Creep-resistant steels
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Creep-resistant steels Edited by Fujio Abe, Torsten-Ulf Kern and R. Viswanathan
Woodhead Publishing and Maney Publishing on behalf of The Institute of Materials, Minerals & Mining WPNL2204
CRC Press Boca Raton Boston New York Washington, DC
WOODHEAD
PUBLISHING LIMITED
Cambridge England
WPNL2204
iv Woodhead Publishing Limited and Maney Publishing Limited on behalf of The Institute of Materials, Minerals & Mining Woodhead Publishing Limited, Abington Hall, Abington Cambridge CB21 6AH, England www.woodheadpublishing.com Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW, Suite 300, Boca Raton, FL 33487, USA First published 2008, Woodhead Publishing Limited and CRC Press LLC © 2008, Woodhead Publishing Limited The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publishers cannot assume responsibility for the validity of all materials. Neither the authors nor the publishers, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. Library of Congress Cataloging in Publication Data A catalog record for this book is available from the Library of Congress. Woodhead Publishing ISBN 978-1-84569-178-3 (book) Woodhead Publishing ISBN 978-1-84569-401-2 (e-book) CRC Press ISBN 978-1-4200-7088-0 CRC Press order number: WP7088 The publishers’ policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elementary chlorine-free practices. Furthermore, the publishers ensure that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Replika Press Pvt. Ltd. India Printed by TJ International Limited, Padstow, Cornwall, England
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Contents
Contributor contact details
xiii
Preface
xix
Part I General 1
Introduction
3
F. ABE, National Institute for Materials Science (NIMS), Japan
1.1 1.2 1.3 1.4 1.5 1.6
Definition of creep Creep and creep rate curves Creep rupture data Deformation mechanism map Fracture mechanism map References
3 3 7 9 11 14
2
The development of creep-resistant steels
15
K.-H. MAYER, ALSTOM Energie GmbH, Germany and F. MASUYAMA, Kyushu Institute of Technology, Japan
2.1 2.2 2.3 2.4 2.5
15 18 19 42
2.6 2.7
Introduction Requirements for heat-resistant steels Historical development of ferritic steels Historical development of austenitic steels Historical development of steel melting and of the purity of heat-resistant steels Summary References
3
Specifications for creep-resistant steels: Europe
78
64 67 70
G. MERCKLING, RTM BREDA Milano, Italy
3.1
Introduction
78
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Contents
3.2 3.3 3.4
3.6 3.7
Specifications and standards The European Creep Collaborative Committee (ECCC) European Pressure Equipment Research Council (EPERC) The latest generation of CEN standards for creep-resistant steels Future trends References
95 150 151
4
Specifications for creep-resistant steels: Japan
155
3.5
81 85 92
F. MASUYAMA, Kyushu Institute of Technology, Japan
4.1 4.2 4.3 4.4 4.5 4.6 4.7
Introduction Types of heat-resistant steels in Japan Specifications for high temperature tubing and piping steels Specifications for steam turbine steels Heat-resistant super alloys Summary References
155 155 158 169 169 169 173
5
Production of creep-resistant steels for turbines
174
Y. TANAKA, Japan Steel Works, Japan
5.1 5.2 5.3 5.4 5.5
Introduction Overview of production technology of rotor shaft forgings for high temperature steam turbines Production and properties of turbine rotor forgings for high temperature applications Future trends References
174 175 192 207 212
Part II Behaviour of creep-resistant steels 6
Physical and elastic behaviour of creep-resistant steels 217 Y. YIN and R.G. FAULKNER, Loughborough University, UK
6.1 6.2 6.3 6.4 6.5 6.6 6.7
Introduction 217 Elastic behaviour 219 Thermal properties of creep-resistant steels 225 Electrical resistivity and conductivity of creep-resistant steels 234 Implications for industries using creep-resistant steels 238 Future trends 239 References 239
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Diffusion behaviour of creep-resistant steels
vii
241
H. OIKAWA and Y. IIJIMA, Tohoku University, Japan
7.1 7.2 7.3 7.4 7.5 7.6 7.7 7.8 8
Introduction Diffusion and creep Diffusion characteristics Roles of atom/vacancy movement in creep Influence of some factors on creep through their effects on diffusion Diffusion data in iron and in some iron-base alloys Concluding remarks References Fundamental aspects of creep deformation and deformation mechanism map
241 241 243 248 250 255 260 263
265
K. MARUYAMA, Tohoku University, Japan
8.1 8.2 8.3 8.4 8.5 8.6 8.7 8.8 8.9
Introduction Stress–strain response of materials Temperature and strain rate dependence of yield stress Deformation upon loading of creep test Creep behavior below and above athermal yield stress Change in creep behavior at athermal yield stress σa Deformation mechanism maps Concluding remarks References
265 265 267 269 270 271 275 278 278
9
Strengthening mechanisms in steel for creep and creep rupture
279
F. ABE, National Institute for Material Science (NIMS), Japan
9.1 9.2 9.3 9.4 9.5 9.6 10
Introduction Basic ways of strengthening steels at elevated temperature Strengthening mechanisms in modern creep-resistant steels Loss of strengthening mechanisms in 9–12Cr steels during long time periods Future trends References Precipitation during heat treatment and service: characterization, simulation and strength contribution
279 279 287 295 301 301
305
E. KOZESCHNIK and I. HOLZER, Graz University of Technology, Austria
10.1
Introduction
305
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Contents
10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9
Microstructure analysis of the COST alloy CB8 Modelling precipitation in complex systems Computer simulation of the precipitate evolution in CB8 Microstructure–property relationships The back-stress concept Loss of precipitation strengthening during service of CB8 Summary and outlook References
