 5.7–7.0 6.3–6.6 6.2–6.7 6.5 4.9–8.2 6.8 2.0–12 7.1 7.2–7.3 5.1–9.6 4.5–11 7.1–9.7 8.3–8.5 8.3–8.4 5.5–11 8.4–8.6 8.6–8.7 7.6–9.9 7.7–10 4.0–14 8.8–9.1 7.6–11 9.3–9.6 9.3–9.9 9.1–10 8.4–11 8.6–11 9.9 9.8–10 10–11 6.8–14 9.3–12 7.6–14 11 8.9–12 9.5–12 9.9–12 11 10–12 10–12 9.3–12 9.8–13 10–12

1.4–1.8 1.2–3.4 2.5–2.6 0.3–4.8 2.7–2.8 3.1 3.3 3.4 3.2–3.9 3.5–3.7 3.4–3.7 3.6 2.7–4.6 3.8 1.1–6.7 3.9 4.0–4.1 2.8–5.3 2.5–6.2 3.9–5.4 4.6–4.7 4.6–4.7 3.1–6.3 4.7–4.8 4.8–4.8 4.2–5.5 4.3–5.7 2.2–7.8 4.9–5.1 4.2–5.9 5.2–5.3 5.2–5.5 5.1–5.6 4.7–6.3 4.8–6.3 5.5 5.4–5.8 5.6–5.9 3.8–7.8 5.2–6.5 4.2–7.5 5.9 4.9–6.9 5.3–6.6 5.5–6.5 6.1 5.6–6.6 5.6–6.5 5.2–6.9 5.4–6.9 5.8–6.7

11–12 12 11–13

6.1–6.9 6.5 5.9–7.1

11–13 10–14 12 8.8–15 11–14 9.4–15 12–13

6.2–7.0 5.6–7.6 6.6 4.9–8.4 5.9–7.5 5.2–8.2 6.5–7.0

12 11–14 11–14 7.6–17

6.8 5.9–7.6 6.0–7.5 4.2–9.4

11 8.5–14 11 7.1–16 9.3–13 11–12 11 11 10–13 7.6–15 11–12 6.3–17 10–13

6.2 4.7–7.8 6.3 3.9–8.7 5.2–7.2 6.1–6.6 6.3 6.4 5.7–7.0 4.2–8.5 6.1–6.8 3.5–9.4 5.7–7.3

Pure Silicon (Si) Pure Osmium (Os) Pure Tungsten (W) Iron-cobalt-nickel alloys Pure Molybdenum (Mo) Pure Arsenic (As) Pure Germanium (Ge) Pure Hafnium (Hf) Pure Zirconium (Zr) Pure Cerium (Ce) Pure Rhenium (Re) Pure Tantalum (Ta) Pure Chromium (Cr) Pure Iridium (Ir) Magnetically soft iron alloys Pure Technetium (Tc) Pure Niobium (Nb) Pure Ruthenium (Ru) Pure Praseodymium (Pr) Beta and near beta titanium Pure Rhodium (Rh) Purr Vanadium (V) Zirconium alloys Pure Titanium (Ti) Mischmetal Unalloyed or low-alloy titanium Alpha beta titanium Molybdenum alloys Pure Platinum (Pt) Alpha and neat alpha titanium High-chromium gray cast iron Ductile high-chromium cast iron Pure Gadolinium (Gd) Pure Antimony (Sb) Maraging steel Protactinium (Pa) Water-hardening tool steel Molybdenum high-speed too steel Niobium alloys Ferritic stainless steel Pure Neodymium (Nd) Cast ferritic stainless steel Hot work tool steel Martensitic stainless steel Cast martensitic stainless steel Cermet Ductile silicon-molybdenum cast iron Iron carbon alloys Pure Terbium (Tb) Cobalt chromium nickel tungsten High-carbon high-chromium cold work tool steel Tungsten high-speed tool steel Commercially pure or low-alloy nickel Low-alloy special purpose tool steel Pure Dysprosium (Dy) Nickel molybdenum alloy steel Pure Palladium (Pd) Pure Thorium (Th) Wrought iron Oil-hardening cold work tool steel Pure Scandium (Sc) Pure Beryllium (Be) Carbide Nickel chromium molybdenum alloy steel

11–14 12–13 4.8–20 10–15 9.9–13 9.0–16 12–13 11–14 10–15 6.0–20 11–15 9.0–17 13 7.0–20 11–16 13 14 12–14 10–17 13–15 8.1–19 14 7.0–20 14 10–19 7.9–21 13–16 14 14–15 12–18 10–20 9.7–19 15 12–19 8.8–22 14–18 13–19 4.5–27 16–18 17 15–19 17–18 16–18 17 9.8–25 16–19 16–19 18

6.2–7.5 6.4–7.4 2.7–11 5.6–8.3 5.5–7.3 5.0–8.9 6.5–7.4 5.9–8.0 5.6–8.6 3.3–11 6.0–8.5 5.0–9.6 7.4 3.9–11 6.1–8.6 7.4 7.5 6.8–7.7 5.6–9.6 7.0–8.2 4.5–11 7.8 3.9–11 7.8 5.3–11 4.4–12 7.0–9.0 7.8 7.7–8.4 6.7–10 5.6–11 5.4–11 8.5 6.7–10 4.9–12 7.5–9.8 7.0–10 2.5–15 8.8–10 9.4 8.3–11 9.2–9.8 9.1–10 9.6 5.4–14 8.9–11 8.9–11 10

(continued) Source: Thermal Properties of Metals and Alloys, ASM, 2002

Material

Shock-resisting tool steel Structural steel Air hardening medium-alloy cold work tool steel High manganese carbon steel Malleable cast iron Mold tool steel Nonresulfurized carbon steel Chromium molybdenum alloy steel Chromium alloy steel Molybdenum/molybdenum sulfide alloy steel Chromium Vanadium alloy steel Cold work tool steel Ductile medium-silicon cast iron Nickel with chromium and/or iron, molybdenum Resulfurized carbon steel High strength low-alloy steel (HSLA) Pure Lutetium (Lu) Duplex stainless steel High strength structural steel Pure Promethium (Pm) Pure Iron (Fe) Metal matrix composite aluminum Cobalt alloys (including Stellite) Pure Yttrium (Y) Gray cast iron Precipitation hardening stainless steel Pure Bismuth (Bi) Pure Holmium (Ho) Nickel copper Pure Nickel (Ni) Palladium alloys Pure Cobalt (Co) Cast austenitic stainless steel Gold alloys High-nickel gray cast iron Bismuth tin alloys Pure Uranium (U) Pure Gold (Au) Pure Samarium (Sm) Pure Erbium (Er) Nickel chromium silicon gray cast iron Tungsten alloys Beryllium alloys Manganese alloy steel Iron alloys Proprietary alloy steel White cast iron Austenitic cast iron with graphite Pure Thulium (Tm) Wrought copper nickel Ductile high-nickel cast iron Pure Lanthanum (La) Wrought high copper alloys Cast high copper alloys Wrought bronze Cast copper Wrought copper Cast copper nickel silver Austenitic stainless steel Cast bronze Wrought copper nickel silver Pure Barium (Ba)

pg 544


Appendix 6:

18/8/2008 4:32PM Plate # 0


Table A6.2 (continued) CTE

CTE 10 6/K

10 6/°F

Material

10 6/K

10 6/°F

Material

18 18 18–20 19 17–21 16–23

10 10 9.9–11 11 9.4–12 8.9–13

16–24 16–24 16–24

8.9–13 8.9–13 9.1–13

20 20–21 18–24

11 11–12 10–13

25–26 25–26 25–26 25–26 26 17–36 25–27 26 20–33 25–28 27–29 28 23–24

14–15 14–15 14–15 14–15 15 9.2–20 14–15 15 11–18 14–16 15–16 16 13–19

20–22 19–23 12–32 22 12–22 22 21–24 22–24 23 23 22–25

11–12 11–13 6.4–18 12 6.9–12 12 12–13 12–13 13 13 12–14

23–25 24–25 21–29 22–28 23–27

13–14 13–14 12–16 12–15 13–15

Cast copper nickel Pure Tellurium (Te) Silver alloys Pure Silver (Ag) Wrought brass 3xx.x series cast aluminum silicon+copper or magnesium 2xxx series wrought aluminum copper Zinc copper titanium alloys 6xxx series wrought aluminum magnesium silicon Pure Strontium (Sr) Cast brass 1xx.x series commercially pure cast aluminum 4xx.x series cast aluminum silicon 2xx.x series cast aluminum copper Pure Gallium (Ga) Manganese (Mn) 4xxx series wrought aluminum silicon Pure Calcium (Ca) 7xxx series wrought aluminum zinc 3xxx series wrought aluminum manganese 8xx.x series cast aluminum tin Unalloyed aluminum ingot 1xxx series commercially pure wrought aluminum 5xx.x series cast aluminum magnesium 7xx.x series cast aluminum zinc Tin lead Zinc aluminum Zinc copper aluminum

29 28–30 26–32 20–40 22–40 30–32 33–35 35 37–49 56 64 69–71 83 90 14–203 125 97–291

16 16–17 14–18 11–12 12–22 17–18 18–19 19 21–27 31 36 38–39 46 50 7.8–113 70 54–162

Cast magnesium aluminum zinc Pure Magnesium (Mg) Wrought magnesium aluminum zinc Cast magnesium aluminum manganese Cast magnesium rare earth Commercially pure tin Commercially pure magnesium Pure Ytterbium (Yb) Pure Indium (In) Lead tin solder Commercially pure or low-alloyed lead Tin silver 9xx.x series cast aluminum plus other elements Pure Lead (Pb) Pure Thallium (Tl) Magnesium alloys Unalloyed or low-alloy zinc 5xxx series wrought aluminum magnesium Pure Cadmium (Cd) Zinc copper Pure Europium (Eu) Pure Selenium (Se) Pure Lithium (Li) Pure Sulfur (S) Pure Sodium (Na) Pure potassium (K) Pure Rubidium (Rb) Pure Plutonium (Pu) Pure Phosphorus (P) Pure Cesium (Cs)

Source: Thermal Properties of Metals and Alloys, ASM, 2002

Table A6.3 Summary of thermal conductivity Thermal conductivity, approximate ranges at room temperature to 100 C (212 F), from lowest to highest value Thermal conductivity W/m K

2–3 0.2–5 3 5 5–8 8 8 6–11 8–11 6–12 8–10 6–14 4–17 11 11 11 11 11–13 9–15 13 13 13 13 13 11–16 10–17

Btu/ (h ft °F)

1–2 0.1–3 2 3 3–5 5 5 3–7 4–7 3–7 5–6 4–8 2–10 6 6 6 7 6–7 5–9 7 7 8 8 8 6–10 6–10

Thermal conductivity Material

Pure Tellurium (Te) Pure Selenium (Se) Pure Sulfur (S) Pure Technetium (Tc) Pure Plutonium (Pu) Permanent magnet iron alloy Pure Manganese (Mn) Beta and near beta titanium Pure Bismuth (Bi) Alpha beta titanium Pure Mercury (Hg) Cobalt chromium nickel tungsten alloys Alpha and near alpha titanium Pure Gadolinium (Gd) Pure Dysprosium (Dy) Pure Terbium (Tb) Pure Cerium (Ce) Cermet Nickel molybdenum alloys Austenitic cast iron with graphite Pure Praseodymium (Pr) Mischmetal Pure Lanthanum (La) Pure Samarium (Sm) Ferritic stainless steel Pure Yttrium (Y)

W/m K

13–15 14 14–16 25 15 15 11–21 16 16 16 16 1–33 17 17 12–24 18 15–22 15–22 33 11–28 8–35 13–30 22 19–25 20–25 14–31

Btu/ (h ft °F)

Material

7–9 8 8–9 15 8 9 6–12 9 9 9 10 0.6–19 10 10 7–14 11 9–13 9–13 19 6–16 5–20 7–17 13 11–15 11–15 8–18

Cobalt chromium tungsten alloys Pure Europium (Eu) Manganese alloy steel Cast martensitic stainless steel Pure Erbium (Er) Pure Promethium (Pm) Austenitic stainless steel Pure Scandium (Sc) Pure Gallium (Ga) Pure Holmium (Ho) Pure Lutetium (Lu) Cobalt alloys Pure Neodymium (Nd) Pure Thulium (Tm) Cast austenitic stainless steel Pure Barium (Ba) Pure Titanium (Ti) Low-alloy titanium Pure Rubidium (Rb) Uranium alloys Nickel with chromium, iron, molybdenum Duplex stainless steel Cast ferritic stainless steel Pure Rhenium (Re) Maraging steel Martensitic stainless steel

pg 545


18/8/2008 4:32PM Plate # 0


Table A6.3

(continued) Thermal conductivity

Thermal conductivity W/m K

Btu/ (h ft °F)

Material

W/m K

Btu/ (h ft °F)

13–32 22–24 19–27 24 8–40 17–42 22–29 17–34 22–30 20–32 24–29 18–36 13–42 25–30 20–36 24–34 10–50 16–45 23–45 34 13–55 34–35 29–41 23–50 11–63 38 39 36–43 40 27–54 35–57 25–58 37–48 31–54 39–46 34–54 44 8–83 12–81 46–57 47 45–50 46–50 47–52 29–70 20–80 50–51 29–72 51 18–86 53–54 38–71 52–57 54–57 58 42–75 44–75 48–71 56–65 52–74 67 60–76 21–116 69–73 71–76 68–80

8–19 13–14 11–16 14 5–23 10–24 13–17 10–20 13–17 12–19 14–17 11–21 8–24 15–17 12–21 14–20 6–22 9–26 13–26 20 7–32 19–20 17–24 13–28 6–37 22 22 21–25 23 16–31 20–33 14–33 21–28 18–31 23–27 20–29 26 5–48 7–47 27–33 27 26–29 27–29 27–30 17–40 12–46 29–30 17–42 30 11–50 30–31 22–41 30–33 31–33 34 24–43 26–43 28–41 32–38 30–43 39 35–44 12–67 40–42 41–44 39–47

Precipitation hardening stainless steel Pure Hafnium (Hf) Pure Zirconium (Zr) Pure Indium (In) Zirconium alloys Hot work tool steel Cast nickel-silver copper alloys Pure Antimony (Sb) White cast iron Cold work tool steel Shock-resisting tool steel Pure Cesium (Cs) Ductile cast iron Pure Uranium (U) High-speed tool steel Lead antimony Nickel steel Nickel copper alloys Cast copper-nickel Low-alloy lead Magnetically soft iron alloy Pure Lead (Pb) Pure Gallium (Ga) Chromium alloy steel Specialty stainless steel Pure Thorium (Th) Pure Ytterbium (Yb) Chromium molybdenum alloy steel Tin antimony Tantalum alloys Lead tin Ultrahigh strength steel Molybdenum alloy steel Pure Vanadium (V) Pure Thallium (Tl) Nickel chromium molybdenum alloy steel Nickel molybdenum alloy steel Nickel specialty alloys Cast bronze Tin lead alloys Pure Protactinium (Pa) Niobium alloys High-manganese carbon steel Resulfurized carbon steel Niobium alloys Gray cast iron High-strength low-alloy steel Mold steel Malleable cast iron Low-alloy nickel Pure Niobium (Nb) Nonresulfurized carbon steel Pure Tantalum (Ta) Tin silver alloys Pure Rubidium (Rb) Cast magnesium aluminum manganese Cast magnesium aluminum zinc Pure Rhenium (Re) Pure Germanium (Ge) Pure Tin (Sn) Pure Chromium (Cr) Pure Palladium (Pd) Wrought nickel-silver copper alloys Pure Platinum (Pt) Pure Lithium (Li) Pure Iron (Fe)

77 35–121 47–110 80 51–113 72–92 69–99 87 88 88 87–91 84–97 56–142 65–138 96–108 100–108 25–189 21–195 108–113 105–125 105–125 116 88–146 125–126 26–230 100–159 92–168

44 20–70 27–64 47 30–65 42–53 40–57 50 51 50 50–53 49–56 32–82 38–80 56–97 58–63 14–109 12–115 62–65 61–72 61–72 67 51–85 72–73 15–133 58–92 53–97

117–146 131–138 133–138 138–155 121–163 141–147 134–155 138–156 145–151 130–178 112–205

68–85 76–80 77–80 80–90 70–94 82–85 77–90 80–90 84–87 75–103 65–118

153–167 109–212 100–234 154–193

88–97 63–123 58–135 89–112

121–227 167–184 146–210 112–251 85–315 167–244

70–131 97–106 85–121 65–145 49–182 97–141

142–284

82–164

209–222

121–128

221–247 80–410 117–431 194–391 292–316 346–391 360–413 417–428 415–493 633–716

128–143 46–237 21–249 112–226 169–182 200–226 208–239 241–247 240–285 366–414

Material

Pure Thorium (Th) Carbide based material Wrought magnesium aluminum zinc Pure Indium (In) Cast rare earth magnesium alloys Pure Nickel (Ni) Pure Cobalt (Co) Pure Osmium (Os) Pure Rhodium (Rh) Pure Gallium (Ga) Pure Chromium (Ch) Pure Cadmium (Cd) Magnesium alloys Tungsten alloys (2xx.x series) Cast aluminum copper Potassium (K) Cast brass Wrought copper nickel Low-alloy zinc Pure Zinc (Zn) Zinc aluminum Pure Ruhenium (Ru) (5xx.x series) Cast aluminum magnesium Pure Calcium (Ca) Wrought bronze (7xx.x series) Cast aluminum zinc (3xx.x series) Cast aluminum silicon plus copper or magnesium Wrought magnesium zinc Pure Sodium (Na) Cast magnesium zinc (4xxx series) Wrought aluminum silicon (4xx.x series) Cast aluminum silicon Pure Iridium (Ir) Molybdenum (Mo) Pure Silicon (Si) Pure Rhodium (Rh) Pure Tungsten (W) (5xxx series) Wrought aluminum magnesium Pure Magnesium (Mg) Beryllium alloys Wrought brass (3xxx series) Wrought aluminum manganese (7xxx series) Wrought aluminum zinc (8xx.x series) Cast aluminum tin Pure Beryllium (Be) (2xxx series) Wrought aluminum copper Cast high copper alloys (1xxx series) Commercially pure wrought aluminum (6xxx series) Wrought aluminum magnesium silicon (1xx.x series) Commercially pure cast aluminum Pure Aluminum (Al) High copper alloys Molybdenum alloys Wrought coppers Pure Gold (Au) Cast copper Silver alloys Pure Silver (Ag) Silver tungsten Silver tungsten carbide

pg 546


Appendix 6:

18/8/2008 4:32PM Plate # 0

pg 547


Table A6.4 Thermal conductivities of carbon and low-alloy steels Conductivity, W/m K, at°C (°F) (a) AISI-SAE grade

1008 1008 1010 1015 1016 1018 1020 1022 1025 1026 1029 1030 1035 1037 1039 1040 1042 1043 1044 1045 1046 1050 1055 1060 1064 1070 1078 1078 1080 1086 1095 1117 1118 1141 1151 1522 1524 1526 1541 1548 1551 1561 1566 2515 4037 4130 4140 4145 4161 4427 4626 5132 5140 8617 8622 8627 8637 8822


0 (32)

Unknown Annealed Unknown Annealed Annealed Annealed Unknown Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Annealed Unknown Annealed Unknown Unknown Unknown Unknown Annealed Unknown Unknown Unknown Unknown Unknown Unknown Unknown Unknown Annealed Annealed Annealed Annealed Unknown Annealed Annealed Annealed Unknown Hardened and tempered Hardened and tempered Hardened and tempered Hardened and tempered Hardened and tempered Unknown Unknown Unknown Hardened and tempered Unknown Unknown Unknown Unknown Unknown

59.5 65.3(b) ... 51.9 51.9 51.9 51.9 51.9 51.9 51.9 51.9 ... ... ... ... ... 51.9 ... ... ... 51.2 51.2 51.2 50.5 51.2 49.9 47.8 49.6 50.5 49.9 ... 51.9(b) 51.5(b) ... ... 51.9 51.9 51.9 51.9 50.5 50.7 51.2 51.2 34.3(b) ...

100 (212)

200 (392)

300 (572)

400 (752)

500 (932)

600 (1112)

700 (1292)

800 (1472)

1000 (1832)

1200 (2192)

57.8 60.3 46.7 51.0 50.2 50.8 51.0 50.8 51.1 50.1 50.1 51.0 50.8 51.0 50.7 50.7 50.7 50.8 50.8 50.8 49.7 49.7 49.7 ... 49.7 48.4 48.2 48.1 ... 48.4 46.7 ... ... 50.5 50.5 50.1 50.1 50.1 50.1 49.0 49.3 49.7 49.7 ... 48.2

53.2 54.9 ... 48.9 47.6 48.9 48.9 48.8 49.0 48.4 48.4 ... ... ... ... ... 48.2 ... ... ... ... 46.8 ... 46.8 ... ... 45.2 ... 46.8 ... ... ... ... 47.6 47.6 48.4 48.4 48.4 48.4 48.3 48.4 ... ... ... 45.6

49.4 ... ... ... ... ... ... ... 46.1 ... ... ... ... ... ... ... 45.6 ... ... ... ... ... ... ... ... ... 41.4 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 38.2(c) ...

45.6 45.2 ... ... ... ... ... ... 42.7 ... ... ... ... ... ... ... 41.9 ... ... ... ... ... ... ... ... ... 38.1 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 39.4

41.0 ... ... ... ... ... ... ... 39.4 ... ... ... ... ... ... ... 38.1 ... ... ... ... ... ... ... ... ... 35.2 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

36.8 36.4 ... ... ... ... ... ... 35.6 ... ... ... ... ... ... ... 33.9 ... ... ... ... ... ... ... ... ... 32.7 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 33.9

33.1 ... ... ... ... ... ... ... 31.8 ... ... ... ... ... ... ... 30.1 ... ... ... ... ... ... ... ... ... 30.1 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

28.5 28.5 ... ... ... ... ... ... 26.0 ... ... ... ... ... ... ... 24.7 ... ... ... ... ... ... ... ... ... 24.3 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

27.6 27.6 ... ... ... ... ... ... 27.2 ... ... ... ... ... ... ... 26.8 ... ... ... ... ... ... ... ... ... 26.8 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

29.7 ... ... ... ... ... ... ... 29.7 ... ... ... ... ... ... ... 29.7 ... ... ... ... ... ... ... ... ... 30.1 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...

...

42.7

...

...

37.3

...

31.0

...

28.1

30.1

...

42.7

42.3

...

37.7

...

33.1

...

...

...

...

40.6

41.8(b)

...

...

...

...

...

...

...

...

...

...

42.7(b)

...

...

...

...

...

...

...

...

...

...

36.8(b) ... 48.6 ...

... 44.1 46.5 44.8

... ... ... ... ...

43.3 37.5(d) 37.5(d) 37.5(d) 37.5(d)

... ... 44.4 43.5

... ... 42.3 ...

... ... 38.5 37.7

... ... 35.6 ...

... ... 31.8 31.4

... ... 28.9 ...

... ... 26.0 ...

... ... 28.1 ...

... ... 30.1 ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

... ... ... ... ...