306 312 315 320 322 324 325 326
11
Grain boundaries in creep-resistant steels
329
R.G. FAULKNER, Loughborough University, UK
11.1 11.2 11.3 11.4 11.5 11.6 12
Introduction Ferritic steels Austenitic steels Grain boundary properties and constitutive creep design equations Future trends References Fracture mechanism map and fundamental aspects of creep fracture
329 330 341 345 346 347
350
K. MARUYAMA, Tohoku University, Japan
12.1 12.2 12.3 12.4 12.5
12.7 12.8 12.9 12.10
Introduction Fracture mechanisms and ductility of materials Stress and temperature dependence of rupture life Fracture mechanism maps Influence of fracture mechanism change on creep rupture strength Influence of microstructural degradation on creep rupture strength Change in creep rupture properties at athermal yield stress Multi-region analysis of creep rupture data Summary References
358 359 361 362 364
13
Mechanisms of creep deformation in steel
365
12.6
350 351 352 355 356
W. BLUM, University of Erlangen-Nuernberg, Germany
13.1 13.2 13.3 13.4
Introduction Initial microstructure Creep at constant stress Transient response to stress changes
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Contents
13.5 13.6 13.7 13.8 13.9 13.10 13.11 13.12 14
Cyclic creep Microstructural interpretation of creep rate Dislocation models of creep In situ transition electron microscope observations of dislocation activity Discussion and outlook Acknowledgments References Appendix: Microstructural model Mikora Constitutive equations for creep curves and predicting service life
ix
374 375 385 389 393 395 395 401
403
S.R. HOLDSWORTH, EMPA – Materials Science & Technology, Switzerland
14.1 14.2 14.3 14.4 14.5 14.6 14.7 14.8
Introduction Constitutive equations Constitutive equation selection Predicting service life Future trends Concluding remarks Nomenclature References
403 405 405 412 416 416 416 417
15
Creep strain analysis for steel
421
B. WILSHIRE and H. BURT, University of Wales Swansea, UK
15.1 15.2 15.3 15.4 15.5 15.6
Introduction Creep-induced strain Patterns of creep strain accumulation Practical implications of creep strain analysis Future data analysis options References
421 422 427 433 441 442
16
Creep fatigue behaviour and crack growth of steels
446
C. BERGER, A. SCHOLZ, F. MUELLER and M. SCHWIENHEER, Darmstadt University of Technology, Germany
16.1 16.2 16.3 16.4 16.5 16.6
Introduction Creep–fatigue experiments Stress–strain behaviour Creep–fatigue interaction, life estimation Multiaxial behaviour Creep and creep–fatigue crack behaviour
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Contents
16.7 16.8 16.9
Concluding remarks Acknowledgements References
468 469 469
17
Creep strength of welded joints of ferritic steels
472
H. CERJAK and P. MAYR, Graz University of Technology, Austria
17.1 17.2 17.3 17.4 17.5 17.6 17.7 17.8 18
Introduction Influence of weld thermal cycles on the microstructure of ferritic heat-resistant steels Weld metal development for creep-resistant steels Creep behaviour of welded joints Selected damage mechanism in creep-exposed welded joints Implications for industries using welded creep-resistant steels Future trends References
472
495 496 498
Fracture mechanics: understanding in microdimensions
504
474 482 483 484
M. TABUCHI, National Institute for Materials Science (NIMS), Japan
18.1 18.2 18.3 18.4 18.5 18.6
Introduction Non-linear fracture mechanics Effect of mechanical constraint Effect of microscopic fracture mechanisms Type IV creep crack growth in welded joints References
504 504 507 509 513 517
19
Mechanisms of oxidation and the influence of steam oxidation on service life of steam power plant components
519
P. J. ENNIS and W. J. QUADAKKERS, Forschungszentrum Juelich GmbH, Germany
19.1 19.2 19.3 19.4 19.5 19.6
Introduction Mechanisms of enhanced steam oxidation Steam oxidation rates Oxidation and service life Development of steam oxidation-resistant steels Outlook
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Contents
19.7 19.8
Sources of further information References
xi
534 534
Part III Applications 20
Alloy design philosophy of creep-resistant steels
539
M. IGARASHI, Sumitomo Metal Industries, Japan
20.1 20.2 20.3 20.4
Introduction Creep-resistant steels for particular components in power plants and the properties required Alloy design philosophies of creep-resistant steels References
539 539 541 570
21
Using creep-resistant steels in turbines
573
T.-U. KERN, Siemens AG Power Generation Group, Germany
21.1 21.2 21.3 21.4 21.5 21.6
Introduction Implications for industries using creep-resistant steels Improving the performance and service life of steel components Next steps into the future Summary References
573 574 583 591 593 593
22
Using creep-resistant steels in nuclear reactors
597
S.K. ALBERT, Indira Gandhi Centre for Atomic Research, India and S. SUNDARESAN, Maharaja Sayajirao University, Baroda, India
22.1 22.2 22.3 22.4 22.5 22.6 22.7
Introduction Radiation damage Embrittlement caused by ageing Use of heat-resistant steels in major reactor types Fabrication and joining considerations Summary References
597 598 611 613 629 631 632
23
Creep damage – industry needs and future research and development
637
R. VISWANATHAN and R. TILLEY, Electric Power Research Institute, USA
23.1 23.2
Introduction Calculational methods for estimating damage
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Contents
23.3 23.4 23.5 23.6 23.7
Non-destructive evaluation methods Accelerated destructive tests High temperature crack growth Future trends References
643 653 658 662 663
Index
667
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Contributor contact details
(* = main contact)
Editors
Chapter 1
Fujio Abe* Structural Metals Center, National Institute for Materials Science 1-2-1 Sengen Tsukuba 305-0047 Japan
Fujio Abe Structural Metals Center, National Institute for Materials Science 1-2-1 Sergen Tsukuba 305–0047 Japan
Email:
[email protected]
Email:
[email protected]
T.-U. Kern Siemens AG, Power Generation Group Dept. Materials Rheinstr. 100 D-45478 Muelheim Germany
Chapter 2 K.-H. Mayer* Am Kirchbühl 1 D-90592 Schwarzenbruck Germany
Email:
[email protected]
R. (Vis)Viswanathan Technical Executive Electric Power Research Institute 3420 Hillview Ave Palo Alto CA 94304 USA
Fujimitsu Masuyama Dept. Applied Science for Integrated System Engineering Kyushu Institute of Technology 1-1 Sensui-cho Tobata Kitakyushu 804-8550 Japan Email:
[email protected];
[email protected]
Email:
[email protected]
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Contributor contact details
Chapter 3
Chapter 6
Gunther Merckling Via Po 84 20032 CORMANO MI Milano Italy
Y.F. Yin and R.G. Faulkner IPTME Loughborough University Ashby Road Loughborough LE11 3TU UK
Email:
[email protected]
Email:
[email protected]
Chapter 4
Chapter 7
Fujimitsu Masuyama Dept. Applied Science for Integrated System Engineering Kyushu Institute of Technology 1-1 Sensui-cho Tobata Kitakyushu 804-8550 Japan
Hiroshi Oikawa* 2-2 Kagitori-3 Sendai 982-0804 Japan
Email:
[email protected]
Email:
[email protected]
Yoshiaki Iijima 37-2 Kamo-1 Sendai 981-3122 Japan Email:
[email protected]
Chapter 5 Yasuhiko Tanaka The Japan Steel Works 4 Chatsu, Muroran, Hokkaido 051-8505 Japan
Chapter 8
Email:
[email protected]
Kouichi Maruyama Graduate School of Environmental Studies Tohoku University 6-6-02 Aobayama Aoba-ku Sendai 980-8579 Japan Email:
[email protected]
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Chapter 9
Chapter 12
Fujio Abe Structural Metals Center, National Institute for Materials Science 1-2-1 Sergen Tsukuba 305–0047 Japan
Kouichi Maruyama Graduate School of Environmental Studies Tohoku University 6-6-02 Aobayama Aoba-ku Sendai 980-8579 Japan
Email:
[email protected] Email:
[email protected],jp
Chapter 10 Ernst Kozeschnik* and Ivan Holzer Institute for Materials Science, Welding and Forming Graz University of Technology Kopernikusgasse 24 A-8010 Austria Email:
[email protected];
[email protected]
Chapter 13 Wolfgang Blum Department of Materials Science Engineering Institute I: General Materials Properties WWI University of Erlangen-Nuernberg Martensstr. 5 91058 Erlangen Germany
Chapter 11 R.G. Faulkner IPTME Loughborough University Loughborough LE11 3TU UK Email:
[email protected]
Email:
[email protected]
Chapter 14 Stuart Holdsworth EMPA – Materials Science & Technology Überlandstrasse 129 CH-8600 Dübendorf Switzerland Email:
[email protected]
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Contributor contact details
Chapter 15
Chapter 18
B.Wilshire* and H. Burt Materials Research Centre School of Engineering University of Wales Swansea Singleton Park Swansea SA2 8PP UK
Masaaki Tabuchi Materials Reliability Centre National Institute for Materials Science (NIMS) 1-2-1 Sengen Tsukuba 305-0047 Japan
Email:
[email protected]
Email:
[email protected]
Chapter 16
Chapter 19
Christina Berger*, Alfred Scholz, F. Mueller, Michael Schwienheer Institute for Materials Technology Darmstadt University of Technology Grafenstr 2 64283 Darmstadt Germany
P. J. Ennis and W. J. Quadakkers Forschungszentrum Juelich GmbH IEF-2 D 52425 Juelich Germany
Email:
[email protected];
[email protected];
[email protected]
Chapter 17 Horst Cerjak* and Peter Mayr Institute for Materials Science, Welding and Forming Graz University of Technology Kopernikusgasse 24 A-8010 Austria
Email:
[email protected]
Chapter 20 M. Igarashi Corporate Research and Development Laboratories Sumitomo Metal Industries Ltd. 1-8 Fuso-cho Amagasaki Hyogo 660-0891 Japan Email:
[email protected]
Email:
[email protected];
[email protected];
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Chapter 21
Chapter 23
T.-U. Kern Siemens AG Power Generation Group Dept. Materials Rheinstr. 100 D-45478 Muelheim Germany
R. (Vis) Viswanathan* and Richard Tilley Technical Executive Electric Power Research Institute 3420 Hillview Ave Palo Alto CA 94304 USA
Email:
[email protected]
Email:
[email protected]
Chapter 22 S. Sundaresan* L&T Visiting Welding Chair Dept. of Metallurgical Engineering Faculty of Technology and Engineering Maharaja Sayajirao University Kala Bhavan Baroda-390001 India S.K. Albert Indira Gandhi Centre for Atomic Research Kalpakkam 603102 India Email:
[email protected] [email protected]
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Preface
Creep-resistant steel that can be used for a long time at elevated temperature is the key to the construction of thermal and nuclear power generation plants, chemical plants and petroleum plants. During the last decade, great progress has been made in developing creep-resistant steels of high strength and corrosion resistance at ever increasing temperatures and in evaluating the steels in terms of the weld characteristics, creep strength and corrosion resistance necessary for constructing plants. Although in the past the driving force for these developments has been primarily to achieve higher efficiencies, the focus has shifted more recently to the reduction of emissions of CO2, dioxins and other environmentally hazardous gases. In the field of thermal power generation, the maximum allowable temperature was about 565°C for conventional low alloy ferritic steels. However, progress in recent years has led to the development of high-strength 9–12% Chromium ferritic steels capable of operating in ultra super critical (USC) power plants at metal temperatures approaching up to 650°C. The creep strength of austenitic creep-resistant steels has been enhanced to enable operation up to temperatures of 675–700°C through the development of high Cr, high nickel steels. In the field of nuclear power, creep-resistant steels, which are excellent both in high-temperature creep strength and in irradiation resistance, have been developed for cladding tubes for 650°C fast breeder reactors. The temperature and pressure used were 454°C and 17 MPa, respectively in the early 1990s for hydrogen refining equipment in chemical plants, when reaction chambers were made of 2.25Cr–1Mo steel, but the subsequent development of high-strength 3Cr–1Mo–V steel and 2.25Cr– 1Mo–V steel raised the limiting temperature and pressure to 482°C and 24 MPa, respectively, by 1995. These figures are now about to reach 510°C and 24 MPa. For power generation from wastes, the development of austenitic creep-resistant steels that have high corrosion resistance enabled the boiler steam temperature to be raised from about 300°C in conventional plants up to about 500°C in more modern plants. In the automotive field, exhaust manifolds used to be made of cast iron to withstand exhaust heat. However, as the exhaust gas temperature rose with improved engine performance,