(a) To obtain conductivities in Btu/ft h F, multiply values in table by 0.5777893; to obtain conductivities in Btu in./ft2 h F, multiply values by 6,933472; to obtain conductivities in cal/cm s C, multiply values in table by 0.0023884. (b) Thermal conductivity at 20 C (68 F). (c) Thermal conductivity at 260 C (500 F). (d) Thermal conductivity at 50 C (120 F)


18/8/2008 4:32PM Plate # 0


Table A6.5 Summary of heat capacity Specific heat capacity, approximate ranges at room temperature to 100 C (212 F), from lowest to highest value Specific heat capacity J/kg K

113–117 92–142 116–117 104–140 121–129 126–130 112–133 130–132 130–131 129–134 128–138 136 134–142 137–140 111–167 130–150 134–151 130–159 131–159 144–148 147 150 154 160 162 165 168 171 182 182 188 190 193–195 150–239 195 195 196 201 204–213 219 213–239 205–238 205–247 230 202–267 236 201–272 235–239 238 201–280 236–246 234–251 244–248 229–264 233–264 254 246–276 268 285 281–289 298–300 308 310–322 310–328 325–331 272–418 330–380

Btu/lb °F)

0.027–0.028 0.022–0.034 0.028 0.025–0.033 0.029–0.031 0.030–0.031 0.027–0.032 0.031–0.032 0.031 0.031–0.032 0.031–0.033 0.032 0.032–0.034 0.033 0027–0.040 0.031–0.036 0.032–0.036 0.031–0.038 0.031–0.038 0.034–0.035 0.035 0.036 0.037 0.038 0.039 0.039 0.040 0.041 0.043 0.043 0.045 0.045 0.046–0.047 0.036–0.057 0.047 0.047 0.047 0.048 0.049–0.051 0.052 0.051–0.057 0.049–0.057 0.049–0.059 0.055 0.048–0.064 0.056 0.048–0.065 0.056–0.057 0.057 0.048–0.067 0.056–0.059 0.056–0.060 0.058–0.059 0.055–0.063 0.056–0.063 0.061 0.059–0.066 0.064 0.068 0.067–0.069 0.071–0.072 0.074 0.074–0.077 0.074–0.078 0.078–0.079 0.065–0.100 0.079–0.091

Specific heat capacity Material

Pure Thorium (Th) Tantalum alloys Pure Uranium (U) Pure Mercury (Hg) Pure Bismuth (Bi) Pure Gold (Au) Pure Lead (Pb) Pure Iridium (Ir) Pure Osmium (Os) Low alloyed lead Commercial gold Pure Platinum (Pt) Pure Tungsten (W) Pure Rhenium (Re) Uranium alloys Pure Thallium (Tl) Tungsten alloys Lead antimony alloys Pure Plutonium (Pu) Pure Hafnium (Hf) Pure Tantalum (Ta) Pure Lutetium (Lu) Pure Ytterbium (Yb) Pure Thulium (Tm) Pure Cerium (Ce) Pure Holmium (Ho) Pure Erbium (Er) Pure Dysprosium (Dy) Pure Europium (Eu) Pure Terbium (Tb) Pure Promethium (Pm) Pure Neodymium (Nd) Mischmetal Tin lead alloys Pure Lanthanum (La) Pure Praseodymium (Pr) Pure Samarium (Sm) Pure Tellurium (Te) Pure Antimony (Sb) Tin silver alloys Lead-tin alloys Pure Tin (Sn) Commercially pure tin Cadmium Pure Cesium (Cs) Pure Gadolinium (Gd) Niobium alloys Pure Silver (Ag) Pure Ruthenium (Ru) Molybdenum alloys Pure Palladium (Pd) Commercial silver Commercial palladium Pure Indium (In) Indium Semi-conductor Grade Pure Rhodium (Rh) Pure Molybdenum (Mo) Pure Niobium (Nb) Pure Barium (Ba) Pure Zirconium (Zr) Pure Yttrium (Y) Pure Strontium (Sr) Pure Germanium (Ge) Pure Arsenic (As) Gallium compounds Zinc alloys Pure Rubidium (Rb)

J/kg K

Btu/lb °F)

337–376 335–390 377 377.0 375–398 374–425

0.080–0.090 0.080–0.093 0.090 0.090 0.090–0.095 0.089–0.102

360–420 387–393 317–462 360–420 372–409 382–400 280–502 371–418 398 377–419 377–420 383–418 377–439 389–420 335–481 376–440 352–473 285–545

0.086–0.100 0.092–0.094 0.076–0.110 0.086–0.100 0.089–0.098 0.091–0.096 0.067–0.120 0.089–0.100 0.095 0.090–0.100 0.090–0.100 0.091–0.100 0.090–0.105 0.093–0.100 0.080–0.115 0.090–0.105 0.084–0.113 0.068–0.130

335–502 418–421 377–565 414–452 368–502 439 420–460

0.080–0.120 0.100–0.101 0.090–0.135 0.099–0.108 0.088–0.120 0.105 0.100–0.110

448 420–481 431–477 456 393–525 444–473 460 460 460

0.107 0.100–0.115 0.103–0.114 0.109 0.094–0.125 0.106–0.113 0.110 0.110 0.110

460 460 460 460 420–502 460–462 453–473 442–487 440–494 470 440–502 460–486 461–486 461–486‘ 448–502 418–540 460–502 430–532 460–502 486 481–502 475–508 460–527

0.110 0.110 0.110 0.110 0.100–0.120 0.110 0.108–0.113 0.106–0.116 0.105–0.118 0.112 0.105–0.120 0.110–0.116 0.110–0.116 0.110–0.116 0.107–0.120 0.100–0.129 0.110–0.120 0.103–0.127 0.110–0.120 0.116 0.115–0.120 0.113–0.121 0.110–0.126

(continued)

Material

Cast beryllium copper nickel alloys Cast brass Cast copper-nickels Cast nickel-silvers Wrought copper-nickels Cobalt chromium nickel tungsten alloys Wrought high-copper alloys Pure Copper (Cu) Pure Selenium (Se) Wrought bronze Gallium (Ga) Pure Zinc (Zn) Zirconium alloys Wrought brass Beryllium copper Wrought nickel silvers Wrought coppers Unalloyed or low-alloy zinc Nickel molybdenum alloys Cast high copper alloys Nickel molybdenum alloy steel Cast bronze Cobalt alloys Nickel with chromium and/or iron, molybdenum Precipitation hardening stainless steel Cobalt chromium tungsten alloys Austenitic stainless steel Pure Cobalt (Co) Nickel steel Chrome-nickel-iron superalloy Air-hardening medium-alloy cold work tool steel Ultrahigh-strength steel Maraging high strength steel Martensitic stainless steel Molybdenum alloy steel Zinc aluminum alloys Pure Iron (Fe) Cast martensitic stainless steel Mold steel High-carbon high-chromium cold work tool steel Molybdenum high-speed tool steel Gray cast iron Cold work tool steel Water-hardening tool steel Nonresulfurized carbon steel Hot work tool steel Pure Nickel (Ni) Pure Chromium (Cr) Chromium alloy steel Permanent magnet iron alloys Duplex stainless steel High strength low alloy steel (HSLA) High-manganese carbon steel Austenitic cast iron with graphite Chromium molybdenum alloy steel Nickel copper alloys Austenitic cast stainless steel Commercially pure or low-alloy nickel Ductile cast iron High strength structural steel Resulfurized carbon steel Pure Manganese (Mn) Malleable cast iron

pg 548


Appendix 6:

18/8/2008 4:32PM Plate # 0

pg 549


Table A6.5 (continued) Specific heat capacity

Specific heat capacity Btu/lb °F)

Material

500 498–506 420–586 460–572 370–674 460–586 510–553 490–580 513–599 515–597 448–670

0.119 0.119–0.121 0.100–0.140 0.110–0.137 0.088–0.161 0.110–0.140 0.122–0.132 0.117–0.139 0.123–0.143 0.123–0.143 0.107–0.160

561–567 623–655 691 712–728 733 741 770 864

0.134–0.135 0.149–0.156 0.165 0.170–0.174 0.175 0.177 0.184 0.206

816–963

0.195–0.230

893–920

0.213–0.220

895–933 875–963 920

0.214–0.223 0.209–0.230 0.220

879–963

0.210–0.230

Cast ferritic stainless steel Pure Vanadium (V) Ferritic stainless steel Alpha and near alpha titanium alloys Alpha beta titanium alloys Cast stainless steels Manganese alloy steel Beta and near beta titanium alloys Pure Titanium (Ti) Unalloyed or low-alloy titanium Nickel chromium molybdenum alloy steel Pure Scandium (Sc) Pure Calcium (Ca) Carbon Pure Silicon (Si) Pure Sulfur (S) Pure Phosphorus (P) Pure Potassium (K) (4xxx series) Wrought aluminum silicon (2xxx series) Wrought aluminum copper (3xxx series) Wrought aluminum manganese Pure Aluminum (Al) (2xx.x series) Cast aluminum copper (8xxx series) Wrought aluminum+other elements (5xxx series) Wrought aluminum magnesium

J/kg K

Btu/lb °F)

Material

887–963

0.212–0.230

800–1050 795–1050 900–963

0.191–0.251 0.190–0.251 0.215–0.230

960–963

0.229–0.230

963 963

0.230 0.230

963 963 966–1050 1010–1050 962–1109

0.230 0.230 0.231–0.251 0.241–0.251 0.230–0.265

1000–1080

0.239–0.258

1025–1064 1050 1060 1222–1234 1285–1620 1500–1630 1886–2070 3300–3515

0.245–0.254 0.251 0.253 0.292–0.295 0.307–0.387 0.358–0.389 0.450–0.494‘ 0.788–0.840

(6xxx series) Wrought aluminum magnesium silicon Cast magnesium aluminum zinc (7xxx series) Wrought aluminum zinc (1xxx series) Commercially pure wrought aluminum, 99.00% or greater (3xx.x series) Cast aluminum silicon+copper or magnesium (4xx.x series) Cast aluminum silicon (5xx.x series) Cast aluminum magnesium (7xx.x series) Cast aluminum zinc (8xx.x series) Cast aluminum tin Cast magnesium rare earth alloys Commercially pure magnesium Wrought magnesium aluminum zinc alloys Cast magnesium aluminum manganese Wrought magnesium zinc alloys Cast magnesium zinc Pure Magnesium (Mg) Pure Sodium (Na) Pure Boron (B) Beryllium alloys Pure Beryllium (Be) Pure Lithium (Li)

J/kg K

Table A6.6 Specific heats of carbon and low-alloy steels Mean apparent specific heat, J/Kg K, at °C (°F) AISISAE grade

1008 1010 1015 1016 1017 1018 1020 1025 1030 1035 1040 1042 1045 1050 1060 1070 1078 1086 1095 1117 1140 1151 1522 1524


Annealed Unknown Annealed Annealed Unknown Annealed Unknown Annealed Annealed Annealed Annealed Annealed Annealed Annealed Unknown Unknown Annealed Unknown Unknown Unknown Unknown Unknown Annealed Annealed

50–100 150–200 200–250 250–300 300–350 350–400 450–500 (122–212) (302–392) (392–482) (482–572) (572–662) (662–752) (842–932)

481 450 486 481 481(a) 486 486 486 486 486 486 486 486 486 502 490 490 500 461(b) 481 461(b) 502(b) 486 477

519 500 519 515 ... 519 519 519 519 519 519 515 519 519 544 532 532 532 ... ... ... ... 519 511

536 520 ... ... ... ... ... 532 ... ... ... 528 ... ... ... ... 548 ... ... ... ... ... ... 528

553 535 ... ... ... ... ... 557 ... ... ... 548 ... ... ... ... 565 ... ... ... ... ... ... 544

574 565 ... ... ... ... ... 574 ... ... ... 569 ... ... ... ... 586 ... ... ... ... ... ... 565

595 590 599 595 ... 599 599 599 599 586 586 586 586 590 ... ... 607 ... ... ... ... ... 599 590

662 650 ... ... ... ... ... 662 . . .. ... ... 649 ... ... ... ... 670 ... ... ... ... ... ... 649

550–600 650–700 700–750 750–800 850–900 (1022– (1202– (1292– (1382– (1562– 1112) 1292) 1382) 1472) 1652)

754 730 ... ... ... ... ... 749 ... ... ... 708 ... ... ... ... 712 ... ... ... ... ... ... 741

867 825 ... ... ... ... ... 846 ... ... ... 770 ... ... ... ... 770 ... ... ... ... ... ... 837

1105 ... ... ... ... ... ... 1432 ... ... ... 1583 ... ... ... ... 2081 ... ... ... ... ... ... 1449

875 ... ... ... ... ... ... 950 ... ... ... 624 . . .. ... ... ... 615 ... ... ... ... ... ... 821

846 ... ... ... ... ... ... ... ... ... ... 548 ... ... ... ... ... ... ... ... ... ... ... 536

(continued) C (75–200 F). (b) Specific heat at 20–100 C (68–212 F). (c) Specific heat at 20–200 C (68–392 F). (d) Value presented is mean value of

(a) Specific heat at 25–95 temperatures between 20 C (68 F) and the higher of the cited temperatures. (e) Specific heat at 10–25 C (50 80 F). (f) Average specific heat from 25 540 C (80 1000 F)


18/8/2008 4:32PM Plate # 0

pg 550


Table A6.6

(continued) Mean apparent specific heat, J/Kg K, at °C (°F)

AISISAE grade


1561 4032 4130

Annealed Unknown Hardened and tempered Hardened and tempered Unknown Normalized and tempered Unknown Unknown Hardened and tempered Unknown Unknown Unknown

4140

4142 4626

4815 5132 5140

8115 8617 8637

50–100 150–200 200–250 250–300 300–350 350–400 450–500 (122–212) (302–392) (392–482) (482–572) (572–662) (662–752) (842–932)

550–600 650–700 700–750 750–800 850–900 (1022– (1202– (1292– (1382– (1562– 1112) 1292) 1382) 1472) 1652)

486 ... 477

519 461(c) 515

... ... ...

... ... 544

... ... ...

... ... 595

... ... 657

... ... 737

... ... 825

... ... ...

... ... 833

... ... ...

...

473(d)

...

...

...

519(d)

...

561(d)

...

...

...

...

... 335(e)

502(c) ...

... ...

... ...

... ...

... ...

... ...

... ...

... 615(f)

... ...

... ...

... ...

481(b) 494 452(d)

... 523 473(d)

... 536 ...

... 553 ...

... 574 ...

... 595 519(d)

... 657 ...

... 741 561(d)

... 1499 ...

... 934 ...

... 574 ...

461(b) 481(a) ...

... ... 502(c)

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... 837 ...

... ... ...

(a) Specific heat at 25–95 C (75–200 F). (b) Specific heat at 20–100 C (68–212 F). (c) Specific heat at 20–200 C (68–392 F). (d) Value presented is mean value of temperatures between 20 C (68 F) and the higher of the cited temperatures. (e) Specific heat at 10–25 C (50 80 F). (f) Average specific heat from 25 540 C (80 1000 F)



21/8/2008 4:56PM Plate # 0

pg 551


APPENDIX 7

AISI to Non-AISI Steel Cross Reference Table A7.1 Cross-references of standard SAE carbon and low-alloy steels to selected chemically similar steels Specifications established by standards organizations from the Federal Republic of Germany, Japan, the United Kingdom, France, Italy, and Sweden. United states (SAE)

Fed. R. of Germany (DIN)

Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon steels 1005

1.0288, D5-2 1.0303, QSt32-2 1.0312, D5-1 1.0314, D6-2 1.0393, ED3 1.0394, ED4 1.1012, RFe120

...

970 015A03

1006

1.0311, D7-1 1.0313, D8-2 1.0317, RSD4 1.0321, St23 1.0334, StW23 1.0335, StW24 1.0354, St14Cu3 1.0391, EK2 1.0392, EK4 1.1009, Ck7

...

970 030A04 970 040A04 970 050A04

1008

1.0010, D9 1.0318, St28 1.0320, St22 1.0322, USD8 1.0326, RSt28 1.0330, St2, St12 1.0333, St3, St13 1.0331, RoSt2 1.0332, StW22 1.0336, Ust4, Ust14 1.0337, RoSt14 1.0344, St12Cu3 1.0347, RRSt13 1.0357, USt28 1.0359, RRSt23 1.0375, Feinstblech T57, T61, T65, T70 1.0385, Weissblech T57, T61, T65, T70 1.0744, 6P10 1.0746, 6P20 1.1116, USD6

G34445 STKM11A (11A)

1449 3CR 1449 3CS 1449 3HR 1449 3HS 1717 ERW101 3606 261

1010

1.0204, UQSt36 1.0301, C10 1.0328, USD10 1.0349, RSD9 1.1121, Ck10 1.1122, Cq10

G4051 S10C G4051 S9Ck

1449 40F30, 43F35, A33-101 AF34 46F40, 50F45, CC10 60F55, 68F62, C10 75F70 (available in HR, HS, CS conditions) 1449 4HR, 4HS, 4CR, 4CS 970 040A10 (En2A, En2A/1, En2B) 970 045A10, 045M10 (En32A) (continued)

...

5598 3CD5

1160

A35-564 XC6FF

5598 3CD6 5771 C8

1147 1225

A35-551 XC10 XC6 XC6FF

5598 3CD8

1142 1146

5331 C10 6403 C10 7065 C10 7846 C10 5598 ICD10 5598 3CD12 5771 C12

1232 1265 1311

7356 CB10FF, CB10FU


21/8/2008 4:57PM Plate # 0

pg 552


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon steels (continued) 970 050A10 970 060A10 980 CEW1 1012

1.0439, RSD13

1013

1.0036, USt37-2 1.0037, St37-2 1.0038, RSt37-2 1.0055, Ust34-1 1.0057, RSt34-1 1.0116, St37-3 1.0218, RSt41-2 1.0219, St41-3 1.0307, StE210-7 1.0309, St35.4 1.0315, St37.8 1.0319, RRStE210.7 1.0356, TTSt35 1.0417 1.0457, StE240.7

1015

1.0401, C15 1.1132, CQ15 1.1135, Ck16A1 1.1140, Cm15 1.1141, Ck15 1.1144 1.1148, Ck15AI

1016

1.0419, RSt44.2 1.0467, 15Mn3 1.0468, 15Mn3A1 1.1142, GS-Ck16

1017

...

1018

1.0453, C16.8

G4051 S12C

...

1449 12HS, 12CS 1501 141-360 970 040A12 (En2A, En2A/1, En2B) 970 050A12 970 060A12

A33-101 AF37 A-35 551 XC12 C12

3059 360 3061 360 3603 360

A35-551 XC12 CC12

5869 Fe360-1KG, Fe360-2KW 6403 Fe35-2 7070 Fe34CFN 7091 Fe34

1233 1234 1330

G4051 F15Ck G4051 S15C

970 040A15 970 050A15 970 060A15 970 080A15, 080M15 970 173H16

XC15

5331 C16 7065 C16 7356 CB15 7846 C15

1370

...

3059 440 3606 440 970 080A15, 080M15 970 170H15 970 17H16

...

G4051 S17C

...

1019

...

...

1020

1.0402, C22 1.0414, D20-2

G4051 S20C G4051 S20CK

1.0427, C22.3 1.0460, C22.8 1.1149, Cm22 1.1151, Ck22

...

1332 1431

...

1370 2101

1449 17HS, 17CS 970 040A17 970 050A17 970 060A17

A35-551 XC18 A35-552 XC18 A35-566 XC18 A35-553 XC18S A35-554 XC18S

...

1312

970 080A17

A33-101 AF42 C20

...

...

...

...

...

970 040A20 970 050A20 (En2C, En2D) 970 060A20

A35-551 XC18 A35-552 XC18

5598 1CD20 5598 3CD20

A35-566 XC18 A35-553 C20 A35-553 XC18S A35-554 XC18S CC20

6922 C21 7356 CB20FF

... 1450

1021

...

...

970 070M20 970 080A20

A35-551 21B3 A35-552 21B3 A35-553 21B3 A35-557 21B3 A35-566 21B3

5332 C20 7065 C20

...

1022

1.0432, C21 1.0469, 21Mn4 1.0482, 19Mn5 1.1133, 20Mn5, GS-20Mn5

...

3111 Type 9 970 120M19 970 170H20

A35-551 20MB5 A35-552 20MB5 A35-553 20MB5 A35-556 20MB5

5771 20Mn4

...

(continued)


Appendix 7:

21/8/2008 4:57PM Plate # 0

pg 553

AISI to Non-AISI Steel Cross Reference / 553

Table A7.1 (continued) United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon steels (continued) 1.1134, Ck19

A35-557 20MB5 A35-566 20MB5 A35-566 20M5

1023

1.1150, Ck22.8 1.1152, Cq22

G4051 S22C

1025

1.0406, C25 1.0415, D25-2, D26-2 1.1158, Ck25

G4051 S25C

1026

1.1155, GS-Ck25 1.1156, GS-Ck24

1029

1.0562, 28Mn4

...

...

1449 2HS, 22CS 970 040A22 (En2C, En2D) 970 050A22 970 060A22 970 080A22 ...

A35-552 XC25 A35-566 XC25 ...

970 070M26 970 080A25 970 080A27

G3445 STKM15A (15A), STKM15C (15C) G4051 S28C

970 060A27 970 080A27 (En5A)

A33-101 AF50 CC28

5332 C20 7065 C20

...

5598 ICD25 5598 3CD25

...

7845 C25 7847 C25

...

...

...

C30

1030

1.0528, C30 1.0530, D30-2 1.1178, Ck30 1.1179, Cm30 1.1811, G-31Mn4

G4051 S30C

1449 30HS, 30CS 970 060A30 970 080A30 (En5B) 970 080M30 (En5)

A35-552 XC32 A35-553 XC32

5332 C30 6403 C30 7065 C30 7845 C30 7874 C30 5598 3CD30 6783 Fe50-3 7065 C31

...

1035

1.0501, C35 1.0516, D35-2 1.1172, Cq35 1.1173, Ck34 1.1180, Cm35 1.1181, Ck35

G4051 S35C

1717 CDS105/106 970 060A35 970 080A32 (En5C) 970 080A35 (En8A) 980 CFS6

A33-101 AF55 A35-553 C35 A35-553 XC38 A35-554 XC38 XC35 XC38TS C35

5333 C33 5598 1CD35 5598 3CD35 7065 C35 7065 C36 7847 C36 7356 CB35

1550 1572

1037

1.0520, 31Mn4 1.0561, 36Mn4

G4051 S35C

3111 type 10 970 080M36 970 170H36

1038

No international equivalents

1039

1.1190, Ck42A1

1040

1.0511, C40 1.0541, D40-2 1.1186, Ck40 1.1189, Cm40

1042

1043

...

...

970 060A40 970 080A40 (En8C) 970 080M40 (En8) 970 170H41

40M5 A35-552 XC38H2 A35-553 38MB5 A35-556 38MB5 A35-557 38MB5 A35-557 XC38H2 XC42, XC42TS

G4051, S40C

1287 1449 40HS, 40CS 3146 Class 1 Grade C 3146 Class 8 970 060A40 970 080A40 (En8C) 970 080M40 (En8)

A33-101 AF60 C40

1.0517, D45-2

G4051 S43C

970 060A42 970 080A42 (En8D)

A35-552 XC42H1 A35-553 C40 CC45 XC42, XC42TS

1.0558, GS-60.3

G4051 S43C

970 060A42 970 080A42 (En8D) 970 080M46

A35-552 XC42H2

(continued)

...

...

...

...

5598 ICD40 5598 3CD40 6783 Fe60-3 6923 C40 7065 C40 7065 C41 ...

7847 C43

...

...

...