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Preface
higher strength was required and so 18Cr–2Mo–Nb and other steels were developed, raising the exhaust gas temperature to 900°C or higher. Recent research on enhancing the creep strength of 9–12Cr steels for 650°C operation has revealed that the formation of even a partially weak microstructure near a grain boundary promotes local creep deformation and causes premature fracture. This suggests the importance of taking into account microstructural evolution phenomena during creep such as precipitation and coarsening of carbonitrides and intermetallic compounds, dynamic recovery and dynamic recrystallization, in the matrix as well as in the vicinity of grain boundaries. Recently, some high-strength 9–12Cr steels have been found to suffer premature loss of creep strength at 550°C or higher often after prolonged use up to relevant times. Therefore, efforts have been made to clarify the mechanisms of creep strength loss, using modern transmission electron microscopy studies. Extrapolation of short duration laboratory data using time–temperature parameter (TTP) methods, such as the Larson–Miller parameter, have been used widely in the past to predict long-term life. However, it has now become clear that conventional TTP methods tend to overpredict the long term strength because of microstructural degradation phenomena. To address this issue, new analysis techniques have been proposed taking the mechanisms of creep deformation and creep rupture into account. Welded structures made of ferritic creep-resistant steels used under high temperature and low stress (about 600°C and 100 MPa or less) are subject to premature brittle creep fracture by the so-called type IV fracture in the finegrained heat-affected zone (HAZ). Therefore, 9–12Cr steels are being investigated to clarify the mechanisms and the means of preventing this form of fracture. Operation of thick section components under thermally cyclic conditions further exacerbates the cracking problem by creep–fatigue interaction. Thus, as plant temperatures are raised to improve energy efficiency, it is becoming increasingly important to establish the foundation of creepresistant steels that can be used safely for a long time without showing deterioration of creep strength and creep ductility. The aim of this book is to consolidate and review the current state of knowledge of creep resistant steels, summarizing the information which is now scattered throughout voluminous scientific journals and a large number of proceedings of international conferences. Each chapter of the book has been written for engineers and researchers in particular by a world renowned expert in the field. Therefore, the book contains not only background on materials but also recent progress from an engineering and technology point of view. It also can be used as a reference source by graduate level students. It is hoped that the book will serve as an authoritative source of information relating to creep of steels. This book consists of three parts: a general Part I on specifications and
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manufacture, Part II on the behaviour of creep-resistant steels and Part III on specific applications. The introductory Part I includes the introductory description of creep and rupture (Chapter 1) and the historical development of creep-resistant steels (Chapter 2). Part I also includes the specifications of creep-resistant steels in Europe (Chapter 3) and in Japan (Chapter 4) and the production of creep-resistant steels for turbines (Chapter 5). Part II on the behaviour of creep-resistant steels covers physical and elastic behaviour (Chapter 6), diffusion behaviour (Chapter 7), fundamental aspects of creep deformation (Chapter 8), strengthening mechanisms (Chapter 9), precipitation (Chapter 10), grain boundaries (Chapter 11), fracture mechanisms and creep fracture (Chapter 12), mechanisms of creep deformation (Chapter 13), constitutive equations for creep curves and the prediction of service life (Chapter 14), creep strain analysis (Chapter 15), creep crack growth and creep-fatigue behaviour (Chapter 16), creep strength of welded joints (Chapter 17), fracture mechanics (Chapter 18), and oxidation and corrosion (Chapter 19). Part III on specific applications includes the alloy design philosophy behind creep-resistant steels (Chapter 20), creep-resistant steels in turbines (Chapter 21), creep-resistant steels in nuclear reactors (Chapter 22), and industry needs and future research trends in understanding creep damage (Chapter 23). We are grateful to all the contributors for their willing participation and for the cooperation they have extended to us in producing this book. We are also grateful to Mr Robert Sitton, Mr Ian Borthwick, Mrs Lynsey Gathercole and Ms Laura Bunney of Woodhead Publishing for their help in the publication of this book. F. Abe T.-U. Kern R. Viswanathan
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Part I General
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Creep-resistant steels
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1 Introduction F . A B E, National Institute for Materials Science (NIMS), Japan
1.1
Definition of creep