21/8/2008 4:57PM Plate # 0

pg 554


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon steels (continued) 1044

1.0517, D45-2

1045

1.0503, C45 1.1184, Ck46 1.1191, Ck45, GS-Ck45 1.1192, Cq45 1.1194, Cq45 1.1201, Cm45 1.1193, Cf45

1046

1049

1.0503, C45 1.0519, 45MnAl 1.1159, GS-46Mn4 ...

... G4051 S45C G5111 SCC5

...

... 970 060A47 970 080A47 970 080M46

...

...

A33-101 AF65 A35-552 XC48H1 A35-553 XC45

3545 C45 5332 C45 7065 C45

A35-554 XC48 XC48TS C45

7845 C45 7874 C45 5598 1CD45 5598 3CD45 7065 C46 7847 C46 ...

3100 AW2 970 080M46

45M4TS A35-552 XC48H1 A35-552 XC48H2 XC48TS

G3445 STKM17A (17A) G3445 STKM17C (17C)

970 060A47

A35-552 XC48H1

6403 C48

970 080A47

A35-554 XC48

7847 C48

... 1672

...

...

XC48TS 1050

1.0540, C50 1.1202, D53-3 1.1206, Ck50 1.1210, Ck53 1.1213, Cf53 1.1219, Cf54 1.1241, Cm50

G4051 S50C G4051 S53C

1549 50HS 1549 50CS 970 060A52 970 080A52 (En43C) 970 080M50 (En43A)

A35-553 XC50

5332 C50 7065 C50 7065 C51 7845 C50 7874 C50 5598 ICD50 5598 3CD50 6783 Fe70-3 7847 C53

1674

1053

1.1210 Ck53 1.1213 Cf53 1.1219 Cf54

G4051 S53C

970 080A52 (En43C)

52M4TS A35-553 XC54

7847 C53

1674

1055

1.0518, D55-2 1.0535, C35 1.1202, D53-3 1.1203, Ck55 1.1209, Cm55 1.1210, Ck53 1.1213, Cf53 1.1219, Cf54 1.1220, D55-3 1.1820, C55W

G4051 S53C G4051 S55C

3100 AW3 970 060A57 970 070M55 970 080A52 (En43C) 970 080A57

A33-101 AF70 A35-552 XC55H1 A35-552 XC55H2 A35-553 XC54 XC55 C55

5598 3CD55 7065 C55 7845 C55 7874 C55 7065 C56 7847 C53

1059

1.0609, D58-2 1.0610, D60-2 1.0611, D63-2 1.1212, D58-3 1.1222, D63-3 1.1228, D60-3

970 060A62

A35-553 XC60

1060

1.0601, C60 1.0642, 60Mn3 1.1221, Ck60 1.1223, Cm60 1.1740, C60W

1449 60HS 1449 60CS 970 060A57 970 080A57

A35-553 XC60

3545 C60 7064 C60 7065 C60 7845 C60 7874 C60 5598 3CD60 7065 C61

1678

1064

1.0611, D63-2 1.0612, D65-2 1.0613, D68-2 1.1222, D63-3 1.1236, D65-3

...

970 060A62 970 080A62 (En43D)

...

5598 3CD65

...

1065

1.0627, C68 1.0640, 64Mn3 1.1230, Federstahldraht FD

...

970 060A67 970 080A67 (En43E)

...

G4051 S58C

(continued)

XC65

...

...

...

...

...


Appendix 7:

21/8/2008 4:57PM Plate # 0

pg 555




Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon steels (continued) 1.1233 1.1240, 65Mn4 1.1250, Federstahldraht VD 1.1260, 66Mn4 1069

1.0615, D70-2 1.0617, D73-2 1.0627, C68 1.1232, D68-3 1.1237 1.1249, Cf70 1.1251, D70-3 1.1520, C70W1 1.1620, C70W2

...

1070

1.0603, C67 1.0643, 70Mn3 1.1231, Ck67

...

1449 70HS, 70CS 970 060A72 970 070A72 (En42) 970 080A72

XC70

3545 C70

1770

1074

1.0605, C75 1.0645, 76Mn3 1.0655, C74 1.122, D73-3

...

970 070A72 (En42) 970 080A72

A35-553 XC75 XC70

3545 C75 7064 C75

1774

1075

1.0614, D75-2 1.0617, D73-2 1.0620, D78-2 1.1242, D73-3 1.1252, D78-3 1.1253, D75-3

...

A35-553 XC75 XC70

3545 C75 7064 C75 5598 3CD70 5598 3CD75

...

1078

1.0620, D78-2 1.0622, D80-2 1.0626, D83-2 1.1252, D78-3 1.1253, D75-3 1.1255, D80-3 1.1262, D83-3 1.1525, C80W1

970 060A78

XC80

5598 3CD80

...

1080

1.1259 80Mn4 1.1265 D85-2

...

1449 80HS, 80CS 970 060A78 970 060A83 970 070A78 970 080A78 970 080A83

XC80

5598 3CD80 5598 3CD85

...

1084

1.1830, C85W

...

970 060A86 970 080A86

XC85

G4801 SUP3

...

...

A35-553 XC68 XC70

...

...

... ...

...

... ...

1085

1.0647, 85Mn3 1.1273, 90Mn4 1.1819, 90Mn4

...

970 080A83

1086

1.0616, C85, D85-2 1.0626, D83-2 1.0628, D88-2 1.1262, D83-3 1.1265, D85-3 1.1269, Ck85 1.1272, D88-3

...

970 050A86

A35-553 XC90

5598 3CD85 5598 3CD90

...

1090

1.1273, 90Mn4 1.1819, 90Mn4 1.1282, D95S3

...

1449 95HS 1449 95CS 970 060A96

...

3545 C90 7064 C90 5598 3CD95

...

1095

1.0618, D95-2 1.1274, Ck101 1.1275, Ck100 1.1282, D95S3 1.1291, MK97 1.1545, C105W1 1.1645, C105W2

G4801 SUP4

1449 95HS 1449 95CS 970 060A99

(continued)

A35-553 XC100

3545 C100 7064 C100

1870


21/8/2008 4:57PM Plate # 0

pg 556


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon-manganese steels 1513

1.0424, Schiffbaustahl CS:DS 1.0479, 13Mn6 1.0496, 12Mn6 1.0513, Schiffbaustahl A32 1.0514, Schiffbaustahl B32 1.0515, Schiffbaustahl E32 1.0549 1.0579 1.0583, Schiffbaustahl A36 1.0584, Schiffbaustahl D36 1.0589, Schiffbaustahl E36 1.0599 1.8941, QStE260N 1.8945, QStE340N 1.8950, QStE380N

...

12M5

1449 40/30 HS 1449 40/30 CS 1453 A2

A33-101 AF50-S A35-501 E35-4 A35-501 E36-2

2772 150M12

A35-501 E36-3

970 130M15 970 130M15 (En201)

1.0471, 21MnS15 1.0529, StE350-Z2 1.1120, GS-20Mn5 1.1138, GS-21Mn5 1.1169, 20Mn6 1.8970, StE385.7 1.8972, StE415.7 1.8978 1.8979

G4106 SMn21

1503 221-460 1503 223-409 1503 224-490 3146 CLA2 980 CFS7

1524

1.0499, 21Mn6A1 1.1133, 20Mn5, GS-20Mn5 1.1160, 22Mn6

G4106 SMn21 G5111 SCMn1

1456 Grade A 970 150M19 (En14A, En14B) 970 175H23 980 CDS9, CDS10

...

...

970 120M28

A35-551 20MB5 A35-552 20M5 A35-556 20M5 A35-552 20MB5 A35-553 20MB5 A35-556 20MB5 A35-557 20MB5 A35-566 20MB5 ...

...

1.0412, 27MnSi5 1.1161, 26Mn5 1.1165, 30Mn5 1.1165, GS-30Mn5 1.1170, 28Mn6

G5111 SCMn2

1453 A3 1456 Grade B1, Grade B2 3100 A5 3100 A6 970 150M28 (En14A, En14B)

1536

1.0561, 36Mn4 1.1165, 30Mn5 1.1165, GS-30Mn5 1.1166, 34Mn5 1.1167, 36Mn5, GS-36Mn5 1.1813, G-35Mn5

G4052 SMn1H G4052 SMn433H G4106 SMn1 G4106 SMn433 G5111 SCMn2

1045 3100 A5, A6 970 120M36 (En15B) 970 150M36 (En15)

A35-552 32M5 A35-552 38MB5 A35-553 38MB5 A35-556 38MB5 A35-557 38MB5

1541

1.0563, E

G4106 SMn2, SMn438 G4052 SMn2H, SMn438H G4106 SMn3, SMn443 G4052 SMn3H, SMn443H G5111 SCMn5

970 135M44

40M5

970 150M40

45M5

1.1127, 36Mn6 1.1168, GS-40Mn5

1548

1.1128, 46Mn5 1.1159, GS-46Mn4

4010 FeG52 6930 20Mn6 7660 Fe510

...

A35-566 25MS5

1527

1.0564, N-80

...

970 125A15

1522

1526

...

1449 40/30 HR

2165 2168

...

2130 4010 FeG60 7874 C28Mn

...

4010 FeG60

...

G5111 SCMn3 ...

2120 2128

A35-552 40M6

...

...

...

...

...

1551

1.0542, StSch80

...

...

24M4TS

...

...

1552

1.0624, StSch90B 1.1226, 52Mn5

...

...

55M5

...

...

(continued)


Appendix 7:

21/8/2008 4:57PM Plate # 0

pg 557




Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Carbon-manganese steels (continued) 1561

1.0908, 60SiMn5

...

...

...

...

...

1566

1.1233 1.1240, 65Mn4 1.1260, 66Mn7

...

...

...

...

...

...

A35-562 10F1

...

...

Resulfurized carbon steels 1108

1.0700, U7S10 1.0702, U10S10

G4804 SUM12

1110

1.0703, R10S10

G4804 SUM11

1117

...

G4804 SUM31

970 210A15 970 210M17 (En32M) 970 214A15 970 214M15 (En202)

1118

...

...

970 214M15 (En201)

1137

...

G4804 SUM41

970 212M36 (En8M) 970 216M36 (En15AM) 970 225M36

35MF4 A35-562 35MF6

1139

1.0726, 35S20

...

970 212A37 (En8BM) 970 212M36 (En8M) 970 216M36 (En15AM) 970 225M36

35MF4 A35-562 35MF6

...

1957

1140


1141

...

G4804 SUM42

970 212A42 (En8DM) 970 216A42

A35-562 45MF4

...

...

1144

1.0727, 45S20

G4804 SUM43

970 212A42 (En8DM) 970 212M44 970 216M44 970 225M44 970 226M44

A35-562 45MF6

970 212M44

45MF4

...

1146

1.0727, 45S20

...

1151

1.0728, 60S20 1.0729, 70S20

...

...

...

...

...

...

...

...

...

...

...

... 4838 CF35SMn10

4838 CF44SMn28

... ...

1973

...

...

...

1973

Resulfurized/rephosphorized carbon steels 1211


1212

1.0711, 9S20 1.0721, 10S20 1.1011, RFe160K

G4804 SUM21

1213

1.0715, 9SMn28 1.0736, 9SMn36 1.0740, 9SMn40

G4804 SUM22

1215

1.0736, 9SMn36

G4804 SUM23

12L14


...

10F2 12MF4 S200

4838 10S20 4838 10S22 4838 CF9S22

970 220M07 (En1A) 970 230M07 970 240M07 (En1A)

A35-561 S250 S250

4838 CF9SMn28 4838 CF9SMn32

1912

970 240M07 (En1B)

A35-561 S300

4838 CF9SMn32 4838 CF9SMn36

...

Alloy steels 1330 1335 1340 1345

4023 4024

No international equivalents 1.5069, 36Mn7 1.5223, 42MnV7 1.0625, StSch90C 1.0912, 46Mn7 1.0913, 50Mn7 1.0915, 50MnV7 1.5085, 51Mn7 1.5225, 51MnV7 1.5416, 20Mo3 1.5416, 20Mo3

... ... ...

... ... ...

... ... ...

... ... ...

... ... ...

... ...

... ...

... ...

... ...

... ...

(continued)


21/8/2008 4:57PM Plate # 0

pg 558


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Alloy steels (continued) 4027

1.5419, 22Mo4

...

4028

...

...

...

...

...

...

... ...

970 605A32 970 605H32 970 605M30 970 605M36 (En16)

...

...

...

4032

1.5411

4037

1.2382, 43MnSiMo4 1.5412, GS-40MnMo4 3 1.5432, 42MnMo7

...

3111 Type 2/1 3111 Type 2/2 970 605A37 970 605H37

...

...

...

4042

1.2382, 43MnSiMo4 1.5432, 42MnMo7

...

...

...

...

...

4047


4118

1.721, 23CrMoB4 1.7264, 20CrMo5

...

7846 18CrMo4

...

4130

...

G5111 SCMnM3

...

970 605M30

G4052 SCM15H G4105 SCM21H G4052 SCM418H G4105 SCM418H

970 708H20 970 708M20

G4105 SCM1 G4105 SCM432 G4105 SCM2 G4105 SCM430 G4106 SCM2

1717 CDS110 970 708A30

A35-552 30CD4 A35-556 30CD4 A35-557 30CD4

30CrMo4 6929 35CrMo4F 7356 34CrMo4KB 7845 30CrMo4 7874 30CrMo4

2233

4135

1.2330, 35CrMo4 1.7220, 34CrMo4 1.7220, GS-34CrMo4 1.7226, 34CrMoS4 1.7231, 33CrMo4

G4054 SCM3H G4054 SCM435H G4105 SCM1 G4105 SCM432 G4105 SCM3 G4105 SCM435

970 708H37 970 708H37

35CD4 A35-552 35CD4 A35-553 35CD4 A35-556 35CD4 A35-557 34CD4

5332 35CrMo4 6929 35CrMo4F 7356 34CrMo4KB 7845 35CrMo4 7874 35CrMo4

2234

4137

1.7225, GS-42CrMo4

G4052 SCM4H G4052 SCM440H G4105 SCM4

3100 type 5 970 708A37 970 708H37 970 709A37

40CD4 42CD4 A35-552 38CD4 A35-557 38CD4

5532 40CrMo4 5333 38CrMo4 7356 38CrMo4KB

...

4140

1.3563, 43CrMo4 1.7223, 41CrMo4

G4052 SCM4H G4052 SCM440H

3100 Type 5 4670 711M40

3160 G40CrMo4 5332 40CrMo4

1.7225, 42CrMo4

G4103 SNCM4

970 708A40

1.7225, GS-42CrMo4

G4105 SCM4

970 708A42 (En19C)

1.7227, 42CrMoS4

G4105 SCM440

970 708H42

40CD4 A35-552 42CD4, 42CDTS A35-553 42CD4, 42CDTS A35-556 42CD4, 42CDTS A35-557 42CD4, 42CDTS

40CD4 A35-552 42CD4, 42CDTS A35-553 42CD4, 42CDTS A35-556 42CD4, 42CDTS A35-557 42CD4, 42CDTS

7845 42CrMo4 7874 42CrMo4

2244

7845 42CrMo4 7847 41CrMo4 7874 42CrMo4

970 708M40 970 709A40 970 709M40 4142

1.3563, 43CrMo4 1.7223, 41CrMo4

...

970 708A42 (En19C) 970 708H42 970 709A42

2244

4145

1.2332, 47CrMo4

G4052 SCM5H G4052 SCM445H G4105 SCM5, SCM445

970 708H45

A35-552 45SCD6 A35-553 45SCD6

...

...

4147

1.2332, 47CrMo4 1.3565, 48CrMo4

G4052 SCM5H G4052 SCM445H

970 708A47

A35-552 45SCD6 A35-553 45SCD6

...

...

(continued)


Appendix 7:

21/8/2008 4:57PM Plate # 0

pg 559




Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Alloy steels (continued) 1.7228, 50CrMo4

G4105 SCM5, SCM445

A35-571 50SCD6

1.7228, GS-50CrMo4 1.7230, 50CrMoPb4 1.7238, 49CrMo4 ...

4150

1.3565,48CrMo4 1.7228, 50CrMo4 1.7228, GS-50CrMo4 1.7230, 50CrMoPb4 1.7238, 49CrMo4

4161

1.7229, 61CrMo4 G4801 SUP13 1.7266, GS-58CrMnMo4 43

4320

4340

...

1.6565, 40NiCrMo6

G4103 SNCM23 G4103 SNCM420 G4103 SNCM420H

...

A35-571 50SCD6

...

3100 BW4 3146 CLA12 Grade C ...

20NCD7 A35-565 18NCD4 A35-565 20NCD7 ...

G4103 SNCM8 4670 818M40 G4103 SNCM439 970 2S.119 G4108 SNB23-1-5 G4108 SNB24-1-5

E4340

1.6562, 40NiCrMo7 3

...

4422

1.5419, 22Mo

...

4427

No international equivalent

...

23D5

...

15ND8

4615

...

...

4617

...

...

970 665A17 970 665H17 970 665M17 (En34)

4620

...

...

970 665A19 970 665H20 970 665M20

4626

...

...

970 665A24 (E35B)

4718

No international equivalent ...

...

...

18NCD4

...

...

...

A35-552 38C2 A35-556 38C2 A35-557 38C2 A35-552 42C2 A35-556 42C2 A35-557 42C2

7356 41Cr2KB

...

45C2

7847 45Cr2

...

50B40

1.7003, 38Cr2 1.7023, 28CrS2

50B44

...

...

...

...

...

...

1.2101, 62SiMnCr4

...

5115

1.7131, 16MnCr5, GS-16MnCr5 1.7139, 16MnCrS5

...

3608 G20Mo5

...


1.7138, 52MnCrB3

...

...

4820

5060

5332 40NiCrMo7 6926 40NiCrMo7 7845 40NiCrMo7 7874 40NiCrMo7 7356 40NiCrMo7KB ...

...


50B50

2523 2523-02

...

4817

1.3561, 44Cr2

3097 20NiCrMo7 5331 18NiCrMo7 7846 18NiCrMo7

...

...


...

...


5046

...

...

4815

50B46

...

...

4720

G4052 SMnC3H G4052 SMnC443H G4106 SMnC3 G4106 SMnC443 G5111 SCMnCr4

...

970 2S.119

...

...

...

2ND8

...

...

...

55C2

...

...

970 526M60 (En11)

61SC7 A35-552 60SC7

...

...

G4052 SCr21H

970 527A17

16MC5

G4052 SCr415H

970 527H17

A35-551 16MCS

(continued)

7846 16MnCr5

2127


21/8/2008 4:57PM Plate # 0

pg 560


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Alloy steels (continued) 1.7142, 16MnCrPb5 1.7160, 16MnCrB5

G4104 SCr21 G4104 SCr415

970 527M17

5117

1.3521, 17MnCr5 1.7016, 17Cr3 1.7131, 16MnCr5, GS-16MnCr5 1.7139, 16MnCrS5 1.7142, 16MnCrPb5 1.7168, 18MnCrB5

...

5120

1.2162, 21MnCr5 1.3523, 19MnCr5 1.7027, 20Cr4 1.7028, 20Cr5 4 1.7121, 20CrMnS3 3 1.7146, 20MnCrPb5 1.7147, GS-20MnCr5 1.7149, 20MnCrS5

G4052 SCr22H G4052 SCr420H G4052 SMn21H G4052 SMn421H G4104 SCr22 G4104 SCr420

5130

1.8401, 30MnCrTi4

G4052 SCr2H G4052 SCr430H G4104 SCr2 G4104 SCr430

970 530A30 (En18A) 970 530H30

28C4

5132

1.7033, 34Cr4 1.7037, 34CrS4

G4104 SCr3 G4104 SCr435

970 530A32 (En18B) 970 530A36 (En18C) 970 530H32

5135

1.7034, 37Cr4 1.7038, 37CrS4 1.7043, 38Cr4

G4052 SCr3H G4052 SCr435H

5140

1.7035, 41Cr4 1.7039, 41CrS4 1.7045, 42Cr4

G4052 SCr4H G4052 SCr440H G4104 SCr4 G4104 SCr440

5147

1.7145, GS-50CrMn4 4

5150

1.7145, GS-50CrMn4 4 1.8404, 60MnCrTi4

5155

1.7176, 55Cr3

G4801 SUP11 G4801 SUP9

5160

1.2125, 65MnCr4

G4801 SUP9A

51B60


E50100

1.2018, 95Cr1 1.3501, 100Cr2

...

E51100

1.2057, 105Cr4 1.2109, 125CrSi5 1.2127, 105MnCr4 1.3503, 105Cr4

...

E52100

1.2059, 120Cr5 1.2060, 105Cr5 1.2067, 100Cr6 1.3505, 100Cr6 1.3503, 105Cr4

...

...

18Cr4 A35-551 16MC5

...

...

...

A35-551 20MC5 A35-552 20MC5

7846 20MnCr5

...

...

...

A35-552 32C4 A35-553 32C4 A35-556 32C4 A35-557 32C4

7356 34Cr4KB 7874 34Cr4

...

3111 Type 3 970 530A36 (En18C) 970 530H36

38C4 A35-552-38Cr4 A35-553 38Cr4 A35-556 38Cr4 A35-557 38Cr4

5332 35CrMn5 6403 35CrMn5 5333 36CrMn4 7847 36CrMn4 7356 38Cr4KB 7845 36CrMn5 7874 36CrMn5 7847 38Cr4

...

3111 Type 3 970 2S.117 970 530A40 (En18D) 970 530H40 970 530M40

A35-552 42C4 A35-557 42C4 A35-556 42C4

5332 40Cr4 7356 41Cr4KB 7845 41Cr4 7874 41Cr4

2245

...

3100 BW2, BW3 3146 CLA 12 Grade A 3146 CLA 12 Grade B

50C4

...

3100 BW2 3100 BW3 3146 CLA 12 Grade A 3146 CLA 12 Grade B

...

...

...

...

2230

A35-571 55C3

...

...

...

...

...

...

A35-565 100C2

...

...

...

...

3160 G90Cr4

...

...

100C6 3097 100Cr6

2258

... 970 527A60 (En48) 970 527H60

970 534A99 (En31) 970 535A99 (En31)

(continued)


Appendix 7:

21/8/2008 4:57PM Plate # 0

pg 561




Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Alloy steels (continued) 1.3514, 101Cr6 1.3520, 100CrMn6 6118


6150

1.8159, GS-50CrV4

8115


81B45


G4801 SUP10

8615

...

...

8617

...

...

8620

1.6522, 20NiCrMo2 1.6523, 21NiCrMo2 1.6526, 21NiCrMoS2 1.6543, 21NiCrMo2 2

8622

1.6541, 23MnNiCrMo5 2

8625

...

8627


8630

1.6545, 30NiCrMo2 2

8637

8640

...

1.6546, 40NiCrMo2 2

G4052 SNCM21H G4052 SNCM220H G4103 SNCM21 G4103 SNCM220

...

...

...

970 735A50 (En47) 970 S.204

...

A35-552 50CV4 A35-553 50CV4 A35-571 50CV4

3545 50CrV4 7065 50CrV4 7845 50CrV4 7874 50CrV4

15NCD2 15NCD4

3097 16NiCrMo2 5331 16NiCrMo2 7846 16NiCrMo2

18NCD4 18NCD6

2772 806M20 970 805A20 970 805H20 970 805M20 (En362)

18NCD4 5331 20NiCrMo2 20NCD2 6403 20NiCrMo2 A35-551 19NCDB2 7846 20NiCrMo2 A35-552 19NCDB2 A35-551 20NCD2 A35-553 20NCD2 A35-565 20NCD2 A35-566 20NCD2

2772 806M22

23NCDB4

970 805A22 970 805H22 970 805M22

A35-556 23MNCD5 A35-556 23NCDB2 A35-566 22NCD2

970 805H25 970 805M25

25NCD4 A35-556 25MNCD6 A35-566 25MNDC6

...


8645


86B45


8650


8655


8660

...

...

...

2506-03 2506-08

...

...

...

...

30NCD2

7356 30NiCrMo2KB

...

...

970 945M38 (En100)

40NCD3

5332 38NiCrMo4 7356 38NiCrMo4KB 7845 39NiCrMo3 7874 39NiCrMo3

... ...

...

3111 Type 7, 2S.147 970 945A40 (En 100C)

40NCD2 40NCD2TS

5333 40NiCrMo4 7356 40NiCrMo2KB 7845 40NiCrMo2 7874 40NiCrMo2 7847 40NiCrMo3

...

40NCD3TS 40NCD3 8642

...

970 805A17 970 805H17 970 805M17 (En 361)

2230

...

970 805A60 970 805H60 (continued)

...

...

...


21/8/2008 4:57PM Plate # 0

pg 562


Table A7.1

(continued)

United states (SAE)


Japan (JIS)

United Kingdom (BS)

France (AFNOR NF)

Italy (UNI)

Sweden (SS14)

Alloy steels (continued) 8720


8740

1.6546, 40NiCrMo2 2

...

3111 Type 7, 2S.147

40NCD2 40NCD2TS 40NCD3TS

8822


9254


9260

...

E9310

1.6657, 14NiCrMo13 4

94B15


94B17


94B30


G4801 SUP7 ...

970 250A58 (En45A) 970 250A61 (En45A)

60S7 61S7

970 832H13 970 832M13 (En36C) S.157

16NCD13

7356 40NiCrMo2KB 7845 40NiCrMo2 7874 40NiCrMo2

... 6932 15NiCrMo13 9335 10NiCrMo13

...

... ...



21/8/2008 5:03PM Plate # 0

pg 563


APPENDIX 8

Non-AISI to AISI Steel Cross Reference Table A8.1 Cross reference to steels by country Country

FRANCE

Designation

AISI

Country

AFNOR 20 MC 5 20 MC 5 20 NCD 2 20 NCD 2 20 NCD 2 20 NCD 2 22 NCD 2 22 NCD 2 22 NCD 2 22 NCD 2 25 CD 4 (S) 25 CD 4 (S) 32 C 4 32 C 4 32 DCV 28 35 CD 4 35 CD 4 35 CD 4 TS 35 CD 4 TS 35 M 5 38 C 4 38 C 4 40 CD 4 40 CD 4 40 CD 4 40 CD 4 40 M 5 40 M 5 42 C 2 42 C 2 42 C 4 42 C 4 42 CD 4 42 CD 4 42 CD 4 42 CD 4 45 C 2 45 C 2 50 CV 4 50 CV 4 55 C 3 55 C 3 55 S 7 55 WC 20 60 S 7 60 S 7 61 SC 7 61 SC 7 90 MV 8 100 C 6 CC 20 CC 35

Designation

AISI

FRANCE (continued) AFNOR 5120 5120H 8617 8617H 8620 8620H 8617 8617H 8620 8620H 4130 4130H 5130H 5132 H10 4135 4135H 4135 4135H 1039 5132H 5135 4137 4137H 4140 4140H 1335 1335H 5140H 5150 5135H 5140 4137 4137H 4140 4140H 5140H 5150 6150 6150H 5155 5155H 9255 S1 9260 9260H 9260 9260H O2 E52100 1020 1035 (continued)

Adapted from Engineering Properties of Steel, ASM International, 1982, p 509–523

CC 55 XC 10 XC 15 XC 15 XC 18 XC 18 XC 18 S XC 25 XC 32 XC 35 XC 38 XC 38 TS XC 38 TS XC 42 XC 42 XC 42 TS XC 42 TS XC 45 XC 45 XC 48 XC 48 XC 60 XC 65 XC 68 XC 90

1060 1010 1015 1017 1015 1017 1023 1023 1034 1034 1034 1038 1038H 1045 1045H 1045 1045H 1045 1045H 1045 1045H 1064 1064 1070 1086

Z 2 CND 17.12 Z 2 CND 19.15 Z 6 CA 13 Z 6 CN 18.09 Z 6 CND 17.11 Z 6 CNN6 18.10 Z 6 CNT 18.10 Z 6 CNU 17.04 Z 8 C 17 Z 8 CD 17.01 Z 10 C 13 Z 10 C 14 Z 10 CF 17 Z 10 CNF 18.09 Z 12 C 13 Z 12 C 13 M Z 12 CN 17.08 Z 12 CNS 25.20 Z 12 CNS 25.20 Z 15 CN 16.02 Z 15 CN 24.13 Z 18 N 5 Z 20 C 13 Z 30 WCV 9 Z 38 CDV 5 Z 40 COV 5

316L 317L 405 304 316 347 321 431 430 434 410 410 430F 303 410 403 301 310 314 431 309S A2515 420 H21 H11 H13


21/8/2008 5:04PM Plate # 0

pg 564


Table A8.1

(continued)

Country

Designation

Country

AISI

GERMANY (continued)

FRANCE (continued) AFNOR Z 80 WCV 18-04-01 Z 80 WKCV 18-05-04-01 Z 85 DCWV 08-04-02-01 Z 85 DCWV 08-04-02-01 Z 85 WDCV 06-05-04-02 Z 90 WDCV 06-05-04-02 Z 100 CDV 5 Z 110 WKCDV 07-05-04-04-02 Z 110 WKCDV 07-05-04-04-02 Z 120 WDCV 06-05-04-03 Z 130 WDCV 06-05-04-04 Z 200 C 12 GERMANY

T1 T4 H41 M1 M2 M3 (Class 1) A2 M41 M42 M3 (Class 2) M3 (Class 2) D3

DIN 1.0204 1.0402 1.0419 1.0501 1.0601 1.0700 1.0702 1.0711 1.0715 1.0718 1.0718 1.0904 1.0909 1.0909 1.0912 1.0912 1.1121 1.1133 1.1141 1.1141 1.1151 1.1157 1.1158 1.1165 1.1165 1.1167 1.1167 1.1172 1.1176 1.1176 1.1181 1.1186 1.1191 1.1191 1.1209 1.1210 1.1221 1.1226 1.1230 1.1231 1.1269 1.1273

1008 1020 1016 1035 1060 1108 1109 1212 1213 12L13 12L14 9255 9260 9260H 1345 1345H 1010 1022 1015 1017 1023 1039 1025 1330 1330H 1335 1335H 1030 1038 1038H 1034 1040 1045 1045H 1055 1050 1064 1548 1065 1070 1086 1090 (continued)


Designation

AISI

DIN 1.1274 1.2080 1.2330 1.2341 1.2343 1.2344 1.2363 1.2365 1.2379 1.2510 1.2550 1.2581 1.2606 1.2625 1.2735 1.2842 1.3202 1.3246 1.3246 1.3249 1.3249 1.3255 1.3265 1.3342 1.3343 1.3344 1.3346 1.3346 1.3348 1.3355 1.3501 1.3503 1.3505 1.4001 1.4002 1.4005 1.4006 1.4016 1.4021 1.4024 1.4057 1.4104 1.4112 1.4113 1.4125 1.4301 1.4303 1.4303 1.4305 1.4306 1.4310 1.4401 1.4404 1.4438 1.4449 1.4510 1.4512 1.4532 1.4541 1.4546 1.4550 1.4568 1.4828 1.4833 1.4841 1.4841

1095 D3 P20 P4 H11 H13 A2 H10 D2 O1 S1 H21 H12 H23 P6 O2 T15 M41 M42 M33 M34 T4 T5 M3 (Class 1) M2 M3 (Class 2) H41 M1 M7 T1 E50100 E51100 E52100 410S 405 416 410 430 420 403 431 430F 440B 434 440C 304 305 308 303 304L 301 316 316L 317L 317 430Ti 409 632 321 348 347 631 309 309S 310 314


Appendix 8:

21/8/2008 5:04PM Plate # 0

pg 565

Non-AISI to AISI Steel Cross Reference / 565

Table A8.1 (continued) Country

GERMANY (continued)

Designation

DIN 1.4935 1.4971 1.4980 1.5069 1.5419 1.5419 1.5419 1.5680 1.5711 1.6511 1.6523 1.6523 1.6523 1.6523 1.6543 1.6543 1.6543 1.6543 1.6543 1.6543 1.6545 1.6545 1.6546 1.6546 1.6546 1.6546 1.6562 1.6562 1.6565 1.6565 1.6755 1.6755 1.7006 1.7006 1.7007 1.7007 1.7030 1.7033 1.7033 1.7034 1.7034 1.7035 1.7035 1.7138 1.7138 1.7147 1.7147 1.7176 1.7176 1.7218 1.7218 1.7220 1.7220 1.7223 1.7225 1.7225 1.7225 1.7225 1.7228 1.7228 1.7228 1.7228 1.7362 1.7511 1.7511 1.8159 1.8159

Country

AISI

ITALY 422 661 660 1340H 4419 4419H 4422 A2515 3140 9840 8617 8617H 8620 8620H 8622 8622H 8720 8720H 8822 8822H 8630 8630H 8640 8640H 8740 8740H E4340 E4340H 4340 4340H 4718 4718H 5140H 5150 50B40 50B40H 5130 5130H 5132 5132H 5135 5135H 5140 50B50 50B50H 5120 5120H 5155 5155H 4130 4130H 4135 4135H 4142H 4137 4137H 4140 4140H 4147 4147H 4150 4150H 501 6118 6118H 6150 6150H

Designation

9 SMn 23 9 SMnPb 23 9 SMnPb 23 10 S 20 20 NiCrMo 20 NiCrMo 2 20 NiCrMo 2 20 NiCrMo 2 25 CrMo 4 25 CrMo 4 25 CrMo 4 KB 25 CrMo 4 KB 30 NiCrMo 2 KB 30 NiCrMo 2 KB 34 Cr 4 KB 34 Cr 4 KB 34 CrMo 4 KB 34 CrMo 4 KB 35 CrMo 4 35 CrMo 4 35 CrMo 4 F 35 CrMo 4 F 38 Cr 4 KB 38 Cr 4 KB 38 CrB 1 KB 38 CrB 1 KB 38 CrMo 4 38 CrMo 4 KB 38 CrMo 4 KB 38 CrMo 4 KB 38 CrMo 4 KB 38 NiCrMo 4 40 Cr 4 40 Cr 4 40 CrMo 4 40 CrMo 4 40 CrMo 4 40 CrMo 4 40 NiCrMo 2 KB 40 NiCrMo 2 KB 40 NiCrMo 2 KB 40 NiCrMo 2 KB 40 NiCrMo 7 40 NiCrMo 7 40 NiCrMo 7 KB 40 NiCrMo 7 KB 41 Cr 4 KB 41 Cr 4 KB 50 CrV 4 50 CrV 4 55 Si 8 58 WCr 9 KU 88 MnV 8 KU 100 Cr 6 C 20 C 35 C 60 CB 10 FU CB 35 G 22 Mn 3 G 22 Mo 5 G 22 Mo 5 G 22 Mo 5 G 40 CrMo 4 G 40 CrMo 4 G 40 CrMo 4 (continued)


AISI

UNI 1213 12L13 12L14 1212 8617H 8617 8620 8620H 4130 4130H 4130 4130H 8630 8630H 5130H 5132 4135 4135H 4135 4135H 4135 4135H 5132H 5135 50B40 50B40H 4142H 4137 4137H 4140 4140H 9840 5135H 5140 4137 4137H 4140 4140H 8640 8640H 8740 8740H E4340 E4340H E4340 E4340H 5135H 5140 6150 6150H 9255 S1 O2 E52100 1020 1035 1060 1008 1030 1022 4419 4419H 4422 4137 4137H 4140


21/8/2008 5:04PM Plate # 0

pg 566


Table A8.1

(continued)

Country

ITALY (continued)

Designation

G 40 CrMo 4 ICL 472 T X 2 CrNi 18 11 X 2 CrNi 18 11 KG X 2 CrNi 18 11 KW X 2 CrNiMo 17 12 X 3 Cr Ni 18 11 X 5 CrNi 18 10 X 5 CrNiMo 17 12 X 5 CrNiMo 18 15 X 6 CrAl 13 X 6 CrNi 23 14 X 6 CrNiTi 18 11 X 6 CrNiTi 18 11 KG X 6 CrNiTi 18 11 KT X 6 CrNiTi 18 11 KW X 8 Cr 17 X 8 CrMo 17 X 8 CrNi 19 10 8 CrNi 19 10 X 8 CrNiNb 18 11 X 10 CrNiS 18 09 X 10 CrS 17 X 12 Cr 13 X 12 CrNi 17 07 X 12 CrS 13 X 16 CrNi 16 X 16 CrNi 23 14 X 16 CrNiSi 25 20 X 16 CrNiSi 25 20 X 20 Cr 13 X 22 CrNi 25 20 X 22 CrNi 25 20 X 28 W 09 KU X 35 CrMo 05 KU X 35 CrMoV 05 KU X35 CrMoW 05 KU X 75 W 18 KU X 78 WCo 1805 KU X 80 WCo 1810 KU X 82 MoW 09 KU X 82 MoW 09 KU X 82 WMo 0605 KU X 150 CrMo 12 KU X 150 WCoV 130505 KU X 210 Cr 13 KU JAPAN

Country

AISI

JAPAN (continued)

UNI

S 15 CK S 15 CK S 17 C S 17 C S 20 C S 20 CK S 22 C S 25 C S 28 C S 38 C S 40 C S 45 C S 45 C S 48 C S 48 C S 53 C S 55 C SCCrM 1 SCCrM 1 SCCrM 3 SCCrM 3 SCM 1 SCM 1 SCM 2 SCM 2 SCM 4 SCM 4 SCM 4 SCM 4 SCM 4 H SCM 4 H SCM 4 H SCM 4 H SCM 5 SCM 5 SCM 5 SCM 5 SCM 5 H SCM 5 H SCM 5 H SCM 5 H SCMn 2 SCMn 2 SCMn 3 SCMn 3 SCPH 11 SCPH 11 SCPH 11 SCr 2 SCr 2 SCr 2H SCr 2H SCr 3 H SCr 3 H SCr 4 H SCr 4 H SCS 19 SKD 1 SKD 5 SKD 6 SKD 12 SKD 61 SKD 62 SKH 2 SKH 3 SKH 4A

4140H 321 304L 304L 304L 316L 304L 304 316 317 405 309S 321 321 321 321 430 434 305 308 347 303 430F 410 301 416 431 309 310 314 420 310 314 H21 H11 H13 H12 T1 T4 T5 H41 M1 M2 D2 T15 D3

JIS S 9 CK S 10 C S 12 C S 15 C S 15 C

Designation

1010 1010 1010 1015 1017 (continued)


AISI

JIS 1015 1017 1015 1017 1023 1023 1023 1025 1025 1034 1040 1045 1045H 1045 1045H 1050 1050 4130 4130H 4315 4135H 4135 4135H 4130 4130H 4137 4137H 4140 4140H 4137 4137H 4140 4140H 4147 4147H 4150 4150H 4147 4147H 4150 4150H 1330 1330H 1335 1335H 4419 4419H 4422 5130H 5132 5130H 5132 5132H 5135 5135H 5140 304L D3 H21 H11 A2 H13 H12 T1 T4 T5


Appendix 8:

21/8/2008 5:04PM Plate # 0

pg 567



JAPAN (continued)

Designation

Country

AISI

SWEDEN

JIS SKH 9 SKH 52 SKH 53 SMn 1 H SMn 1 H SMn 2 SMn 2 SMn 2 H SMn 2 H SMnC 21 SNCM 8 SNCM 8 SNCM 21 SNCM 21 SNCM 21 SNCM 21 SNCM 21 H SNCM 21 H SNCM 21 H SNCM 21 H SUH 309 SUH 310 SUH 409 SUH 616 SUM 11 SUM 12 SUM 21 SUM 22 SUM 22 L SUM 22 L SUM 23 L SUM 24 L SUM 24 L SUP 4 SUP 10 SUP 10 SUP 11 SUP 11 SUS 301 SUS 303 SUS 304 SUS 304 L SUS 305 SUS 305 SUS 305 J1 SUS 305 J1 SUS 316 SUS 316 L SUS 317 SUS 321 SUS 347 SUS 403 SUS 405 SUS 410 SUS 410S SUS 416 SUS 420 J1 SUS 430 SUS 430 SUS 431 SUS 434 SUS 440 B SUS 440 C SUS Y 310 SUS Y 310 SUS Y 316

M2 M3 (Class 2) M3 (Class 2) 1330 1330H 1335 1335H 1335 1335H 1022 4340 4340H 8617 8617H 8620 8620H 8617 8617H 8620 8620H 316 316L 409 422 1109 1109 1212 1213 12L13 12L14 12L13 12L13 12L14 1095 6150 6150H 50B50 50B50H 301 303 304 304L 305 308 305 308 316 316L 317 321 347 403 405 410 410S 403 420 430 430F 431 434 440B 440C 310 314 316

1370 1370 1450 1550 1665 1672 1672 1678 1770 1778 1870 1914 1914 2090 2120 2120 2225 2225 2230 2230 2234 2234 2242 2244 2244 2244 2244 2260 2302 2303 2320 2325 2332 2337 2338 2346 2347 2348 2352 2367 2383 2722

AISI

1015 1017 1020 1035 1064 1045 1045H 1064 1070 1070 1095 12L13 12L14 9255 1335 1335H 4130 4130H 6150 6150H 4135 4135H H13 4137 4137H 4140 4140H A2 410 420 430 434 304 321 347 303 316 316L 304L 317L 430F M2

UNITED KINGDOM B.S.

(continued) Adapted from Engineering Properties of Steel, ASM International, 1982, p 509–523

Designation

SS14

040 A 20 060 A 35 060 A 62 060 A 96 070 M 20 080 A 32 080 A 35 080 A 37 080 A 40 080 M 36 2 S 93 2 S 117 2 S 117 2 S 119 2 S 119 2 S 130 2 S 516 2 S 516 2 S 517 2 S 517 3 S 95 3 S 95

1020 1035 1060 1095 1020 1035 1035 1035 1040 1035 1040 5135H 5140 4340 4340H 348 1345 1345H 1345 1345H 4340 4340H


21/8/2008 5:04PM Plate # 0

pg 568


Table A8.1 Country

(continued) Designation

UNITED KINGDOM B.S. (continued) 5 S 80 120 M 36 150 M 36 220 MO7 250 A 53 250 A 58 250 A 58 302 S 17 303 S 21 304 S 12 304 S 14 304 S 15 304 S 16 304 S 18 304 S 22 304 S 25 304 S 40 310 S 24 310 S 24 316 S 12 316 S 14 316 S 16 316 S 18 316 S 22 316 S 24 316 S 25 316 S 26 316 S 29 316 S 30 316 S 30 316 S 31 316 S 37 316S 40 316 S 41 316 S 82 317 S 12 321 S 12 321 S 18 321 S 22 321 S 27 321 S 40 321 S 49 321 S 50 321 S 59 321 S 87 347 S 17 347 S 17 347 S 18 347 S 40 403 S 17 405 S 17 409 S 17 410 S 21 416 S 21 420 S 29 420 S 37 430 S 15 431 S 29 434 S 19 530 A 30 530 A 32 530 A 32 530 A 36 530 A 36 530 A 40 530 A 40

Country

AISI

Designation

UNITED KINGDOM B.S. (continued) 530 H 30 530 H 32 530 H 32 530 H 36 530 H 36 530 H 40 530 H 40 530 M 40 530 M 40 534 A 99 535 A 99 640 M 40 708 A 37 708 A 37 708 A 42 708 A 42 708 A 42 708 A 42 708 M 40 708 M 40 708 M 40 708 M 40 708 A 40 709 M 40 709 M 40 709 M 40 735 A 50 735 A 50 805 A 20 805 A 20 805 A 20 805 A 20 805 A 20 805 A 20 805 H 20 805 H 20 805 H 20 805 H 20 805 M 20 805 M 20 805 M 20 805 M 20 816 M 40 817 M 40 817 M 40 3111 Type 6 3111 Type 6 ANC 1 Grade A ANC 3 Grade B BA 2 BD 2 BD 3 BH 11 BH 12 BH 13 BH 21 BM 1 BM 1 BM 2 BM 34 BM 34 BO 1 BO 2 BT 1 BT 4 BT 5

431 1039 1039 1213 9255 9260 9260H 304 303 304L 304L 304 304 304 304L 304 304 310 314 316L 316L 316 316 316L 316L 316 316 316L 316 316L 316L 316L 316 316 316L 317L 321 321 321 321 321 321 321 321 321 347 348 348 348 410S 405 409 410 416 403 420 430 431 434 5130 5130H 5132 5132H 5135 5135H 5140


AISI

5130 5130H 5132 5132H 5135 5135H 5140 5135H 5140 E52100 E52100 3140 4135 4135H 4137 4137H 4140 4140H 4137 4137H 4140 4140H 4137H 4137 4140 4140H 6150 6150H 8622 8622H 8720 8720H 8822 8822H 8617 8617H 8620 8620H 8617 8617H 8620 8620H 9840 4340 4340H 4340 4340H 410 347 A2 D2 D3 H11 H12 H13 H21 H41 M1 M2 M33 M34 O1 O2 T1 T4 T5


Appendix 8:

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pg 569



Designation

UNITED KINGDOM B.S. (continued) BT 15 CDS-18 CDS-20 CDS 105/106 CDS 110 CDS 110 En. 44 B En. 47 En. 47 En. 56 A En. 56 B En. 58 B En. 58 C En. 58 E En. 58 F En. 58 G En. 58 H S. 139 S. 139 S. 525 S. 527 S. 536 S. 537 Type 3 Type 3 Type 7 Type 7 Type 7 Type 7 Type 8 Type 8 UNITED STATES

Country

AISI

UNITED STATES (continued)

T15 420 321 1039 4130 4130H 1095 6150 6150H 410 403 321 321 304 347 347 316 E4340 E4340H 348 348 304L 316L 5132H 5135 8640 8640H 8740 8740H E4340 E4340H

AMS 5010 D 5024 C 5032 5040 5042 5044 5045 5047 5053 5060 5069 5070 5075 5077 5080 5082 5085 5110 5112 5112 E 5115 5120 D 5121 5122 5132 5331 5333 5342 5343 5344 5354 5355 5369

1212 1137 1020 1010 1010 1010 1020 1010 1010 1015 1018 1022 1025 1025 1035 1035 1050 1080 1090 1086 1070 1074 1095 1095 1095 4340 8615 630 630 630 615 630 651 (continued)


Designation

AISI

AMS 5376 5501 5502 5503 5504 5505 5506 5507 5508 5510 5511 5512 5513 5514 5515 5516 5517 5518 5519 5520 5521 5522 5523 5524 5525 5526 5527 5528 5529 5531 5532 5546 5547 5548 5549 5554 5556 5557 5558 5559 5560 5561 5565 5566 5567 5568 5570 5571 5572 5573 5574 5575 5576 5577 5578 5579 5585 5591 5592 5594 5602 5604 5610 5611 5612 5613