Plastic deformation is irreversible and it consists of time-dependent and time-independent components. In general, creep refers to the time-dependent component of plastic deformation. This means that creep is a slow and continuous plastic deformation of materials over extended periods under load. Although creep can take place at all temperatures above absolute zero Kelvin, traditionally creep has been associated with time-dependent plastic deformation at elevated temperatures, often higher than roughly 0.4Tm, where Tm is the absolute melting temperature, because diffusion can assist creep at elevated temperatures. For detailed description of mechanical equation of state, creep behavior of metals and alloys, dislocation motion during creep, mechanisms of creep, creep damage and fracture, the reader is referred to standard text books on creep.1–6
1.2
Creep and creep rate curves
Creep tests can be conducted either at constant load or at constant stress. For experimental convenience, most frequently the creep tests of engineering steels are conducted at constant tensile load and at constant temperature. The test results can be plotted as creep curves, which represent graphically the time dependence of strain measured over a reference or gauge length. Figure 1.1 shows schematically three types of creep curves under constant tensile load and constant temperature conditions and also their creep rates ε˙ = dε/ dt, where ε is the strain and t the time, as a function of time. Textbooks on creep of metals and alloys generally describe that three stages of creep, consisting of primary or transient, secondary or steady-state and tertiary or acceleration creep that appear after instantaneous strain ε0 upon loading as shown in Fig. 1.1(a), when the test temperature is high enough or at a high homologous temperature. The homologous temperature is defined as the ratio T/Tm, where T is the test temperature in absolute Kelvin and Tm the 3 WPNL2204
Creep-resistant steels εr
(d) (a) Three stages creep curve
ε2 ε1 ε0 Strain
t1
t2
tr
εr (b) Two stages creep curve
εm
log (strain rate or creep rate)
4
Steady-state creep rate
. εs
t1
t2
tr
(e)
Minimum creep rate
. εmin
ε0
tm
tr
tm
tr
(f)
(c) Logarithmic creep curve ε0 Time
Time
1.1 (a), (b) and (c) Creep curves of engineering steels under constant tensile load and constant temperature and (d), (e) and (f) their creep rate curves as a function of time.
absolute melting temperature. The instantaneous strain ε0 contains elastic strain and possibly plastic strain depending on the stress level. In the primary creep stage between ε0 and ε1, the creep rate, ε˙ , decreases with time, as shown in Fig. 1.1(d). The decreasing creep rate in the primary creep stage has been attributed to strain hardening or to a decrease in free or mobile dislocations. In the secondary creep stage between ε1 and ε2, the creep rate remains constant. This creep rate is designated as a steady-state creep rate, ε˙ s , which is given by ε˙ s = (ε2 – ε1)/(t2 – t1) and is commonly attributed to a state of balance between the rate of generation of dislocations contributing to hardening and the rate of recovery contributing to softening. At high homologous temperatures, creep mainly involves diffusion and hence the recovery rate is high enough to balance the strain hardening and results in the appearance of secondary or steady-state creep. In the tertiary creep stage, the creep rate increases with time until rupture at rupture time tr and rupture strain, εr. It should be remembered that under the constant tensile load, the stress continuously increases as creep proceeds or as cross-section decreases and a pronounced effect of increase in stress on the creep rate appears in the tertiary creep stage. Necking of the specimens before rupture causes a significant increase in stress. The increase in creep rate with time in the tertiary creep stage can follow from increasing stress or from microstructure evolution including
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5
damage evolution taking place during creep. Microstructure evolution usually consists of dynamic recovery, dynamic recrystallization, coarsening of precipitates and other phenomena, which cause softening and result in a decrease in resistance to creep. Damage evolution includes the development of creep voids and cracks, often along grain boundaries. The extent and shape of the three creep stages described above can vary markedly depending on test conditions of stress and temperature, as shown schematically in Fig. 1.2, where the final point in each curve represents creep rupture. With increasing stress and temperature, the time to rupture and the extent of secondary creep usually decrease but the total elongation increases. Under certain conditions, the secondary or steady-state creep stage may be absent, so that immediately after the primary creep stage the tertiary creep stage begins at tm, as shown in Fig. 1.1(b) and 1.1(e). In this case, the minimum creep rate, ε˙ min , can be defined instead of the steady-stage creep rate, ε˙ s . Similar to the steady-stage creep rate, ε˙ s , the minimum creep rate, ε˙ min , can be explained by the process where hardening in the primary stage is balanced by softening in the tertiary stage. In many cases, there is substantially no steady-state stage in engineering creep-resistant steels and alloys. Many researchers have shown that there is an ever-evolving microstructure during creep for engineering creep-resistant steels and alloys. This suggests that there is no dynamic microstructural equilibrium in engineering creep-resistant steels and other alloys during creep, which characterizes steady-state creep of simple metals and alloys. Therefore, the term ‘minimum creep rate’ has been favored by engineers and researchers who are concerned with engineering creep-resistant steels and alloys. The stress dependence of minimum or steady-state creep rate is usually expressed by a power law as:
Strain
Arrows: increasing stress and temperature
Time
1.2 Schematic creep curves varying with stress and temperature.
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Creep-resistant steels
ε˙ min or ε˙ s = A σ n
[1.1]
A = A′ exp (– Qc/RT)
[1.2]
where n is the stress exponent, Qc the activation energy for creep, R the gas constant and T the absolute temperature. The parameter A′ includes microstructure parameters such as grain size and so on. Equation [1.1] is often referred to Norton’s law. It is well known that the minimum or steadystate creep rate is inversely proportional to the time to rupture tr as:
ε˙ min or ε˙ s = C /( t r ) m = A ′ σ n exp (– Qc / RT )
[1.3]
where C is a constant depending on total elongation during creep and m is a constant often nearly equal to 1. Equation [1.3] is often referred to as the Monkman–Grant relationship, which has been experimentally confirmed not only for simple metals and alloys but also for a number of engineering creepresistant steels and alloys. Equation [1.3] suggests that the minimum or steady-state creep rate and the time to rupture vary in a similar manner to stress and temperature. At low homologous temperatures, with T/Tm often less than roughly 0.3, where diffusion is not important, only the primary stage appears. Usually only limited strains well below 1% occur that do not lead to final rupture, as shown in Fig. 1.1(c) and 1.1(f). This deformation process is designated as logarithmic creep. Considerable efforts have been made to describe the creep curves, namely, the time dependence of creep strain. There are several model equations available for characterizing the primary, secondary and tertiary creep stage characteristics, ranging in complexity from simple phenomenological to physically based constitutive. Recent progress on the suitability of some of these to specific materials classes and analytical applications is reviewed by Holdsworth et al. [7]. Although Fig. 1.1 shows the idealized creep and creep rate curves, engineering creep-resistant steels sometimes exhibit complicated behavior, especially under low stress and long time conditions, reflecting complex microstructural evolution during creep. Complicated behavior is clearly demonstrated by creep rate curves rather than creep curves. Figure 1.3 shows an example of complicated creep rate curves of 1Cr–0.5Mo steel at 550°C.8 At high stresses above 108 MPa, the creep rate curves are relatively simple and consist of the primary and tertiary stages but there is no substantial steady-state stage, similar to Fig. 1.1(e). The shape of creep rate curve becomes gradually complicated with decreasing stress. At low stresses below 88 MPa, two minima appear in the creep rate curves. This suggests that new strengthening effects such as the precipitation of new phases seem to operate after an extended period, causing a decrease in creep rate again after an growing the previous acceleration creep. The subsequent loss of the existing
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7
10–2 265MPa 10–3
216MPa 108MPa 88MPa 74MPa
Creep/h–1
10–4
61MPa 10
–5
10
–6
137MPa
10–7 10–8 10–1
53MPa
1Cr–0.5Mo steel (JIS STBA22) 550°C 100
101
102 103 Time/h
104
105
106
1.3 Creep rate versus time curves of 1Cr–0.5Mo steel at 550°C (823K).
strengthening effects by microstructural evolution such as the coarsening of new phases causes an increase in creep rate again after reaching a second minimum. Eventually the creep rate versus time curves exhibit oscillated shapes under low stress and long time conditions, reflecting complex microstructural evolution during creep. Similar oscillated shapes have sometimes been observed in other low alloyed steels. In fundamental investigations of creep, creep tests are often conducted at constant stress. The applied stress does not change during the creep test provided that the reduction in cross-sectional area is uniform along the whole gauge length. The stress can be kept constant during creep using proper loading mechanisms. When we need to avoid any influence of oxidation, creep tests are usually conducted in vacuum or in an inert atmosphere. Otherwise the influence of oxidation in reducing the cross-sectional area has to be considered, especially at higher temperatures and longer times for low alloyed steels.
1.3
Creep rupture data
Elevated-temperature components used under creep conditions are designed using allowable stress under creep conditions, which is usually determined on the basis of 100 000 h creep rupture strength at the operating temperature, and sometimes also for 200 000–300 000 h creep rupture strength. The 100 000 h creep rupture strength at a temperature T is defined as the stress at which creep rupture, the last point in Fig. 1.1(a) and (b), occurs at 100 000 h. Generally creep rupture data are represented in graphic form showing the
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Creep-resistant steels
relationship between the stress σ and the time to rupture tr. Figure 1.4 shows an example of creep rupture data for 1Cr–0.5Mo steel from the National Institute for Materials Science (NIMS) Creep Data Sheet.9 This figure contains 309 data points for 11 heats. The material specification defines the chemical composition, heat treatment conditions and so on. In terms of the chemical composition of 1Cr–0.5Mo steel (JIS STBA 22), the Cr concentration is specified as the range 0.80–1.25%, the concentration of Mo in the range 0.45–0.65%, and so on. Practically, the melting of steels causes a difference in the concentration of alloying elements, so that No. 1 ingot contains 1.0%Cr and 0.50%Mo, but No. 2 ingot contains 0.90%Cr and 0.60%Mo and so on, in which the two ingots satisfy the materials specification of 1Cr–0.5Mo steel (JIS STBA 22). Usually, such a small variation in chemical composition causes a difference in creep strength. The 100 000 h creep rupture strength is evaluated to be, for example, 61 MPa at 550°C. The creep rate curves shown in Fig. 1.3 were obtained for one heat of the 1Cr–0.5Mo steel shown in Fig. 1.4. The creep rupture data in Fig. 1.4 exhibit rather complicated curves showing inverse sigmoidal bending at intermediate stress levels of about 130 MPa. It should be noted that two minima appear in the creep rate curves at intermediate stress levels and below, while only one minimum appears at higher stress levels, Fig. 1.3. 500 500°C 550°C 600°C 650°C
400 300
Stress (MPa)
200
100
500°C
80 60 50 40
550°C
30 650°C
n = 309 20 10
102
103 104 Time to rupture (h)
600°C 105
1.4 Creep rupture data for 1Cr–0.5Mo steel at 500–650°C.
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Introduction
9
Recently, long-term creep rupture test data and creep strain data beyond 100 000 h have become available for a number of creep-resistant steels in several materials test institutions in the world, for example, in NIMS, Japan. For long-term creep and creep rupture data, the reader is referred to the NIMS Creep Data Sheets, for example.10 NIMS Creep Data Sheets contain a full set of data, such as creep rupture data, often exceeding 100 000 h, minimum creep rates, short-time tensile data, evaluation of short-time tensile strength and long-term creep rupture strength by curvilinear regression analysis and optical micrographs, together with the details of materials production procedures and chemical compositions. Microstructure Data Sheets, have also been published as the Metallographic Atlas of Long-Term Crept Materials,11 another series of NRIM Creep Data Sheets. The Metallographic Atlas not only contains series micrographs that show microstructural evolution during creep for up to 100 000 h, but that also show related data such as time–temperature–precipitation (TTP) diagrams, histograms describing the distributions of precipitates and creep-voids, and creep damage parameters, using specimens in the Creep Data Sheets. Furthermore, the Atlas of Creep Deformation Property12 was published as Creep Strain Data Sheets for Grade 91 steel (9Cr–1Mo–V–Nb), providing creep curves, creep rate curves and related data. As can be seen from Equation [1.3], stress and temperature are important variables that influence creep rate and time to rupture. In addition, creep and creep rupture properties are markedly affected by not only microstructure variables but also by external variables. The external variables include prestraining (cold-working), additional heat treatments, oxidation and corrosion, stress mode such as uniaxial or multiaxial loading, and superimposition of cyclic loading (creep–fatigue mode). High-temperature structure components in plants are usually used under the complicated conditions described above over long duration up to 300 000 h or longer.
1.4
Deformation mechanism map
Ashby13 proposed the concept of a deformation mechanism map, based on the assumption that all six deformation mechanisms concerned are mutually independent and operate in a parallel way. The six deformation mechanisms include (1) defect-less flow, (2) glide motion of dislocations, (3) dislocation creep, (4) volume diffusion flow, (5) grain boundary diffusion flow and (6) twinning. The twinning can supply only a limited amount of deformation and usually does not appear in the deformation mechanism map. It should be noted that Ashby considered steady-state flow only but no fracture. As illustrated schematically in Fig. 1.5, the deformation mechanism map is constructed with axes of normalized stress σ/G, where G is the shear modulus and T /Tm is the homologous temperature. The map is divided into fields. Within a
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Creep-resistant steels
Defect-free flow
Normalized stress, σ /G
10–1
Ideal strength Dislocation glide
10–3
Dislocation creep 10–2 Coble creep
10–5
Creep rate 10–8 10–10/S Nabarro creep
10–7 0
0.2 0.4 0.6 0.8 Homologous temperature, T / Tm
10–4
10–6 1.0
1.5 Schematic deformation mechanism map with contours of constant creep rate.
field, one mechanism is dominant, that is, it supplies a greater strain rate than any other mechanisms. The upper limit of the boundary is set by a theoretical or ideal strength of roughly G/20 to G/30. At stresses lower than the ideal strength, the deformation takes place by dislocation glide, as in short-time tensile tests. At stresses lower than yield stress, dislocation creep can take place with the aid of diffusion: probably dislocation core diffusion at low homologous temperatures and volume diffusion at high homologous temperatures. Sometimes the dislocation creep field is further divided into two fields: low- and high-temperature dislocation creep fields. At further low stresses, volume diffusion creep (Nabarro–Herring creep) and grain boundary diffusion creep (Coble creep) dominate. The boundaries between adjacent fields in the creep region indicate the conditions under which two mechanisms contribute equally to the overall creep rate. Using an appropriate constitutive equation for creep rates as functions of stress and temperature, we can calculate the creep rates and can draw the boundaries. This also allows us to plot the contours of constant creep rate onto the map, as shown schematically in Fig. 1.5. The locations of the boundaries between adjacent creep fields differ for different materials and also depend on microstructure valuables such as grain size. Experimentally, the deformation mechanism map can be constructed by the measurements of stress and temperature dependence of strain rates or creep rates caused by the individual mechanisms. It should be also noted that the time dependence is not included in the deformation mechanism map. As
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11
already shown in Fig. 1.2, the creep rate of engineering creep-resistant steels varies in complex manner with time because of complicated microstructure evolution during creep exposure at elevated temperatures. Therefore, in the case of engineering creep-resistant steels, the deformation mechanism map can be applied to predict a dominant deformation mechanism at the beginning of creep under specific stress and temperature conditions.
1.5
Fracture mechanism map
Ashby14 also proposed the concept of fracture mechanism map for face centred cubic (fcc) metals and alloys with axes of normalized stress σ/G and homologous temperature T/Tm, which provides us with information about the dominant mechanism resulting in fracture in a shorter time than any other mechanisms. The fracture mechanism map is more important than the deformation mechanism map in practice, because the former relates to damage and fracture processes, which provide us with useful guidelines for assessment of damage evaluation and the remaining life estimation of components in plants. Because the minimum or steady-state creep rate and the time to rupture vary in a similar manner stress and temperature, as suggested by Eqn [1.3], approaches similar to those employed in the construction of a deformation mechanism map can be adopted for the construction of a fracture mechanism map. Figure 1.6 shows schematically the fracture mechanism map for fcc metals, where a cleavage fracture field does not appear. The ideal strength appears 10–1 Dynamic fracture