661 304 501 430 410 410 420 316L 615 321 304L 347 304 305 302 302 301 301 301 632 310S 314 309S 316 660 651 651 631 631 661 661 633 634 633 634 633 347 321 347 321 304 304 304 304 631 321 347 310S 316 309S 347 321 310S 651 661 410 330 634 501 630 416Se 403 403 410


21/8/2008 5:04PM Plate # 0

pg 570


Table A8.1

(continued)

Country


Designation

Country

AISI


AMS 5615 5615 5618 5620 5620 5621 5622 5626 5627 5628 5630 5631 5632 (Type 1) 5632 (Type 2) 5636 5637 5639 5640 (Type 1) 5640 (Type 2) 5641 5643 5644 5645 5646 5647 5648 5649 5650 5651 5652 5653 5654 5655 5657 5673 5674 5678 5680 5681 5685 5686 5688 5689 5690 5691 5694 5695 5697 5716 5720 5721 5722 5731 5732 5734 5735 5736 5737 5738 5742 5743 5744 5745 5768 5769 5774

414 615 440C 420F 420F(Se) 420 630 T1 430 431 440C 440A 440F 440F(Se) 302 302 304 303 303Se 303Se 630 631 321 347 304L 316 316F 309S 310S 314 316L 347 422 632 631 347 631 347 347 305 305 302 321 316 316 310 310 304 330 651 651 651 660 660 660 660 660 660 303Se 634 634 634 633 661 661 633


Designation

AISI

AMS 5775 5776 5780 5781 5794 5795 5804 5805 5812 5813 5817 5821 5824 5825 6260 F 6265 B 6272 6274 6275 6275 A 6276 6277 6280 6281 6290 6294 6322 6323 6325 6327 6342 C 6350 6355 6356 6358 6359 6360 6361 6362 6365 C 6370 6371 6372 C 6373 6381 6382 6390 6395 6414 6415 6437 6440 6441 6442 B 6443 6444 6444 6446 6447 6448 6449 6450 6455 6466 6467 6485

633 410 634 634 661 661 660 660 632 632 615 410 631 630 E9310 E9310 8617 8620 94B17 94B15 8620 8620 8630 8630 4615 4620 8740 8740 8740 8740 9840 4130 8630 4130 8740 4340 4130 4130 4130 4135 4130 4130 4135 4130 4140 4140 4140 4140 4340 4340 H11 E52100 E52100 E50100 E51100 E52100 E52100 E51100 E52100 6150 E51100 6150 6150 502 502 H11


Appendix 8:

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pg 571




Designation

Country

AISI


AMS 6487 6488 6530 6550 7240 7301 7304 7470

H11 H11 8630 8630 1060 6150 1095 615

ASME 5041 SA182 SA182 SA182 SA182 SA182 SA182 SA182 SA182 SA182 SA182 SA182 SA193 SA193 SA193 SA193 SA194 SA194 SA194 SA194 SA194 SA194 SA194 SA194 SA194 (Type 3) SA194 (Type 6) SA194 (Type 8) SA213 SA213 SA213 SA213 SA213 SA213 SA213 SA213 SA213 SA213 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240 SA240

1006 304 304L 304N 310 316 316L 316N 321 347 348 430 305 316 321 347 303 303Se 305 316 321 347 416 416Se 501 410 304 304 304L 304N 310 316 316L 316N 321 347 348 302 304 304L 304N 305 309S 310S 316 316L 316N 317 317L 321 347 348 405 410 410S 430 (continued)


Designation

AISI

ASME SA240 (XM-8) SA240 (XM-21) SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 SA249 (XM-19) SA249 (XM-29) SA268 SA268 SA268 SA268 SA268 SA268 SA268 (XM-8) SA268 (XM-27) SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 SA312 (XM-19) SA312 (XM-29) SA320 SA320 SA320 SA320 SA320 (B8) SA320 (B8C) SA358 SA358 SA358 SA358 SA358 SA358 SA358 SA358 SA358 SA376 SA376 SA376 SA376 SA376 SA376 SA376 SA387 (Type 5) SA387 (Type 5) SA403 SA403 SA403 SA403

S43035 304HN 304 304L 304N 309 310 316 316L 316N 317 321 347 348

329 405 409 410 430 446 S43035 304 304L 304N 309 310 316 316L 316N 317 321 347 348

303 303Se 316 321 304 347 304 304N 309 310 316 316N 321 347 348 304 304N 316 316N 321 347 348 501 502 304 304L 309 310


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pg 572


Table A8.1

(continued)

Country


Designation

Country

AISI


ASME SA403 SA403 SA403 SA403 SA403 SA403 SA403 SA403 (XM-19) SA409 SA409 SA409 SA409 SA409 SA409 SA409 SA409 SA412 SA412 SA412 (XM-19) SA430 SA430 SA430 SA430 SA430 SA430 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 SA479 (XM-8) SA479 (XM-19) SA479 (XM-27) SA564 SA564 (XM-25) SA638 SA688 SA688 SA688 SA688 SA688 (XM-29) SA705 SA705 SA705 (XM-12) SA705 (XM-13) SA705 (XM-25) SA737 (XM-27)

316 316L 316N 317 321 347 348 304 309 310 316 317 321 347 348 201 304 304N 316 316N 321 347 302 304 304L 304N 310S 316 316L 316N 321 347 348 405 410 430 S43035

630 660 304 304L 316 316L 630 631

ASTM A26 A29 A29 A29 A29 A29 A29 A29

1064 1005 1006 1008 1010 1012 1015 1016 (continued)


Designation

AISI

ASTM A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A29 A36 A57 A59 A107 A107 A107 A108 A108 A108 A108 A108 A108 A108 A108 A108 A108

1017 1018 1019 1020 1021 1022 1023 1025 1026 1030 1034 1035 1038 1038H 1039 1040 1044 1045 1045H 1046 1050 1055 1059 1060 1064 1065 1070 1074 1080 1086 1090 1095 1108 1109 1116 1119 1132 1137 1141 1144 1211 1212 1213 12L13 12L14 1215 1547 1548 15B48H 9260 1064 9260 1117 1118 1141 1008 1010 1015 1016 1017 1018 1020 1030 1035 1040


Appendix 8:

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Designation

Country

AISI


ASTM A108 A108 A108 A108 A108 A108 A108 A108 A108 A108 A131 A131 (A) A131 (AH32, DH32, EH32) A131 (AH36, DH36, EH36) A131 (B) A131 (CS, DS) A131 (D) A131 (E) A135 A139 A139 (B) A139 (C) A139 (D) A139 (E) A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A167 A176 A176 A176 A176 A176 A176 A176 A176 A177 A181 A182 A182 A182 A182 A182 A182 A182 A182 A182 A182 A182 A193 A193 A193

1117 1118 1137 1141 1144 1211 1212 1213 12L14 1215

301 302 302B 304 304L 305 308 309 309S 310 310S 316 316L 317 317L 321 347 348 403 405 409 410 410S 430 442 446 301 1034 304 304L 304N 310 316 316L 316N 321 347 348 430 304 316 321 (continued)


Designation

AISI

ASTM A193 A193 A193 A194 A194 A194 A194 A194 A194 A194 A194 A194 A194 A194 (grade 8F) A213 A213 A213 A213 A213 A213 A213 A213 A213 A213 A228 A229 A230 A231 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 A240 (XM-8) A240 (XM-21) A249 A249 A249 A249 A249 A249 A249 A249 A249 A249 A249 A249 A249 A268 A268 A268 A268 A268

347 410 501 303 303Se 304 316 321 347 410 416 416Se 501 303Se 304 304L 304N 310 316 316L 316N 321 347 348 1086 1065 1064 6150 302 304 304L 304N 305 309S 310S 316 316L 316N 317 317L 321 348 405 410 410S 430 S43035 304HN 304 304L 304N 305 309 310 316 316L 316N 317 321 347 348 329 405 409 430 430Ti


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pg 574


Table A8.1

(continued)

Country


Designation

Country

AISI


ASTM A268 A268 A268 (XM-8) A269 A269 A269 A269 A269 A269 A269 A270 A271 A271 A271 A273 A274 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 A276 (XM-21) A276 (XM-27) A284 (C) A295 A295 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304 A304

443 446 S43035 304 316 316L 317 321 347 348 304 304 321 347 1026 9840 302 302B 304 304L 304N 305 308 309 309S 310 310S 314 316 316L 316N 317 321 347 348 403 405 410 414 420 430 431 440A 440B 440C 446 304HN 304HN 304HN E50100 E51100 1330H 1335H 1340H 1345H 15B48H 4027H 4028H 4037H 4042H 4047H 4118H 4130H 4135H 4137H 4140H


Designation

AISI


4142H 4145H 4147H 4150H 4161H 4320H 4340H 4419H 4620H 4626H 4718H 4720H 4815H 4817 4817H 4820H 50B40H 50B44H 5046H 50B46H 50B50H 50B60H 5120H 5130H 5132H 5135H 5140H 5150H 5155H 5160H 51B60H 6118 6150H 81B45H 8617H 8620H 8622H 8625H 8627H 8630H 86B30H 8637H 8640H 8642H 8645H 86B45H 8650H 8655H 8660H 8720H 8740H 8822H 9260H 94B15H 94B17H 94B30H E4340 E4340H E9310H 1137 1141 1144 304 304L 304N 309


Appendix 8:

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pg 575




Designation

Country

AISI



310 316 316L 316N 317 321 347 348 302 304 305 316 631 202 302 302B 303 303Se 304 304L 305 308 309 309S 310 310S 314 316 316L 317 321 347 348 403 405 410 414 416 416Se 420 430 430F 430F(Se) 431 440A 440B 440C 446 501 502 303 303Se 304 316 321 347 1330 1335 1340 1345 3140 4023 4024 4027 4028 4037


Designation

AISI


4042 4047 4118 4130 4137 4140 4142 4145 4147 4150 4161 4320 4340 4419 4615 4620 4626 4718 4720 4815 4817 4820 50B40 50B44 50B46 50B50 50B60 5120 5130 5132 5135 5140 5150 5155 5160 51B60 6118 6150 81B45 8615 8617 8620 8622 8625 8627 8630 8637 8640 8642 8645 8650 8655 8660 8720 8740 8822 9255 9260 94B17 94B30 94B40 9840 E9310 E51100 E52100 1330


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pg 576


Table A8.1

(continued)

Country


Designation

Country

AISI



1335 1340 1345 3140 4023 4024 4027 4028 4037 4042 4047 4118 4130 4137 4140 4142 4145 4147 4150 4161 4320 4340 4419 4615 4620 4626 4718 4720 4815 4817 4820 50B60 5120 5130 5132 5135 5140 5150 5155 5160 51B60 6150 8617 8620 8622 8625 8627 8630 8637 8640 8642 8645 8655 8660 8720 8740 8822 9260 94B17 94B30 E4340 E52100 4135 304N 309 310


Designation

AISI

ASTM A358 A358 A358 A358 A358 A368 A368 A368 A368 A376 A376 A376 A376 A376 A376 A376 A387 (5) A387 (5) A403 A403 A403 A403 A403 A403 A403 A403 A403 A403 A403 A409 A409 A409 A409 A409 A409 A409 A412 A412 A429 A429 A429 (XM-19) A430 A430 A430 A430 A430 A430 A453 A453 A457 A458 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473

316 316N 321 347 348 302 304 305 316 304 304N 316 316N 321 347 348 501 502 304L 304N 309 310 316 316L 316N 317 321 347 348 304 309 310 316 317 321 347 201 202 201 202 304 304N 316 316N 321 347 651 660 651 651 202 302 302B 303 303Se 304 304L 305 308 309 309S 310 310S 314 316


Appendix 8:

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pg 577




Designation

Country

AISI


ASTM A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A473 A477 A478 A478 A478 A478 A478 A478 A478 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 A479 (XM-8) A492 A492 A492 A493 A493 A493 A493 A493 A493 A493 A493 A493 A493 A505 A505 A505 A505 A505 A505 A505

316L 317 321 347 403 405 410 410S 414 416 416Se 420 430 430F 430F(Se) 431 440A 440B 440C 446 501 502 651 302 304 304L 305 316 316L 317 302 304 304L 304N 310S 316 316L 316N 321 347 348 403 405 410 430 S43035 302 304 305 302 304 305 321 347 384 410 430 431 440C 4118 4130 4137 4140 4142 4145 4147


Designation

AISI


4150 4320 4340 4615 4620 4718 4815 4820 5130 5132 5140 5150 5160 6150 8615 8617 8620 8630 8640 8642 8645 8650 8655 8660 8720 8740 9260 E4340 E51100 E52100 1005 1006 1008 1010 1012 1015 1016 1017 1018 1019 1020 1021 1022 1023 1025 1026 1030 1035 1038 1039 1040 1044 1045 1046 1050 1055 1060 1070 1080 1090 1095 1547 1548 302 304 304L


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pg 578


Table A8.1

(continued)

Country


Designation

Country

AISI



305 309 309S 310 310S 316L 317 321 329 347 403 405 410 414 416Se 430 440A 443 446 1025 1030 1016 1017 1018 1019 4130 8620 1008 1010 1012 1015 1017 1018 1019 1020 1021 1022 1025 1026 1030 1035 1040 1045 1050 1330 1335 1340 1345 3140 4023 4024 4027 4028 4037 4042 4047 4118 4130 4135 4137 4140 4142 4145 4147 4150 4320


Designation

AISI


4340 4422 4427 4720 4817 4820 50B40 50B44 5046 50B46 50B50 50B60 5120 5130 5132 5135 5140 5150 5155 5160 51B60 81B45 8630 8637 8640 8642 8645 86B45 8650 8660 8720 8740 8822 9260 94B15 94B17 94B30 94B40 9840 E4340 E9310 E50100 E51100 E52100 4023 4320 4620 4720 4820 E52100 1017 1018 1020 1022 1030 1035 1038 1006 1008 1010 1012 1015 1016 1018 1019 1021


Appendix 8:

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Designation

Country

AISI



1022 1026 1030 1035 1038 1030 1035 1038 1039 1040 1335 1340 4037 4137 4140 4142 4340 1016 1018 1019 1021 1022 1008 1010 1012 1015 1016 1017 1018 301 302 304 304L 305 309 309S 309S(Cb) 310S 316L 317 347 430 430Ti 630 631 632 634 422 615 661 1008 1010 1012 1015 1017 1020 1023 1025 1044 1008 1010 1012 1015 1016 1017 1018


Designation

AISI

ASTM A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A576 A579 A579 A579 (grade 61) A579 (grade 62) A579 (grade 63) A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A580 A581 A581 A581 A581 A581 A581 A581 (XM-2) A582 A582 A582 A582 A582 A582

1019 1020 1021 1022 1023 1025 1026 1030 1035 1038 1039 1040 1044 1045 1046 1050 1055 1060 1070 1080 1090 1095 1547 1548 632 634 633 631 431 302B 304 304L 305 308 309 309S 310S 314 316L 317 347 348 403 405 410 414 420 430 431 440A 440B 440C 446 303 303Se 416 416Se 430F 430F(Se) 303MA 303 303Se 416 416Se 420F(Se) 430F


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pg 580


Table A8.1

(continued)

Country


Designation

Country

AISI

ASTM A582 A582 (XM-2) A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A600 A632 A632 A632 A632 A632 A632 A633 A638 A639 A646 A646 A646 A646 A651 A651 A651 A651 A651 A651 (XM-8) A659 A659 A659 A659 A659 A659 A659 A666 A666 A666 A666 A666 A681 A681 A681 A681 A681 A681 A681 A681

430F(Se) 303MA M1 M2 M3 (Class 1) M3 (Class 2) M4 M6 M7 M10 M30 M33 M34 M36 M41 M42 M43 M44 M46 M47 T1 T2 T4 T5 T6 T8 T15 304 304L 310 316L 317 348 347 660 661 4130 4140 4340 E52100 304 409 430 430Ti 434 S43035 1015 1016 1017 1018 1020 1021 1023 201 202 301 302 304 A2 A3 A4 A5 A6 A7 A8 A9



Designation

AISI


A10 D2 D3 D4 D5 D7 H10 H11 H12 H13 H14 H19 H21 H22 H23 H24 H25 H26 H41 H42 H43 O1 O2 O6 O7 P2 P3 P4 P5 P6 P20 P21 S1 S2 S4 S5 S6 S7 1030 1035 1040 1045 1050 1055 1060 1064 1065 1070 1074 1080 1086 1095 W1 W5 304 304L 316L 630 631 632 633 634 630 631 632


Appendix 8:

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pg 581




Designation

Country

AISI


ASTM A705 A711 A711 A711 A711 A711 A711 A711 A711 B511 B512 B535 B536 B546

634 3140 4135 4720 8660 E9310 E50100 E51100 E52100 330 330 330 330 330

FED QQ-S-633 (C12L13) QQ-S-635 (C1030) QQ-S-635 (C1035) QQ-S-635 (C1045) QQ-S-635 (C1050) QQ-S-637 QQ-S-637 QQ-S-637 (C1008) QQ-S-637 (C1109) QQ-S-637 (C1116) QQ-S-637 (C1117) QQ-S-637 (C1118) QQ-S-637 (C1119) QQ-S-637 (C1132) QQ-S-637 (C1137) QQ-S-637 (C1144) QQ-S-637 (C1211) QQ-S-637 (C1212) QQ-S-637 (C1913) QQ-S-698 (C1008) QQ-S-698 (C1015) QQ-S-700 (C1025) QQ-S-700 (C1030) QQ-S-700 (C1035) QQ-S-700 (C1045) QQ-S-700 (C1050) QQ-S-700 (C1055) QQ-S-700 (C1065) QQ-S-700 (C1074) QQ-S-700 (C1080) QQ-S-700 (C1086) QQ-S-700 (C1095) QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763

12L13 1030 1035 1045 1050 1141 1215 1008 1109 1116 1117 1118 1119 1132 1137 1144 1211 1212 1213 1008 1015 1025 1030 1035 1045 1050 1055 1065 1074 1080 1086 1095 202 302 304 304L 305 309 310 316 316L 317 321 347 403 405 410 414 420 (continued)


Designation

AISI

FED QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-763 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-S-766 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-570 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590

430 440A 440B 440C 446 201 202 301 302 304 304L 309 310 316 316L 321 347 348 420 430 446 A2 A3 A4 A5 A6 A7 A8 A9 A10 D2 D3 D4 D5 D7 H10 H11 H12 H13 H14 H19 H21 H22 H23 H24 H25 H26 H41 H42 H43 O1 O2 O6 O7 S1 S2 S4 S5 S6 M1 M2 M3 (Class 1) M3 (Class 2) M4 M6 M7


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pg 582


Table A8.1

(continued)

Country


Designation

Country

AISI


FED QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-T-590 QQ-W-412 QQ-W-412 (II) QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-423 QQ-W-461 STD-66 STD-66 STD-66 STD-66

M10 M30 M33 M34 M36 M41 M42 M43 M44 M46 T1 T2 T4 T5 T6 T8 T15

302 304 305 310 316 321 347 410 416 420 430 1006 202 304 310 430

MIL SPEC MIL-C-24111 MIL-F-20138 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862 MIL-S-862

630 304 302 303 303Se 304 304L 309 310 316 316L 317 321 347 403 405 410 416 416Se 420 430 430F 430F(Se) 431 440A 440B 440C 440F 440F(Se) (continued)


Designation

AISI

MIL SPEC MIL-S-862 MIL-S-866 MIL-S-866 MIL-S-980 MIL-S-5000 MIL-S-5059 MIL-S-5059 MIL-S-5059 MIL-S-6049 MIL-S-7420 MIL-S-7493 (A4615) MIL-S-7493 (A4620) MIL-S-8503 MIL-S-11310 (CS1005) MIL-S-11310 (CS1006) MIL-S-11310 (CS1008) MIL-S-11310 (CS1010) MIL-S-11310 (CS1012) MIL-S-11310 (CS1017) MIS-L-11310 (CS1018) MIL-S-11310 (CS1020) MIL-S-11310 (CS1022) MIL-S-11310 (CS1025) MIL-S-11310 (CS1030) MIL-S-11310 (CS1040) MIL-S-11310 (ORD4150) MIL-S-11713 (2) MIL-S-16788 (C10) MIL-S-16974 MIL-S-16974 MIL-S-16974 (Gr. 1060) MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-16974 MIL-S-18411 MIL-S-18411 MIL-S-18733 MIL-S-20166

446 1016 8615 E52100 E4340 301 304 316 8740 E52100 4615 4620 6150 1005 1006 1008 1010 1012 1017 1018 1020 1022 1025 1030 1040 4150 1070 1095 1015 1050 1060 1080 1330 1335 1340 3140 4130 4135 4140 4145 4340 8620 8625 8630 8640 8645 1117 12L13 4135


Appendix 8:

21/8/2008 5:04PM Plate # 0

pg 583




Designation

AISI

MIL SPEC (CS1116) MIL-S-22141 MIL-S-22698 (B) MIL-S-23195 MIL-S-23195 MIL-S-23195 MIL-S-23195 MIL-S-23196 MIL-S-23196 MIL-S-23196 MIL-S-23196

1116 E52100 304 304L 347 348 304 304L 347 348


Country


Designation

AISI

MIL SPEC MIL-S-25043 MIL-S-46042 MIL-S-46049 MIL-S-46049 MIL-S-46409 MIL-S-81506 MIL-S-81591 MIL-T-6845 MIL-T-8504 MIL-T-8506 MIL-W-46078

631 651 1065 1074 1065 630 630 304 304 304 631

Name ///sr-nova/Dclabs_wip/Failure_Analysis/5113_585.pdf/Appendix_9/


18/8/2008 4:38PM Plate # 0


APPENDIX 9

Iron-Carbon Equilibrium Diagram

Fig. A9.1

Iron-carbon equilibrium diagram from Metal Progress Materials and Process Engineering Databook, American Society for Metals, 1968, p39

pg 585

Name ///sr-nova/Dclabs_wip/Failure_Analysis/5113_587-599.pdf/F5113_Appendix-10/


18/8/2008 4:39PM Plate # 0


APPENDIX 10

Isothermal Diagrams of Selected Steels A selection of isothermal diagrams for

Carbon steels (1019, 1030, 1050, 1060, 1080) Cr-Mo steels (4130, 4140) Ni-Cr-Mo steels (4340, 8620) Ni-Mo Steel (4640) Cr steel (5160, 52100)

REFERENCE

1. From Atlas of Isothermal Transformation and Cooling Transformation Diagrams, American Society for Metals, 1977

pg 587



Fig. A10.1

Carbon steels, 1019. Source: Ref 1

18/8/2008 4:40PM Plate # 0

pg 588


18/8/2008 4:40PM Plate # 0

Appendix 10: Isothermal Diagrams of Selected Steels / 589

Fig. A10.2


pg 589



Fig. A10.3


18/8/2008 4:40PM Plate # 0

pg 590


18/8/2008 4:40PM Plate # 0


Fig. A10.4


pg 591



Fig. A10.5


18/8/2008 4:40PM Plate # 0

pg 592


18/8/2008 4:41PM Plate # 0


Fig. A10.6

Chromium-molybdenum steels, 4130. Source: Ref 1

pg 593



Fig. A10.7

Chromium-molybdenum steels, 4140. Source: Ref 1

18/8/2008 4:41PM Plate # 0

pg 594


18/8/2008 4:41PM Plate # 0


Fig. A10.8

Ni-Cr-Mo steels, 4340. Source: Ref 1

pg 595



Fig. A10.9

Nickel-molybdenum steels, 4640. Source: Ref 1

18/8/2008 4:41PM Plate # 0

pg 596


18/8/2008 4:41PM Plate # 0


Fig. A10.10

Chromium steels, 5160. Source: Ref 1

pg 597



Fig. A10.11

Nickel-chromium-molybdenum steels, 8620. Source: Ref 1

18/8/2008 4:42PM Plate # 0

pg 598


18/8/2008 4:42PM Plate # 0


Fig. A10.12

Chromium steels, 52100. Source: Ref 1

pg 599


18/8/2008 4:47PM Plate # 0



APPENDIX 11

Continuous Cooling Diagrams of Selected Steels Selected continuous cooling transformation diagrams for:

Carbon steels with nominal carbon contents of 0.8, 0.44, 0.86, 0.96 wt% C Mn steels (13/4 M, 11/2 Mn) Mn-Mo, Mn-Ce, Mn-Ni-Mo, and Mn-Ni-CrMo steels Silicon steels (Si-Mn, Si-Cr, Si-Cr-Mo) 1/4 Mo steel Nickel steels (31/2 Ni, 11/2 Ni-Mn, 13/4 Ni-Mo, 11/4 NiCr)

Ni-Cr-Mo steels (1/2 Ni-Cr-Mo, 11/2 Ni-CrMo, 31/2 Ni-Cr-Mo) Chromium steels (1/2 Cr, 1Cr, 11/4 Cr-Mo, 1Cr-V, 11/2 Cr-Al-Mo)

REFERENCE

1. From Atlas of Continuous Cooling Transformation Diagrams for Engineering Steels, American Society for Metals, 1977

pg 601


18/8/2008 4:48PM Plate # 0

pg 602


C

Si

0.18 0.20

Mn

P

S

Cr

Mo

Ni

Al

Nb

V

0.45 0.020 0.020

900

800

Start

A

Ac3

10% 50%

700

Ac1

F

90%

Finish

P 600 °C

B 500

400

300

M

200

100 1000

500

200

100

50

Cooling rate at 800 °C

0 mm 0.1 Bar diameter

0.2 5

0.5

1 10

10

2

5

20

20

10

20

50

100

100 150 200 300

50 50

20

100

10

5

2

1

°C per min

150 200

300

200 500 500

500

1000 2000 mm Air

mm Oil mm Water

800 Hardness after transformation

700

60

600 500 HV 400

50 HRC 40

300 200

30 As cooled

100

Fig. A11.1

0.18 C (1017–1022), analysis wt%, austenitized at 900 C, previous treatment rolled

20 10


18/8/2008 4:48PM Plate # 0

pg 603

Appendix 11: Continuous Cooling Diagrams of Selected Steels / 603

Si

C

0.44 0.28

Mn

P

S

Cr

Mo

0.81 0.035 0.037 0.14 0.04

Ni

Al

Nb

V

Cu

0.15

Sn

0.12 0.016

900

800 Ac3 700

A

Start 10%

Ac1

50% 90%

F Finish P

600 °C

B

500

400

300

Ms

M 200

100

1000

500

200

100

50

0 mm 0.1

0.2

Bar 5 diameter 10 800

0.5

1

2

5 50

20

10 20

50

Hardenability band BS 970 080H46

10

20 100

100

20

10

5

50 150 200

150 200

100 300 300

200

500

500 HV 400

1

1000 2000 mm Air

500 mm Oil

500

mm Water Hardness after transformation

700 600

2

°C per min


As cooled T 500 °C 1 h T 550 °C 1 h T 600 °C 1 h T 650 °C 1 h T 700 °C 1 h

60

50 HRC 40

300

30

200

20 10

100

Fig. A11.2



18/8/2008 4:48PM Plate # 0

pg 604


Si

C

0.86 0.20

Mn

P

S

Cr

Mo

Ni

Al

Nb

V

0.60 0.020 0.020

900

800 Ac1a Ac1b

A + C

Start

700 P

Finish

600 °C 500

B

400

10% 50% 90%

300 Ms 200 10% 50%

M 100 90% 1000

500

200

100

50


0.2

5

0.5 10

10

1

2

5

20 20

10 50

50

20

20

50

100

100 150 200 300 100

10

5

2

1

°C per min


150

200 300

200

500

500 500

1000 2000 mm Air

mm Oil mm Water


700

60

600 500 HV 400

As cooled

50 HRC 40

300

30

200

20 10

100

Fig. A11.3



18/8/2008 4:49PM Plate # 0

pg 605


Si

C

Mn

0.96 0.20

P

S

Cr

Mo

Ni

AI

Nb

V

0.60 0.020 0.020

900

800

700

Ac1a Ac1b

A

C

Start P

Finish

600 °C

B

500

10% 50% 90%

400

300

200 Ms 10% 50%

100

M 1000 90%

500

200

100

50


20 10 °C per min

5

100

500

2

1


0.2

5

0.5

1

10 10

2

5

20 20

10

20

50 50

100 100

50

150 200 300 150 200 300

200 500 500

1000 2000 mm Air

mm Oil mm Water


700

60

600 As cooled 500 HV

T 400 °C 1 h T 500 °C 1 h

400

HRC 40

300

30 20 10

200 100

Fig. A11.4

50



18/8/2008 4:49PM Plate # 0

pg 606


Si

C

0.46 0.25

Mn

P

S

Cr

Mo

Ni

AI

Nb

V

1.80 0.020 0.015

900

800 Ac3 700

Ac1

A Start

600

10% 50% 90%

F Finish

°C

P

500

400

B

300 Ms

200 M 100 1000

500

200

100

50

20 10 °C per min



0.2

5

10 10

800

0.5

1

2

5

20 20

50 50

10 100

20

50

150 200

100 150 200

100 300

300

5

200

2

1

1000

2000 mm Air

500

500 mm Oil

500

Hardenability band SAE 1345H


700

60

600 50

500 HV

HRC 400 300

40 As cooled 30 20 10

200 100

Fig. A11.5

13/4 Mn (1547) (1345), analysis wt%, austenitized at 850 C, previous treatment rolled


18/8/2008 4:49PM Plate # 0

pg 607


Si

C

0.44 0.20

Mn

P

S

Cr

Mo

Ni

AI

Nb

V

1.50 0.020 0.250

900

800 Ac3 700

A

Ac1

Start °

10% 50% 90%

F

600 Finish °C

P

500

B

400

300 Ms

200 M 100 1000

500

200

100

50


20 10 °C per min

5

100

500

2

1

0 0.2

mm 0.1 Bar diameter

5

0.5 10

10

1

2

5

20 20

50 50

10

20

50

100 150 200 100

800

150 200

300 300

200 500

1000

2000 mm Air

mm Oil

500


700

60

600 50

500 As cooled

HV 400

40

300

30 20 10

200 100

Fig. A11.6

HRC

11/2 Mn+S (1139), analysis wt%, austenitized at 850 C, previous treatment rolled


18/8/2008 4:49PM Plate # 0

pg 608


C

Si

0.38 0.25

Mn

P

S

Cr

Mo

1.50 0.020 0.020

Ni

AI

Nb

V

0.45

900

800 Ac3

700

Ac1

Start

A

10%

F

P

600

50%

°C 500 90%

400 Finish

B Ms 300

200 M 100 1000

500

200

100

50

20

10

5

2

1

°C per min



0.2

5

0.5 10

10

1

2

5

20 20

10

50 50

100 100

20 150 200 150 200

50

100

200

300 300

500 500

500

1000

2000 mm Air

mm Oil mm Water

800 700 600 500


As cooled T 550 °C 1 h T 600 °C 1 h T 650 °C 1 h T 700 °C 1 h

HV

Hardness after transformation 60

50 HRC

400

40

300

30 20 10

200 100

Fig. A11.7

11/2 Mn Mo, analysis wt%, austenitized at 845 C, previous treatment rolled


18/8/2008 4:50PM Plate # 0

pg 609


Si

C

Mn

0.20 0.25

P

S

Cr

1.25 0.025 0.015 1.15

Mo

Ni

0.02

0.15

AI

Nb

V

900 Ac3 800 Start

A

Ac1

10%

700

50% 90%

F Finish 600 P

°C 500 B 400 M s

300 M 200

100 1000

500

200

100

50


0.2

5

0.5 10

10

1

2

5

20 20

50 50

10

20 100

100

20

10

5

2

1

°C per min


50 150 200

150 200

100 300 300

200 500 500

500

1000 2000 mm Air

mm Oil mm Water


700 600 500 HV

60

Hardenability band DIN 17210 20MnCr5

50 HRC

400

40

300 As cooled 200 100

Fig. A11.8

11/4 Mn Cr, analysis wt%, austenitized at 870 C, previous treatment rolled

30 20 10


18/8/2008 4:50PM Plate # 0

pg 610


C

Si

Mn

P

S

Cr

0.19 0.20 1.60 0.020 0.020

Mo

Ni

0.25

0.55

AI

Nb

V

900 Ac3 800 A 700

Start 10%

Ac1

50%

F

P

600 °C 500 90%

B Ms 400 Finish 300 M 200

100 1000

500

200 100

50



5 10

0.2

0.5 10

1

2

5

20 20

50 50

10 100 100

20

50

150 200 150 200

20 10 °C per min

5

100

500

300 300

200 500 500

2

1

1000 2000 mm Air

mm Oil mm Water


700

60

600 50 HRC

500 HV 400 300

40 As cooled

200 100

Fig. A11.9

11/2 Mn Ni Mo, analysis wt%, austenitized at 870 C, previous treatment rolled

30 20 10


18/8/2008 4:50PM Plate # 0

pg 611


Si

C

Mn

0.43 0.24

P

S

Cr

Mo

Ni

1.35 0.025 0.025 0.45 0.20

AI

Nb

V

0.75

900

800

700

Ac3 Ac1

A Start 10%

F

600

P

°C 500 50%

400

90%

B

300 Ms

Finish

200 M 100 1000

500

200 100

50

20

10

5

2

1

°C per min


0 0.2

mm 0.1 Bar diameter

5 10

0.5 10

1

2

5

20 20

50 50

10 100 100

20

50

150 200 150 200

100 300

300

800 700

200 500 500

500

1000

2000 mm Air

mm Oil mm Water Hardness after transformation 60

As cooled

600

T 550 °C 1 h

500

T 620 °C 1 h T 650 °C 1 h

40

300

30 20 10

200 100

Fig. A11.10

50 HRC

HV 400

11/2 Mn Ni Cr Mo, analysis wt%, austenitized at 850 C, previous treatment rolled and softened 650 C 1 h


18/8/2008 4:51PM Plate # 0

pg 612


Si

C

0.40 1.75

Mn

P

S

Cr

Mo

Ni

Al

Nb

V

0.85 0.030 0.030

900 Ac3 800

Start

A

Ac1

10% 50% 90%

700

Finish

F P

600 °C 500 B 400 Ms 300

200

M

100 1000

500

200

100

50

0 mm 0.1

0.2

0.5

1

2

5

10

20

10

5

2

1

°C per min


20

50

100

200

500

150 200

300

500

1000

2000 mm Air

Bar diameter

5

10 10

20 20

50 50

100 100

150 200

300

800 700

500

mm Oil mm Water

Hardness after transformation

Hardenability band NF A-35 41S7

60

600 500 HV 400

50 As cooled

40

300

30 20 10

200 100

Fig. A11.11

HRC

13/4 Si Mn, analysis wt%, austenitized at 910 C, previous treatment rolled


18/8/2008 4:51PM Plate # 0

pg 613


Si

C

Mn

0.45 3.40

900 Ac3

P

S

Cr

Mo

Ni

Al

Nb

V

0.60 0.015 0.010 8.50

Start 800 A + C

10%

P

50%

Finish

700

600 °C 500

400

300 B 200

Ms

100

M

1000

500

200

100

50

20 10 °C per min

50

100


5

2

1


0.2

5

0.5 10

10

1

2

5

20 20

50 50

10 100

20

150 200 300

100 150

200 300

200

500 500

500

1000

2000 mm Air

mm Oil mm Water

800 Hardness after transformation 60

700 600

As cooled 50 HRC

500 HV 400

40

300

30

200

20 10

100

Fig. A11.12

31/2 Si Cr, analysis wt%, austenitized at 1050 C, previous treatment rolled and softened 650 C 1 h


18/8/2008 4:51PM Plate # 0

pg 614


Si

C

0.60 1.90

Mn 0.85

P

S

Cr

0.025 0.025 0.30

Mo

Ni

Al

Nb

V

0.25

900

800

Ac3

Start

Ac1

10% 50% 90%

A 700

Finish P

600 °C 500

400 B 300 Ms 200

M

100

1000 500

200

100

50



0.2

5

0.5 10

10

1

2

5

20 20

50 50

10

20 100

100

50 150 200

150 200

20 10 °C per min 100 300

300

200

500

500 500

5

2

1

1000 2000 mm Air

mm Oil mm Water


60

600 50 HRC

500 HV 400 As cooled

40

300

30

200

20 10

100

Fig. A11.13

2 Si Cr Mo, analysis wt%, austenitized at 910 C, previous treatment rolled


18/8/2008 4:52PM Plate # 0

pg 615


Si

C

Mn

0.40 0.30

P

S

Cr

Mo

0.80 0.025 0.020

Ni

AI

Nb

V

0.26

900

800

Ac3 A

700

Start

Ac1

10% 50% 90%

F 600

P

Finish

°C B

500

400 Ms 300 M 200

100 1000

500

200

50

100

20

10

5

2

1

°C per min


0 0.2

mm 0.1 Bar diameter

5

0.5

1

10 10

2

5

20 20

10

50 50

20 100

100

50

100

200

150 200

300

500

150

200

300

500

1000

2000 mm Air

mm Oil

500

mm Water

800 700



60

600 500 HV 400

As cooled

50 HRC 40

300

30

200

20 10

100

Fig. A11.14

1/4

Mo (4037–4042), analysis wt%, austenitized at 810 C, previous treatment rolled


18/8/2008 4:52PM Plate # 0

pg 616


Si

C

Mn

0.40 0.26

P

S

Cr

0.62 0.007 0.005 0.23

Mo

Ni

0.10

3.45

AI

Nb

V

900

800 Ac3

A

700 Ac1 Start

600

10%

°C

F

50%

500 B 90%

400 Finish 300

Ms

200 M 100 1000 500

200

100

50

20 10 °C per min


5

2

1

1000

2000 mm Air

0 0.2

mm 0.1 Bar diameter

5

0.5 10

10

1

2

5

20 20

50 50

10 100 100

20

50

150 200 150 200

100 300

300

200 500 500

500

mm Oil mm Water


700

60

600 As cooled 500 HV 400

T 575 °C 1 h T 625 °C 1 h

50 HRC 40

300

30

200

20 10

100

Fig. A11.15

31/2 Ni, analysis wt%, austenitized at 860 C, previous treatment rolled


18/8/2008 4:52PM Plate # 0

pg 617


Si

C

0.16 0.25

Mn

P

S

Cr

1.40 0.020 0.015 0.20

Mo

Ni

0.05

1.50

AI

Nb

V

900

800 Ac3 A 700

Start

Ac1

10% 50%

600

F

°C 500 90%

B

400 Ms

h

Finis

300

M

200

100 1000

500

200

100

50


20 10 °C per min

5

100

500

2

1

0 0.2

mm 0.1 Bar diameter

5

0.5 10

10

1

2

5

20 20

50 50

10

20 100

100

150 150

50 200 200

300 300

200 500 500

1000

2000 mm Air

mm Oil mm Water


60

600 500 HV 400

50 As cooled

HRC 40

300

30

200

20 10

100

Fig. A11.16

11/2 Ni Mn, analysis wt%, austenitized at 840 C, previous treatment rolled


18/8/2008 4:53PM Plate # 0

pg 618


Si

C

0.40 0.15

Mn

P

S

Cr

0.48 0.016 0.040 0.15

Mo

Ni

0.25

1.75

AI

Nb

V

900

800 Ac3 700

A

Ac1

Start 10% 50% 90%

F 600

Finish

P

°C 500 B 400

300

Ms

200 M 100 1000

500

200 100

50

20 10 °C per min

50

100


5

2

1

0 0.2

mm 0.1 Bar 5 diameter

0.5 10

10

1

2

5

20 20

50 50

10

20

100 100

150

200 300

150 200

300

200 500 500

500

1000

2000 mm Air

mm Oil mm Water


700 600 500 HV 400

60 As cooled T 500 °C 1 h T 600 °C 1 h

50 HRC 40

300

30

200

20 10

100

Fig. A11.17

13/4 Ni Mo, analysis wt%, austenitized at 845 C, previous treatment rolled


18/8/2008 4:53PM Plate # 0

pg 619


C

Si

Mn

0.40 0.23

P

S

Cr

Mo

Ni

0.75 0.020 0.020 0.65

Al

Nb

V

1.30

900

800 Ac3 700

A Start

Ac1

10% 50% 90%

F Finish

600 P

°C 500

B

400

300

Ms

200 M 100 1000

500

200 100

50

20 10 °C per min

50

100


5

2

1

0 mm 0.1

0.2

Bar 5 diameter

0.5 10

10

1

2

5

20 20

50 50

10 100 100

20 150 150

200 200

300 300

200 500 500

500

1000

2000 mm Air

mm Oil mm Water


700 As cooled 600

T 550 °C 1 h T 600 °C 1 h

500

T 650 °C 1 h

50 HRC

HV 400

40

300

30

200

20 10

100

Fig. A11.18

11/4 Ni Cr, analysis wt%, austenitized at 850 C, previous treatment rolled


18/8/2008 4:53PM Plate # 0

pg 620


Si

C

0.24 0.20

Mn

P

S

Cr

0.80 0.020 0.020 0.50

Mo

Ni

0.20

0.55

AI

Nb

V

900

800

Ac3

Start

A

10%

Ac1

50%

700 90%

F 600

P Finish

°C 500 B 400 Ms

300 M 200

100 1000 500

200

100

50


20

10

5

2

500

1000

1

°C per min


0.2 5

0.5

1

10 10

2

5

10

20 20

50 50

20 100 100

50

100

200

150 200 300

500

150 200 300

500

2000 mm Air

mm Oil mm Water


700 600

60 Hardenability band BS 970 805H25

500

HV

50 HRC

400

40

300

30 As cooled

200 100

Fig. A11.19

1/2

Ni Cr Mo (8622–8627) (8720) (8822), analysis wt%, austenitized at 830 C, previous treatment rolled

20 10


18/8/2008 4:54PM Plate # 0

pg 621


Si

C

Mn

0.40 0.25

P

S

Cr

0.60 0.020 0.020 1.20

Mo

Ni

0.15

1.50

AI

Nb

V

900

800 Ac3 700

Ac1

A

F P

600 °C

Start

500

10%

400 Ms

90%

50%

B 300

Finish M

200

100 1000 500

200

100

50

Cooling rate at 700 °C 0 mm 0.1 Bar 5 diameter 10

0.2

0.5 10

1

2

20 20

5 50

50

10 100 100

20 150 200 150

20

10

5

2

1

°C per min 50 300

200 300

100 500 500

200

500

1000 2000 mm Air

mm Oil mm Water

800 700 600 500 HV 400

Hardness after transformation As cooled T 550 °C 1 h T 600 °C 1 h T 650 °C 1 h

60

50 HRC 40

300

30

200

20 10

100

Fig. A11.20

11/2 Ni Cr Mo, analysis wt%, austenitized at 850 C, previous treatment rolled, softened 650 C 1 h


18/8/2008 4:54PM Plate # 0

pg 622


Si

C

0.13 0.20

Mn

P

S

Cr

0.50 0.020 0.020 0.85

Mo

Ni

0.18

3.40

Al

Nb

V

900

800 Ac3

700

A

Ac1

Start 10% 50%

600 F

°C 500

400

90%

Ms

B

300

Finish M

200

100 1000

500

200

100

50

20


10

5

2

1

°C per min

0 mm 0.1

0.2

Bar 5 diameter 10

0.5 10

1

2

5

20 20

50 50

10

20 100

100

50

150 200 150 200

100 300 300

200

500

1000 2000 mm Air

500 500

mm Oil mm Water


700

60

600 500 HV 400

Hardenability band BS 970 832H13 As cooled

50 HRC 40

300

30

200

20 10

100

Fig. A11.21

31/2 Ni Cr Mo (9310), analysis wt%, austenitized at 820 C, previous treatment blank carburized 900 C 4 h A.C.


18/8/2008 4:54PM Plate # 0

pg 623


Si

C

Mn

0.59 0.25

P

S

Cr

Mo

Ni

0.60 0.025 0.025 0.65

Al

Nb

V

0.20

900

800 Ac3 Ac1

A Start

700 Finish

P 600 °C 10%

500

50%

90%

400 B 300 Ms 200

M

100

1000

500

200

100

50

Cooling rate at 700 °C 0 mm 0.1 Bar 5 diameter

0.2

0.5 10

10

1

2

20 20

5 50

50

10 100 100

20

10

5

2

1

°C per min 20 150 150

50 200 200

100 300 300

200

500

1000

2000 mm Air

500 500

mm Oil mm Water


700

60

600 500 HV


50 HRC

400 300

40 As cooled

200 100

Fig. A11.22

21/2 Cr (5060) (5155–5160), analysis wt%, austenitized at 830 C, previous treatment rolled

30 20 10


18/8/2008 4:55PM Plate # 0

pg 624


Si

C

Mn

0.39 0.20

P

S

Cr

Mo

Ni

Al

Nb

V

0.70 0.020 0.020 1.05

900

800 Ac3 A

Ac1

Start

700 Finish F

600

P °C 500

10%

B 50%

90%

400 Ms 300

200 M 100 1000

500

200

100

50

Cooling rate at 750 °C 0 mm 0.1

0.2

Bar 5 diameter 10

0.5 10

1 20

20

5

2 50 50

10

20 100

100

50 150 200

150 200

20 10 °C per min

5

100

500

300 300

200

2

1

1000 2000 mm Air

500 500

mm Oil mm Water

800 700 600 500 HV 400

Hardenability band ISO R 683 VII 3

As cooled T 500 °C 1 h T 550 °C 1 h T 600 °C 1 h T 650 °C 1 h T 700 °C 1 h

60

50 HRC 40

300

30 20 10

200 100

Fig. A11.23


1 Cr (5140), analysis wt%, austenitized at 870 C, previous treatment rolled


18/8/2008 4:55PM Plate # 0

pg 625


Si

C

0.42 0.25

Mn

P

S

Cr

Mo

0.85 0.020 0.020 1.15

Ni

Al

Nb

V

0.20

900

800

700

Ac3 Start

A

Ac1

10% 50% 90%

Finish

F 600 P

°C 500

400 B 300

Ms

200 M 100 1000 500 200 Cooling rate at 750 °C

100

50

20

10

5

2

1

°C per min

0 0.2


0.5 10

10

1

2

20 20

5 50

50

10 100

20

50

150 200

100 150

200

100 300 300

200

500

1000 2000 mm Air

500 500

mm Oil mm Water

800


700 600 500 HV 400

As cooled T 450 °C 1 h T 500 °C 1 h T 550 °C 1 h T 600 °C 1 h T 650 °C 1 h T 700 °C 1 h


300

50 HRC 40 30 20 10

200 100

Fig. A11.24

60

11/4 Cr Mo (4140–4142), analysis wt%, austenitized at 860 C, previous treatment rolled


18/8/2008 4:55PM Plate # 0

pg 626


Si

C

Mn

0.50 0.25

P

S

Cr

0.75 0.025 0.025 0.95

Mo

Ni

0.05

0.15

Al

Nb

V 0.20

900

800 Ac3 Ac1

Start

700

10%

A

90%

50%

P Finish

600 °C 500 B 400

300

Ms

M 200

100 1000 500 200 Cooling rate at 750 °C

100

50

10

20

5

2

1

°C per min

0 0.2


0.5 10

10

1

2

20 20

5

10

50 50

20

100 100

50

150 200

150

200

100 300

300

200

500

1000 2000 mm Air

500 500

mm Oil mm Water


60

600 500 HV 400

Hardenability band ISO R 683 XIV 13

40

300

As cooled

200 100

Fig. A11.25

50 HRC

1 Cr V, (6150), analysis wt%, austenitized at 875 C, previous treatment rolled

30 20 10


18/8/2008 4:56PM Plate # 0

pg 627


C

Si

Mn

0.42 0.30

P

S

Cr

0.65 0.020 0.020 1.65

Mo

Ni

Al

0.33

Nb

V

1.00

900

Start

800 Ac1

10% 50% 90%

A 700

Finish

P

F

600 °C 500

400

300

B Ms

200 M 100 1000 500

200

100

50

10

20

5

2

1

°C per min

Cooling rate at 800 °C 0 mm 0.1 Bar diameter

5

0.2

0.5 10

10

1

2

5

20 20

50 50

10 100 100

20 150 150

50 200 200

100

200

500

1000

2000 mm Air

300 300

mm Oil mm Water


60

600 500 HV 400

As cooled Hardenability band SIS 14 29 40

40

300

30 20 10

200 100

Fig. A11.26

50 HRC

11/2 Cr Al Mo, analysis wt%, austenitized at 900 C, previous treatment softened 650 C 1 h

Name ///sr-nova/Dclabs_wip/Failure_Analysis/5113_629-640.pdf/F5113_Index/

20/8/2008 5:06PM Plate # 0



Index A abrasive erosion, 129 abrasive wear, 129, 130(F), 137, 195(T), 315(F) acute-angled keyway, 11, 12(F) adhesion colloids, 138 adhesive wear, 129(F), 137, 315 aerospace applications (case studies) cracking in a main landing gear attach pin, 354–355, 357(F), 358(F), 359(F) failure analysis of a catapult holdback bar, 351–354(F), 355(F), 356(F) failure analysis of a main landing gear lever, 362–364(F), 365(F), 366(F) failure analysis of a nose landing gear (NLG) piston axle, 367–372(F), 373(F) failure analysis of an aircraft hoist sling during static test, 373–375(F), 376(F), 377(F) failure analysis of an inboard flap hinge bolt, 364–366(F), 367(F), 368(F), 369(F) failure analysis of an internal spur gear, 375–378(F), 379(F) fracture analysis of ASI 420 stainless steel roll pin, 359–362(F), 363(F) main landing gear (MLG) axle, 378–380(F), 381(F), 382(F) MLG linear actuating rod and cylinder, 355–359(F), 360(F), 361(F), 362(F) multiple-leg aircraft-handling sling, 372–373(F), 374(F), 375(F) nondestructive testing and failure analysis of fin attach bolts after full-scale fatigue testing, 380–392(F&T) aging tendency, 35–36 Agricultural Ammonia Institute, 75 AISI to non-AISI steel cross reference, 551–562(T) Aluminizing, 33 American Association of State Highway Transportation Officials/American Welding Society (AASHTO/ AWS), 503 American Iron and Steel Institute (AISI), 311(T) annealing austenitic stainless steels, 39 definition, 3(T) diffusion (homogenizing annealed), 3(T) double-annealing, 395–396 full, 3(T) incorrect, 107(F) intercritical ( partial), 3(T)

pg 629

isothermal, 3(T) powder, 395 recrystallization, 3(T) soft, 3(T) solution, 38, 39–40 spheroidized, 107(F) stress, 459 subcritical, 3(T), 16, 266–267 temper, 346 time range (normal), 40 transformation, 346 water, 285 anode, 131–132 anodic stress, 74(F) Association Francaise de Normalisation (AFNOR), 493 asymptotic, definition, 116 austenitic stainless steels, 4, 5(T), 36, 39 austenitizing temperatures for steel, 537–538(T) austenitizing temperatures for direct-hardening carbon and alloy steels (SAE), 537–538(T) reheating (austenitizing) temperatures for hardening of carburized carbon and alloy, 538(T) autofrettage, 66

B backscattered “Z” contrast, 113 bainite, 2, 209, 210(F), 279(F), 341, 403–404. See also lower bainite; upper bainite banding, 103(F), 278, 279(F), 353 beach marks, 77, 78(F), 97, 113, 118, 505, 506(F) Biot modulus, 6 blind holes, 14, 33(T), 152(F) blowholes, 155, 156 boost-diffuse cycle, 184–185 boriding, 33 brittle fracture, 30, 33–34, 53–59(F), 120–122(F) brittle temperature range (BTR), 517 brittle transgranular fracture, 295 buckling, 5, 30, 56, 88(F), 89, 128(F) bulk composition evaluation, 114

C carbides cementite, 222–223(F) film or flake, 227, 228(F)



20/8/2008 5:06PM Plate # 0


630 / Index

carbides (continued) globular, 223–225(F) network, 225–227(F&T), 228(F) overview, 222 spheroidal, 224 carbon steels coefficients of linear thermal expansion, 541–543(T) heat capacity, summary of, 548–549(T) physical properties of, 541–550(T) specific heats of, 549–550(T) thermal conductivities of, 547(T) thermal conductivity, summary of, 545–546(T) thermal expansion, summary of, 544–545(T) carbonitrided components carburizing and carbonitriding, comparison of, 178(T) case, definition, 183 case depth, 183, 209 case depth, improper, 208 case hardness, 209 core hardness, 209 core microstructure, 210–211, 212(F) dimensional stability, 200–204(F) overcarbonitriding, 211, 213, 214(F), 215(F) overview, 177–178(F&T) quenching, 191 quenching cracks, 205 residual stresses, 196 carbonitriding case hardening, 397–398(F) case studies, 411–412(F), 413–414(F) dimensional change in, 411–412(F) fatigue property characteristics, 246–250(F&T) process, 177–178(F&T) SAE 1118 steel, 491(F) steel selection, 181 stress, 196 carburized components carbides, 222–227(F&T) carburizing and carbonitriding, comparison of, 178(T) carburizing process, 184–185 case, definition, 183 case crushing, 231–232(F) case depth, 183–185(F), 186(F) case depth, improper, 207–208(F) case hardness, 209, 210(F) core microstructure, 210–211, 212(F) decarburization, 213, 215–217(F) design, 179–181(F) dimensional stability, 200–204(F) grain size, 217–219(F) internal oxidation, 219–222(F) macropitting, 230–231(F) micropitting, 230 noncarbide inclusions, 228–229(F) overcarburizing, 211, 213, 214(F), 215(F) overview, 177–178(F&T) partial melting, 233–234(T) pitting corrosion, 232–233(F) quenching, 185, 187–191

pg 630

quenching cracks, 204–207(F) residual stresses, 196–200(F&T) retained austenite, 191–196(F&T) steel selection and hardenability, 181–196(F&T) survey summary of sources of gear failures, 178(T) transition zone, 211, 213(F) carburized steels carbon potential, 216 case depth, 209 case structure, 218 corrosion resistance, 233 decarburization, 95(F) film or flake carbides, 227 gas-carburized, 199–200 incorrectly, 187(F) intergranular cracking, 96, 118 intergranular oxidation, 219, 220(F) network carbides, 225, 226 partial melting, 233–234(T) pitting, 226, 233 “puzzle piece”, 108(F) retained austenite, 191–196(F&T) spheroidized carbide particles, 224 case crushing, 231–232(F) casting process, failures clean steel, 165–166 cold joints, 163–165 component lifetime factors, 151 decarburization during microfusion, 162–163, 165(F) improper cast design, 151–153(F), 154(F) inclusions, 165–175(F&T) macroinclusions, 166–167 microinclusions, 167–168(T) porosity, effects due to, 154–162(F), 163(F) cathode, 131–132 cause-and-effect (CE) matrix, 400–401, 402–403(F) cavitation creep, 128 cementite, 222–223(F) definition, 1 characteristics vs. defect, 88 Charpy impact test, 49, 55, 58, 295(F), 297, 297(F) Charpy V-notch (CVN) impact toughness testing, 317(F), 343(F), 513 chevron marks, 48, 53–54(F), 119(T), 367, 370(F) chicken-wire pattern, 208 chromizing, 33 cladding, 39 clean steel, 165–166, 229 clogging, 167 coefficient of thermal expansion, 4–5 cold cracking, 36 cold forming, 44 cold joints, 163–165 cold junction, 163 cold work tools, heat treating failures of, 314–330(F&T) characteristics, 315–317(F) chemical composition, 314–317(F) design-related failures, 317–319(F) hardening temperatures, incorrect, 323–324, 325(F), 326(F)



20/8/2008 5:06PM Plate # 0

pg 631


Index / 631

heat treating processing, 323 incorrect EDM, 321–323(F) incorrect grinding, 319–321(F) large (blocky) carbides, 315 retained austenite content, excessive, 324, 326–329(F), 330(F) short or absent tempering, 329–330, 331(F), 332(F) surface damage, 319–323(F) tempering temperatures, incorrect, 323–324, 325(F), 326(F) component design design, overview, 1 design aspects general thermal stress, 7, 32 overview, 2, 4 peak stress, 7, 32, 38 primary stress, 7, 8, 32 residual stress, 7–8 secondary stress, 7, 32 distortion, techniques for controlling, 16–18 hardening, 2 heat treat processing, introduction to, 1–2(T) heat treated components, 8–14(F) heat treatment design, 29–31 heat treatment failures, examples, 18–29(F) heat treatment, modeling of, 31–33(F&T) heat-transfer theory, 4–7(F&T) quenching medium, 17 quenching techniques, 17–18 risk-based approach to, 40–41(F) steel grade/condition, 14–16 tempering, 2 welded components, 33–36 contact fatigue pitting. See macropitting contact loading, 95, 98(F), 206(F) continuous cooling diagrams, 601–627(F) 1/4 Mo steel, 615(F) carbon steels, 602–605(F) chromium steels, 623–627(F) Mn steels, 606–607(F) Mn-Cr steels, 609(F) Mn-Mo steels, 608(F) Mn-Ni-Cr-Mo steels, 611(F) Mn-Ni-Mo steels, 610(F) nickel steels, 616–619(F) Ni-Cr-Mo steels, 620–622(F) silicon steels, 612–614(F) continuous cooling transformation (CCT) diagrams, 32, 255–257(F), 339–340(F) contour hardening, 432, 470–471 corrosion crevice, 131 definition, 131 galvanic, 131–132 pitting, 131 stress-corrosion cracking (SCC), 131 types of, 131 corrosion fatigue, 100–101(F), 118, 120 corrosive wear, 129, 138 cotter holes, 152

crack blunts, 116 brittle, 92, 121(F), 275, 506 brittle fracture, 30 carbides, 315 circumferential, 360(F), 382 coloration, 60(F), 90 crack path, 96 crack-free welding, 507 ductile, 92, 120(F) ductile fracture, 33 extension, 56 faces, 44(F), 50, 63(F), 64(F) fatigue, 101, 171(F), 172(F), 173(F), 229(F), 242(F). See also fatigue crack initiation fatigue failures, 30 forging, 142 formation, 228 formulation, 99, 103, 143(F), 153(F) fracture, 90–97(F), 98(F), 118 grinding, 207–208(F) growth mechanism, 51 growth rates, 49, 50, 80 heat checking, 334(F) intergranular path, 306(F) lateral oxide formation, 221 longitudinal, 169, 170(F), 357, 360(F) macrobrittle, 92 macroductile, 92 mechanical, 348(F) metallography, 50–51 microscale crack path, 97 networks, 33 normalization, 153(F), 154(F) nucleation, 128, 130, 159, 160, 168(F), 170, 171(F) optical micrograph, 27(F) order, 48 origin, 21, 25(F), 90, 162(F) oxide filled, 104(F) pearlitic matrix, 53 pre-existing, 116 profile, 101(F) propagation, 301, 482, 506. See also crack propagation quench, 21, 24(F), 25(F), 27(F), 102(F), 154(F), 411(F). See also quench cracking resistance, 107 resistance to propagation, measuring, 34 secondary, 27(F), 50, 101(F), 274(F), 370, 372(F), 382(F) shank, 388, 389(F) sharp corner, 318(F) stage I, 78 stage II, 78 stage III, 78 surface, 482 temperature, 480, 481(F) through-the-wall, 76 tip, 78, 80, 112, 116, 126(F), 302 transgranular, 162(F) transgranular path, 304–305(F&T) vertical, 64(F)



20/8/2008 5:06PM Plate # 0


632 / Index

crack extension, 38, 56, 118 crack initiation carbides, 315 classifying, 118 electroplating, 45 fatigue crack initiation, 122–124(F) fatigue cycles, 482 fatigue failures, structural features of, 78 high-temperature grain-boundary oxidation, 281, 283(F) hot work tool steels, 344 inhibiting, 196 internal oxidation formation, 221 microcracks, 206 nitriding, 343 retard, 123 S-N curves, 482 subsurface, 93, 94(F), 229(F) crack propagation, 124–127(F) beach marks, 77, 78(F) brittle fracture, 121, 175 carbides, 280 cleavage crack propagation, 509 cleavage fracture, 54(F) definition, 482 ductile fracture, 120(F) ductile structural materials, 506 fatigue crack propagation, 162, 170, 172(F), 173(F) grain-boundary, 338, 341(F) hydrogen embrittlement, 301(F) intergranular brittle fracture, 121–122(F) nitrided tools, 344, 348(F) rate of, 80 resistance, 34 SCC fractures, 73 stress analyses, 116 cracking cold, 36 delayed. See underbead cracking fatigue, 506(F) ferrite vein, 512–513(F) flat, 82 HAC, 509–513(F&T) hydrogen-induced, 36 intergranular, 59, 63(F), 96(F), 100(F), 205–206(F), 279(F) microcracking, 183, 192, 206(F), 218 premature, 333–335, 336(F), 337(F), 338(F) quench, 59–65(F), 273–283(F&T) SCC, 28, 36, 70, 72–76(F&T), 77(F) solidification, 515–517(F) toe, 503(t), 504(t) underbead, 507, 510–511(F) creep, 128–129(F) creep rupture, 76–77(F), 78(F) critical flaw size, 49, 116 critical scuffing temperature, 230 cyclic fracture propagation, 80, 82(F) cyclic loading, 77, 118, 122(F)

pg 632

D damage, definition, 87 damage mechanism, definition, 113 damage mode, definition, 113 decarburization, 215–217(F) on carburized steel, 95(F) deformation, 88–90(F) buckling, 89 elastic, 89, 98(F) gradual onset, 88–89(F) plastic, 89 sudden, 88(F), 89 die-casting dies, 331, 334(F), 341, 343 dimensional stability distortion, 200–202(F) isothermal-transformation diagrams, 202(F) warpage, 200, 202 discontinuities, 133, 503–505(F&T) distortion bending, 468 buckling, 128(F) creep, 128–129(F) reasons for, 127–129(F) residual stresses, 129 techniques for controlling, 16–18 yielding, 127–128 double-frequency heating, 428 ductile crack path. See microvoid coalescence (MVC) ductile dimples. See microvoid coalescence (MVC) ductile fracture, 28(F), 51–53(F), 91(F), 92–93(F), 96, 120(F), 226–227

E elastic collapse. See buckling electrical discharge machining (EDM), 20–21 electrochemical theory, 514 electron-dispersive x-ray (EDX), 169 electroplating, 28–29, 45 embrittlement, 293–303(F&T) definition, 38 temper (two-step embrittlement), 296 tempered martensite embrittlement, (TME), 294–296(F) endogas, 177 energy-dispersive x-ray analysis, 113 energy-dispersive x-ray spectroscopy (EDS), 26 environmentally assisted failure. See corrosion etching, 11(F), 64(F), 104(F), 210(F) acid, 320 cooper sulfide, 303 corrosion resistance during, 337(F) nital, 171, 174(F), 216, 324, 325(F) retained austenite, 326, 327(F) stain, 113 temper, 45(F), 62(F), 357(F) Euler buckling, 30, 89 exogenous inclusions, 165, 229 exothermic sleeves, 158



20/8/2008 5:06PM Plate # 0

pg 633


Index / 633

F failure, definition, 40, 87 failure analysis, guidelines of background data, collection of, 111–112 chemical analysis, 114–115 conclusions, formulation of, 117 evidence, analysis of, 116–117 failed part, preliminary examination of, 112 failure mechanism, determination of, 113–114 fracture mechanics analysis, 90, 116 macroscopic examination/analysis, 112 mechanical testing, 112 metallographic specimens, 113 microchemical analysis, 114 microscopic examination/analysis, 112–113 nondestructive testing, 112 photographic examination, 112 the report, 117–118 specimens, 112 stress analysis, 114–115 testing, 116 failure analysis, stages of background information, collection of, 47 macroscopic examination, 50 mechanical testing, 49 metallography, 50–51(F) microscopic examination, 50(F) nondestructive testing, 48–49 preliminary visual examination, 47–48 specimens, collection/preservation of, 49–50 specimens, sectioning of, 50 failure mechanism, definition, 113 failure mode and effects analysis (FMEA), 40, 399 failure mode, definition, 113 failures, mechanisms/causes of, 43–86(F&T), 87–109(F) corrosion and environmental damage, 99–101(F) deformation, 88–90(F) design deficiencies, 43 failure analysis, 47–51(F) fracture mechanism, 51–83(F) fractures, 90–97(F) heat treating process, 106–108(F) heat treatment, poor response from, 101–103(F), 104(F) manufacture, 44–46(F) material issues, 43–44(F) processing, 44–46(F) service conditions, 46–47(F) wear, 97–99(F) fatigue, 77–83 fatigue crack initiation, 122–124(F) carburized steels, 185 inhibit, 177 structural surface anomalies, 199–200 fatigue fracture, 122–127 fatigue crack initiation, 122–124(F) fatigue crack propagation, 124–127(F) nitrided layers, 241–253(F&T) fatigue intrusions and extrusions, 80(F), 122(F)

fatigue resistance, 241–244(F&T) after nitriding/nitriding treatments, 242–244(F&T) fatigue striations aluminum alloy, 125(F) formation of, 80, 83(F) interstitial-free steel, 125(F) representative, 381(F), 390(F) stage II crack growth, 78(F), 81(F) typical, 354(F), 355(F), 356(F) feeders, 158 ferrite vein cracking, 512–513(F) fine grinding, 472–477(F) finite element analysis (FEA), 116 fisheyes, 115, 243(F), 511–512(F) flat cracking, 82, 83(F) forging, failure in case studies avoidance of flow through, lap, and crack, 145–148(F) crankcase underfill, 138–139(F) spade bit, 140–142(F) summarized, 133(T) trim tear, 142–143(F) tube bending, 139–140(F), 141(F) upset forging, 143–145(F), 146(F) discontinuities, 133 factors in analysis of cold forging failures, 134(T) factors in analysis of hot forging failures, 135(T) forging process design, 134–138(F&T) forging tolerances, 135–137 lubricant performance, 138 lubrication, 137, 138 wear types, 137–138 Fourier’s first law, 5 Fourier’s second law, 4 fracture mechanics analysis, 115–116 fracture mechanism, 51–83(F) brittle fracture, 53–59(F) ductile fracture, 51–53(F) fatigue, 77–83(F) intergranular brittle fracture, 59–77(F&T) fracture modes brittle, 28(F) ductile, 28(F) grain-boundary, 300 identification chart, 119(T) intergranular, 26, 28(F), 80, 82(F) microvoid coalescence (MVC), 51, 52(F), 507–508(F) TME, 65 transgranular, 80, 82(F), 296 fracture surface matching, 118 fracture toughness, 49, 55, 185, 522–523(T) fracture toughness testing, 49 fractures brittle, 96. See also brittle fracture corrosion-fatigue cracks, 118, 120 cyclic loading, 118 examining, 90–93(F) fracture modes, 19(T), 118 hydrogen embrittlement, 96. See also hydrogen embrittlement (HEM)



20/8/2008 5:06PM Plate # 0


634 / Index

fractures (continued) intergranular fracture, 96. See also intergranular brittle fracture macroscale brittle fracture, 91(F), 92(F) macroscale ductile fracture, 91(F), 93(F) macroscale features, 91–93(F) microscale features, 95–97(F) origin of, 118 process of, 118 river line features, 97(F) slant, 56. See also brittle fracture stress versus strength, 93–95(F) stress-corrosion cracks, 118, 120 striations, 118 fretting, 97, 98(F), 130–131(F)

G galling or seizing, 129 galvanic corrosion, 131–132 gap-by-gap hardening, 429, 430(F), 439(F), 470–471 gas bubble, 155, 167 gas holes (blowholes), 154, 155, 156 gas metal arc welding (GMAW), 511 gas porosity, 154–156(F) glow discharge optical emission spectroscopy, 220 gouging, 119(T) grain size, influence of, 217–219(F) grain-boundary sliding, 128(F) graphite products, 138 gray cast iron, 121(F) Great Boston Molasses Disaster, 56–57(F) grid crossings, 153(F) grinder burn, 98 grinding fatigue strength, 80 incorrect, 319–321(F) manufacture and processing, 44, 45(F) porosity, eliminating, 504 quench cracking, 62(F) residual stresses after, 472–477(F) stock removal, 16 surface damage, 319 surface oxidation, removing, 222 thermal defects, preventing, 208 grinding burn, definition, 207 grinding burns, 207, 208 grinding cracks, 207–208(F), 303–304(F) grinding grains, 473–474(F) Grossman H-values (numbers), 270(T)

H halos (or fisheyes), 115. See also fisheyes heat checking, 331, 334(F) heat checking cracks, 334(F) heat treat processing, introduction to, 1–2

pg 634

heat treated steel parts component characteristics, 104–106(F) corrosion damage, 99–101(F) deformation, 88–90(F) environmental damage, 99–101(F) fracture, 90–97(F) heat treatment, poor response to, 101–108(F) wear, 97–99(F) heat treater, 89, 91, 92, 104, 106, 107–108 heat treating process atmosphere, improper, 107–108(F) design engineers and, 106 errors, 106–107 heat treaters and, 106 heating rate, inadequate, 107–108(F) temperature, inadequate, 107–108(F) time at temperature, inadequate, 107–108(F) heat treatment component design process, 29–30 design, 29–31 failures, examples of, 18–29(F) heat treating errors, 18 heating errors, 18 material behavior, 30–31 modeling of, 31–33(F&T) phase 2 design review to avoid failures, 33(T) poor response to component characteristics, 104–106(F) raw material characteristics contributing to, 101–103(F), 104(F) solid-phase transformation model, 31–32 temperature errors, 18 thermomechanical modeling, 32–33 welded components, 36–40 heat-affected zone (HAZ), 35 heat-transfer theory, 4–7(F&T) Hertzian stresses, 97, 98(F), 206(F), 230 high-speed steels, 313–314(F) high-velocity oxy-fuel (HVOF) coating, 62 holes, types of, 152(F). See also individual types hot tops, 158 hot work tools, heat treating failures of, 330–349(F&T) characteristics of, 330–333(T), 334(F), 335(F) chemical composition, 330–333(T), 334(F), 335(F) cooling (slow) during quenching, 338–343(F) heat treating procedures, inadequate, 335–337(F) heating, excessive, 344–349(F) incorrect hardening temperatures, 337–338, 339(F), 340(F), 341(F), 343(F) incorrect tempering temperatures, 337–338, 339(F), 340(F), 341(F), 343(F) nitrided tools, 343–344(F), 348(F) premature cracking, 333–335, 336(F), 337(F), 338(F) hydrogen embrittlement (HEM) electroplating, 28–29, 45 intergranular brittle fracture, 59(F), 68–70(F), 71(F), 72(F), 73(F) TE, interaction with, 301–302(F) hydrogen traps, 508(T)


20/8/2008 5:06PM Plate # 0


pg 635


Index / 635

hydrogen-assisted cracking (HAC) examples failure to pass bend tests due to hydrogen, 513(F) ferrite vein cracking in high-heat output welds, 512–513(F) fisheyes on fracture surface, 511–512(F) underbead cracking, 510–511(F) ferrite vein cracking, 512–513(F) fisheyes, 511–512(F) hydrogen-assisted reduced ductility, 508(F), 513(F), 514(F) underbead cracking, 503(T), 504(F), 509–511 weld metal HAC, 511 hydrogen-assisted cracking theory, 506–509(F&T) hydrogen-assisted reduced ductility, 508(F), 513(F), 514(F) hydrogen-induced cracking, 36

I impact testpieces, 121(F) impact tests, usefulness of, 482 impact toughness testing, 49 inclusions case studies failure in the axle of a reduced section in a rotating component, 170–171, 172(F&T), 173 failure of a 52100 steel axle, 171, 173–175(F) failure of a steam turbine rotor blade, 168–170(F&T), 171(F) categories of, 53 classifications, 165 definition, 165 noncarbide inclusions, 228–229(F) stringer, 166, 280–281(F) indigenous inclusions, 165 induction coils materials for, 423–424(F) multiple-turn induction coil, 424–425(F) types of, 423 induction hardening fatigue strength, 481–485(F&T) fine grinding, 472–477(F) induction heating, 420–422(F), 440–444(F), 444–449(F&T) induction surface heating. See induction surface heating induction surface-hardened layer, 477–481(F&T) machine parts, 422–432(F), 485–491(F&T) magnetic flux concentrators, 437–440(F) overview, 417–419 quenching systems for, 449–452(F) steels, 419–420(F&T) stresses/residual stresses, time variation of, 452–466(F) surface hardening, 421–422, 466–472(F) induction heating coils for, 423(F) definition, 398 features of, 420–422(F) gear wheels, 429, 430(F), 431(F) Jominy curves, 4150 steel, 492, 493(F)

machine parts, 440–444(F) power supplies, 421 rotation velocities, 452 supervising, 456(F) temperature cycles, 456–457(F), 458(F) time-temperature dependence in, 444–449(F&T) use of, 417 induction scanning, 426, 427(F) induction surface hardening advantages of, 421–422 gears, 491–497(F) residual stresses, 472–477(F) workpiece distortion in, 466–472(F) induction surface heating, 430(F), 432–437(F), 442(F), 456, 457(F), 465(F), 471(F) induction surface-hardened layer, 444, 477–481(F&T) inhomogeneity, 462, 463(F), 479 Instron TT-DM machine, 242 intergranular brittle fracture causes of, 59, 279. See also individual causes creep process and, 128(F) creep rupture, 76–77(F), 78(F) hydrogen embrittlement, 59(F), 68–70(F), 71(F), 72(F), 73(F) liquid metal embrittlement, 66–67(F&T) quench cracking, 59–65(F) SAE 3161, 122(F) solid metal embrittlement, 67–68(F&T), 69(F) stress-corrosion cracking (SCC), 70, 72–76(F&T), 77(F) temper embrittlement, 65–66(F) TME and, 65(F), 295–296 intergranular cracking, 59, 63(F), 96(F), 100(F), 205–206(F), 279(F) interlath cleavage, 296 internal oxidation, 219–222(F) internal porosity, 158–159(F) International Organization for Standardization (ISO), 493 ion vapor-deposited (IVD), 362, 364(F), 365, 367(F), 369(F) iron-carbon equilibrium diagram, 585(F) isothermal diagrams carbon steels, 588–592(F) chromium steels, 597(F), 599(F) chromium-molybdenum steels, 593–594(F) nickel-chromium-molybdenum steels, 598(F) nickel-molybdenum steels, 596(F) Ni-Cr-Mo steels, 595(F) isothermal-transformation diagrams. See time-temperature transformation (TTT) diagrams

J jig-bore grinding, 2(F), 20 “job shop” heat treater, 87 Jominy distance (J-distance), 270

L lap, definition, 145–146, 277 lath martensite, 187, 188(F)


20/8/2008 5:06PM Plate # 0



636 / Index

Liberty ships, 54 linear thermal expansion, 5(T), 541(T), 544(T) liquid metal embrittlement (LME), 66–67(F&T), 302–303(T) liquid shrinkage, 157(F) liquidation temperature, definition, 233(T) low-alloy steels coefficients of linear thermal expansion, 541–543(T) cold forging, strain-limiting criteria for, 142(F) core microstructure, 210, 212(F) embrittlement, 302(T) HAZ maximum hardness, 35–36 heat capacity, summary of, 548–549(T) heat treatment procedures, 36–37 linear thermal expansion, 5(T) quench age embrittlement, 59 SCC and, 36, 73, 74(T) specific heats, 549–550(T) thermal conductivities, 547(T) thermal conductivity, summary of, 545–546(T) thermal expansion, summary of, 544–545(T) thermal hardening, 418 lower bainite, 2, 210(F), 270, 292, 299, 403(F) lubricant performance, 138 lubrication graphite products, 138 inadequate, 46–47(F) micropitting, 230 wear and, 137–138(F)

M machine parts, induction hardening of carburized parts, residual stresses in, 489–491(F&T) contour hardening, 432 double-frequency heating, 428 gap-by-gap, 429, 430(F), 439(F), 470–471 gear wheels, 427–432(F), 439(F) induction scanning, 426, 427(F) progressive hardening, 426–427(F) quenching, 442–444(F) quenching techniques, 425–426(F) scan-hardening process, 425(F) single-shot induction hardening, 422, 423, 424(F), 427, 428(F), 440(F), 496 single-shot induction heating, 424(F), 448(F) stress profiles, 485–491(F&T) tooth-by-tooth, 430(F), 439(F) tooth-gap hardening, 431–432(F) macroinclusions, formation of, 166–167 macropitting, 230–231(F) macroscale brittle fracture, 91(F), 92(F) magnetic flux concentrators, 437–440(F) martempering, 17–18, 60 martensite, 285–289(F) definition, 2 diffusionless process, 286 morphology of, 187(F) slipped martensite, 287 tempered martensite, 289

pg 636

mechanical cracks, 348(F) mechanism and mode, difference between, 113–114 mechanism of failure, definition, 87 metric conversion guide base SI units, 521(T) conversion factors, 522–523(T) derived SI units, 521–522(T) supplementary SI units, 521(T) microchemical analysis, 114 microcracking, 183, 192, 206(F), 218 microinclusions, formation of, 167–168(T) micropitting, 230 microscale fracture features, 95–97(F) microscopic examination, 50(F), 112–113 microvoid coalescence (MVC), 96 modulus of elasticity, 89, 529

N National Safety Transportation Board, 75 nil-ductility temperature (NDT), 55–56 niobium (formerly columbium), 39, 314 nitrided layers, fatigue fracture of carbonitriding, 246–250(F&T) fatigue evaluation, 244–246(F&T), 247(F&T) fatigue resistance, 241–244(F&T) nitrided steels, 244–246(F&T), 247(F&T) nitrocarburizing, 33 node dislocation, 152, 153(F) non-AISI to AISI steel cross reference, 563–583(T) noncarbide inclusions, 228–229(F) nondestructive testing, 48–49 dye-penetrant method, 48 eddy-current methods, 48 magnetic particle inspection, 48 ultrasonic testing (UT), 48–49 North American Die Casting Association (NADCA), 336–337 notch, 12(F), 55, 56, 70(F), 179(F), 266(F)

O Occupational Safety and Health Administration (OSHA), 40–41 overcarbonitriding, 211, 213, 214(F), 215(F) overcarburizing, 211, 213, 214(F), 215(F)

P pancake (flattened disc), forging, 144(F) Paris regime, 126–127(F) partial melting, 233–234(T) part-way downs, 146–147 passing hole, 152(F) pearlite crack/void formation, 53 decarburization on carburized steel, 95(F) definition, 1



20/8/2008 5:06PM Plate # 0

pg 637


Index / 637

phase transformations, 255–263(F&T) CCT diagrams, 255–257(F) distortion mechanism, 261 metallurgical crystal structure, 257(F), 258(F) residual stresses, relief of, 263(F) TTT diagrams, 255, 256(F) volume changes during, 261–263(F) volumetric change upon quenching, 257–261(F&T), 262(F) phosphorus, 55, 65, 115, 295–296, 297–298 phosphorus segregation, 185, 295–296, 297–298 pinholes, 154, 155–156(F), 404–405(F), 406(F) pitting, definition, 113–114 pitting corrosion, 232–233(F) plastic blunting process, 78, 80 plastic deformation avoiding, 8, 30 brittle fracture and, 34, 53, 54(F), 120 buckling and, 128 creep and, 128 ductile fracture and, 51, 96, 120 formation of, 7, 52, 286–287, 331, 446 fracture surface matching, 118 shot peening and, 199(F) plastic mold steels, 313(F) plate martensite, 187, 188(F) polishing, 51, 113, 313(F), 390(F), 482 porosity blowholes, treating, 156 casting component feeding, 159 effects due to, 154–162(F), 163(F) gas, 154–156(F) internal, from nucleation, 158–159 internal, from the surface, 158, 159(F) pinholes, treating, 156 shrinkage pores, 156–158(F) postweld heat treatment (PWHT), 37 powder metal steel components case hardening carbonitriding, 397–398(F) carburizing, 397 induction hardening, 398(T) steam treatment, 398–399(F) definition, 395, 396(F) failure analysis techniques, 399–401(F&T) CE matrix, 400–401, 402–403(F) failure mode and effects analysis (FMEA), 399 process maps, 400, 401(T) steps, 400(F) tools, 400(F) powder metallurgy process, 395–396, 397(F) powder metal steel failures, case studies blistered sintered products, 405–406 cracks after induction hardening, 409–410(F) cracks after quenching, 410–411(F) dimensional change in carbonitriding, 411, 412(F) dimensional change in steam treatment, 414(F), 415(F) dimensional instability during shrink fitting, 406–407 fracture of steam-treated part, 407–408(F)

low breaking load after carbonitriding, 413–414(F) low core hardness after steam treatment, 415 low surface hardness after carbonitriding, 411–412(F) oxidation after sintering, 408–409(F) pinholes after sintering, 404–405(F), 406(F) variation in bore diameter after heat treatment, 412–413 wear after sinter hardening, 401, 403–404(F) wear after sintering, 407(F), 408(F) powder metallurgy process blending, 396 compaction, 396 powder production, 395–396 sintering, 397 progressive hardening, 420, 426, 427(F), 446, 455, 469, 496. See also gap-by-gap hardening “puzzle piece” carbide, 108(F)

Q quench cracking, 59–65(F) definition, 59–60 intergranular fracture, 96 localized overheating, 62 martempering, 60 mitigation of, 60 stress raisers role in, 272–273 quench cracking, case studies as-quenched 4340 steel, 273–274(F) cracking of 4140 block forging after quenching and tempering, 274–275(F&T) decarburization and oxidized grain boundary, 281–283(F) network carbides and coarse grain size, 278–280(F) presence of a seam defect, 276(F) presence of chemical segregation, 278, 279(F) presence of slag inclusions and a lap defect, 276–278(F&T) presence of stringer inclusions and chemical segregation, 280–281(F&T) use of improper steel alloy and presence of voids in a steel brazed joint, 275–276(F&T) quench nonuniformity, 189 quenching carbonitrided components, 191 carburized components, 185, 187–191 component design, 263–265(F), 266(F&T) contamination, 190 dry die quenching, 18 failures due to, 255–284(F&T) immersion quenching, 451 induction hardening, 449–452(F) martempering, 17–18 phase transformations, 255–263(F&T) press quenching, 18 quench nonuniformity, 189 quenchant nonuniformity, 271–272(F) quenchant selection, 270–271(T) retained austenite, 191–196(F&T)



20/8/2008 5:06PM Plate # 0


638 / Index

quenching (continued) slack quenching, 278 spray quenching, 430(F), 444(F), 451–452 steel grade/condition, 266–270(F&T) atmosphere control, 269 component support/loading, 267(F), 268(F) heating control, 269 machining, 267 retained austenite, 269–270(T) surface condition, 267–269(F&T) quenching cracks, 204–207(F), 341

R ratchet mark, definition, 505–506 ratchet marks, 126(F), 388(F), 389(F), 506(F) recognized and generally accepted good engineering practice (RAGAGEP), 40 rehardening burn, 207 reoxidation, 166–167 residual stresses, 196–200(F&T), 472–477(F) retained austenite, 191–196(F&T), 269–270(T) risers, 158 river patterns, 54, 55(F), 97(F), 376(F) rolling, 130 hot rolling, 89 surface rolling, 122, 123(F), 124

S scanning electron microscope (SEM), 21, 50(F) shank crack, 388, 389(F) sharp asperity, 130(F) shear lip, 50, 51, 118, 119(T), 273 Sherlock Holmes rule, 114 shot blasting, 198 shot peening, 198–199(F), 200(T) shrinkage pores, 159, 160(F) examples: failure analysis of a mill gear, 159–161(F&T), 161–162(F), 163(F), 164(F&T) silicon, 56, 183, 223, 293, 296, 297(F), 334 siliconizing, 33 single-shot induction hardening, 422, 423, 424(F), 427, 428(F), 440(F), 496 single-shot induction heating, 424(F), 448(F) sintering ferrous PM parts, 397 oxidation after, 408–409(F) pinholes after, 404–405, 406(F) PM process map, 401(T), 402(F) wear after, 407, 408(F) slack quenching, 278 slant fracture, 56, 505 sliding, 56, 128(F), 129, 130(F), 195, 196(F), 199 slip band extrusions, 78, 80(F) slip band intrusions, 78, 80(F), 122 Smith curve, 247–248, 250(F)

pg 638

soft skin layers, 199–200 softening replica tape, 49–50 solid metal embrittlement, 67–68(F&T), 69(F) solid shrinkage, 158 solidification cracking, 515–517(F) solidification shrinkage, 157–158(F) solid-phase transformation model, 31–32 spalling failure, 20(F), 23(F), 138, 231, 322(F), 323 spheroidization, 107(F) spray quenching, 430(F), 444(F), 451–452 stainless steels, PWHT of, 37–40 austenitic chromium-nickel, 38–39 chromium steels, 37 duplex austenitic-ferritic chromium-nickel, 39–40 fully austenitic, 39 martensitic, 38 soft martensitic, 38 stabilized, 39 unstabilized austenitic, 39 static fatigue test, 301 static tests, 301 steady-state creep rate, 77, 139 steel hardness conversions approximate equivalent hardness numbers for Brinell hardness numbers for steel, 532–533(T) approximate equivalent hardness numbers for steel, 535(F) approximate equivalent hardness numbers for Vickers (diamond pyramid) hardness numbers for steel, 533–534(T) approximate Rockwell B hardness conversion numbers for nonaustenitic steels, 530(T) approximate Rockwell C hardness conversion numbers for nonaustenitic steels, 531(T) examples of published hardness conversion equations, 529(T) steel susceptibility ratio, 285 steel welds, failure analysis of discontinuities, 503–505(F&T) examples fatigue cracking of welded pipe flange, 506(F) lack of penetration, 505 lack of side-wall fusion, 505 porosity in weld metal, 504, 505(F) SCC of a weld, 514–515(F) solidification cracking of steel weld, 516(F), 517(F) fatigue, 505–506(F) hydrogen-assisted cracking (HAC), 509–513(F&T) hydrogen-assisted cracking theory, 506–509(F&T) solidification cracking, 515–517(F) stress-corrosion cracking (SCC), 513–515(F) steels alloyed, tempering, 293, 294(F&T), 295(F) carbon. See carbon steels carbon, isothermal diagrams 1019, 588(F) 1030, 589(F) 1050, 590(F) 1060, 591(F) 1080, 592(F)



20/8/2008 5:06PM Plate # 0

pg 639


Index / 639

carbon steels with nominal carbon contents 0.18 C (1017–1022), 602(F) 0.44 C (1039–1046), 603(F) 0.86 C (1080–1090), 604(F) 0.96 C (1090–1095), 605(F) carburized. See carburized steels chromium, 37–39 1 Cr (5140), 624(F) 1 Cr V, (6150), 626(F) 11/4 Cr Mo (4140–4142), 625(F) 11/2 Cr Al Mo, 627(F) 21/2 Cr (5060) (5155–5160), 623(F) 5160, 597(F) 52100, 599(F) chromium-molybdenum 4130, 593(F) 4140, 594(F) clean, 165–166, 229 cold work tool, 312(F) continuous cooling diagrams, 601–627(F) high-speed, 313–314 hot work tool, 312–313(F) induction hardening, 419–420(F&T) interstitial-free, 124, 125(F) isothermal diagrams, 587–599(F) low-alloy steels. See low-alloy steels Mn 11/2 Mn+S (1139), 607(F) 13/4 Mn (1547) (1345), 606(F) Mn-Cr, 11/4 Mn Cr, 609(F) Mn-Mo, 11/2 Mn Mo, 608(F) Mn-Ni-Cr-Mo, 11/2 Mn Ni Cr Mo, 611(F) Mn-Ni-Mo, 11/2 Mn Ni Mo, 610(F) Mo, 1/4 Mo (4037–4042), 615(F) nickel 1/2 Ni Cr Mo (8622–8627) (8720) (8822), 620(F) 11/4 Ni Cr, 619(F) 11/2 Ni Cr Mo, 621(F) 11/2 Ni Mn, 617(F) 13/4 Ni Mo, 618(F) 31/2 Ni, 616(F) 31/2 Ni Cr Mo (9310), 622(F) nickel-chromium-molybdenum, 8620, 598(F) nickel-molybdenum, 4640, 596(F) Ni-Cr-Mo, 4340, 595(F) nitrided, 244–246(F&T), 247(F&T) plastic mold, 313(F) as quenched 4340, 273–274(F) silicon 13/4 Si Mn, 612(F) 2 Si Cr Mo, 614 31/2 Si Cr, 613(F) tool, 311–314(F&T), 314–330(F&T) stress analysis, 115–116 concentrations, 32–33, 45, 46(F), 98(F) machine parts, 485–491(F&T) versus strength, 93–95(F) types of, 7–8 stress raisers, 179(F), 272–273

stress-corrosion cracking (SCC), 28 austenitic stainless steels, 36 intergranular brittle fracture, 70, 72–76(F&T), 77(F) stress-sorption theory, 514 striations, 78, 80, 97, 124, 506. See also fatigue striations stringer inclusion, 166, 280–281(F) stringers, 25(F), 103(F), 280, 281(F) structural flaw, 241 subcritical annealing, 16, 266–267 sudden onset deformation, 89 sulfidizing, 33 superposition principle, 127 surface (or skin) effect, 440 surface carbon content decarburization, 213, 215–217(F) overcarbonitriding, 211, 213, 214(F), 215(F) overcarburizing, 211, 213, 214(F), 215(F) surface fatigue wear, 138 surface integrity, 418, 475 surface rolling, 122, 123(F), 124 surface structure anomalies, 199–200 Systeme International d’Unite´s (SI), 521

T TBE, 299 temper colors for steels color chart (back cover) time-temperature effect, 539, 540(F) temper embrittlement (TE), 296–299(F) HEM, interaction with, 301–302(F) LME, interaction with, 302–303(T) mechanical tests for, 299–301(F&T) mechanism of, 65–66(F) molybdenum, effect on, 297–298 phosphorus, effect on, 297–298 vanadium, effect on, 298 welded components, 34–35 temperature conversion table, 525–528 tempered martensite embrittlement, (TME), 65(F), 294–296(F) tempering alloyed steels, 293, 294(F&T), 295(F) case study: grinding cracks, 303–304(F) case study: transgranular and intergranular crack path, 304–305(F&T), 306(F) colors of tempering heats, 289(T) embrittlement, 293–303(F&T) heating times, 290(T) martensite, 285–289(F) mechanical properties, effect on, 289–290(F), 291(F), 292(F) reactions, 290–291 stages, 291–293(F) tempering embrittlement, 285 tempering resistance, 36(F), 313, 331, 334, 338, 339 tensile, definition, 92 thermal expansion austenite versus ferrite, 5


20/8/2008 5:06PM Plate # 0



640 / Index

thermal expansion (continued) austenitic stainless steels, 39 carbon and low-alloy steels, specific heats of, 549–550(T) coefficients of, 4–5, 6(F), 9(F), 32 coefficients of linear thermal expansion for carbon and low-alloy steels, 541–543(T) ferrous materials, 5(T) heat capacity, summary of, 548–549(T) induced strain, 7 residual stress, influence on, 5, 32 summary of, 544–545(T) thermal conductivities of carbon and low-alloy steels, 547(T) thermal conductivity, summary of, 545–546(T) thermal shock, 5–6(T), 32 thermochemical modeling, 33 thermomechanical modeling, 32–33 through hardening, 395, 396, 413, 421, 425(F), 442(F), 490 through-hardened, 413(F), 484, 485(F) time-temperature transformation (TTT) diagrams, 2, 32, 202(F), 255, 256(F) time-temperature-austenitizing (TTA) diagram, 432(F), 433(F) toe cracking, 503(T), 504(F) tool steels, failure analysis in classification of, 311–314(F&T) cold work, 312(F) cold work tools, 314–330(F&T) high-speed steels, 313–314(F) hot work, 312–313(F) hot work tools, 330–349(F&T) plastic mold steels, 313(F) tooth-by-tooth induction hardening, 430(F), 439(F) tooth-gap hardening, 431–432(F) torsional fatigue, 177, 483 total oxidation potential (TOP), 221(F) total-life fatigue analysis, 177 transition temperature, 35, 55, 66, 67(F) transmission electron microscopy, 65, 220, 288 tribological systems, 475 twinning density, 288 two-step embrittlement, 296

pg 640

U ultrasonic testing (UT), 512 underbead cracking, 507, 510–511(F) undercut, 20(F), 46(F), 503(T) undercutting, 45–46(F) unnotched impact testing, 337 upper bainite, 2, 270, 306(F), 403(F) upset forging, 143–145(F&T), 146(F)

V volume wear, 138

W wear, 97–99(F) wear, definition, 129 wear-assisted failure, 129–131(F) welded components, failure aspects of aging tendency, 35 austenitic stainless steels, SCC of, 36 brittle fracture, 33–34 cold cracking, 36 hardening tendency, 35–36 segregation tendency, 36 temper embrittlement, 34–35 welded components, heat treatment procedures austenitic-ferritic dissimilar joints, 40 carbon steels, 36–37 low-alloy steels, 36–37 stainless steels, PWHT of, 37–40

Y yield point, 483(T) yield strength, 144(T), 161(T), 250(T), 507 yielding, 127–128

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