Normalized stress, σ / G
Ideal strength
Ductile fracture
10–3 Transgranular creep fracture
10–5 Intergranular creep fracture
10–7 0
Rupture 4
10 h
103 h
105 h rupture strength
0.2 0.4 0.6 0.8 Homologous temperature, T / Tm
1.0
1.6 Schematic fracture mechanism map with contours of constant time to rupture for fcc metals.
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Creep-resistant steels
as the upper limiting fracture strength which will overcome interatomic forces in defect-free materials. At stresses lower than the ideal strength, fracture takes place in a ductile, transgranular way, designated ductile fracture, and often designated ductile transgranular fracture. In the creep regime, two fields of transgranular creep fracture and intergranular creep fracture appear at high and low stresses, respectively. At high temperature and relatively high strain rate, dynamic recrystallization can allow materials to deform extensively so that deformation becomes localized in a neck and failure eventually occurs by the specimen necking until the cross-sectional area has gone to zero, usually called the field of rupture. Because grain boundaries become highly mobile under conditions of dynamic recrystallization, the development of creep voids and cavities is suppressed. Figure 1.7 shows schematically the three fracture mechanisms in creep regime: intergranular creep fracture, transgranular creep fracture and rupture.14 Contours of constant time to rupture can be also plotted onto the map, as shown schematically in Fig. 1.6. Although the axes of most of the fracture mechanism maps are stress and temperature, axes of stress and time are also used. Figure 1.8(a) and 1.8(b) show examples of fracture mechanism maps for 1Cr–1Mo–0.25V steel for a turbine rotor, plotted for stress–time to rupture and for stress–temperature coordinate systems, respectively.15 In these figures, the stress–time to rupture plots and stress–temperature plots for constant times to rupture of 100–100 000 h are superimposed. It should be noted that the axes of stress and temperature but not those of normalized stress σ/G and homologous temperature T/Tm are used in these figures because the objective of constructing fracture mechanism maps is primarily for use in assessment of reliability, such as in the design and remaining life prediction of a steam turbine rotor. The intergranular creep fracture field is located in a long time-
Intergranular creep fracture (voids) (wedge cracks)
Growth of voids by power-law creep (transgranular) (intergranular)
Rupture due to dynamic recovery or recrystallization
(a)
(b)
(c)
1.7 Schematic drawing of three fracture mechanisms in a hightemperature creep regime.
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600
Transgranular creep fracture
Stress (MPa)
400
500°C
525°C
625°C
550°C 575°C Intergranular creep fracture (cavitation)
100 675°C
600°C
Rupture (recrystallization) 102
650°C 103 104 Time to rupture (h)
600
105
Tensile
400
streng
th
Transgranular creep rupture 200
10 2
Ductility minimum 100
h
4
10
h
Intergranular creep fracture (cavitation)
10 3 h
e ur pt gth Ru ren st
Stress (MPa)
Ductility minimum
450°C
200
40
13
5
10
h
40 450
500
550 600 Temperature (°C)
Rupture (recrystallization) 650
1.8 Fracture mechanism maps for 1Cr–1Mo–0.25V steel, as functions of time to rupture and of temperature.
to-rupture region at 500–575°C. The rupture field appears at temperatures higher than 600°C. The region of practical importance for 1Cr–1Mo–0.25V steel turbine rotor in power plants is the low stress and long time-to-rupture region at temperatures of 550°C or lower, which belongs to the intergranular creep fracture field. This suggests that precise measurements of the development of creep voids at grain boundaries during creep contributes to the improvement in the reliability of the remaining life estimation.
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1.6
Creep-resistant steels
References
1. Finnie I. and Heller W. R., Creep of Engineering Materials, McGraw-Hill, New York, 1959. 2. Garofalo F., Fundamentals of Creep and Creep-Rupture in Metals, The Macmillan Company, New York, 1965. 3. Penny R. K. and Marriott D. L., Design for Creep, McGraw-Hill, London, 1971. 4. Evans R. W. and Wilshire B., Creep of Metals and Alloys, The Institute of Metals, London, 1985. 5. Cadek J., Creep in Metallic Materials, Elsevier, Amsterdam, 1988. 6. Viswanathan R., Damage Mechanisms and Life Assessment of High-Temperature Components, ASM International, Ohio, 1989. 7. Holdsworth S. R., Baker A., Gariboldi E., Holmstrom S., Klenk A., Merckling G., Sandstrom R., Schwienheer M. and Spigarelli S., ‘Factors influencing creep model equation selection’, Proceedings of ECCC Creep Conference, 12–14 September 2005, The Institute of Materials, London, UK, 2005, 380–393. 8. Kushima H., Kimura K., Abe F., Yagi K., Irie H., Maruyama K., ‘Effect of microstructural change on creep deformation behaviour and long-term creep strength of 1Cr–0.5Mo Steel’, Tetsu-to-Hagane, 2000, 86, 131–137. 9. NIMS (formerly NRIM) Creep Data Sheets No.1. Tokyo, Tsukuba, National Institute for Materials Science, 1996. 10. Series of NIMS (formerly NRIM) Creep Data Sheets No. 1–48. Tokyo, Tsukuba, National Institute for Materials Science, 2007. 11. Series of NIMS Metallographic Atlas of Long-Term Crept Materials No. M1-M6. Tokyo, Tsukuba, National Institute for Materials Science, 2007. 12. NIMS Atlas of Creep Deformation Property, No. D-1. Tokyo, Tsukuba, National Institute for Materials Science, 2007. 13. Ashby M. F., ‘A first report on deformation-mechanism maps’, Acta Metallurgica, 1972, 20, 887–897. 14. Ashby M. F., Gandhi C. and Taplin D. M. R., ‘Fracture-mechanism maps and their Construction for FCC Metals and Alloys’, Acta Metallurgica, 1979, 27, 699–729. 15. Shinya N., Kyono J. and Kushima H., ‘Creep fracture mechanism map and creep damage of Cr–Mo-V turbine rotor steel’, ISIJ International, 2006, 46, 1516–1522.
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15
2 The development of creep-resistant steels K.-H . M A Y E R , ALSTOM Energie GmbH, Germany and F . M A S U YA M A , Kyushu Institute of Technology, Japan
2.1
Introduction
The development of creep-resistant steels is a result of continuous technological progress throughout the 20th century. The urgent need to improve the creep strength of steels was based on endeavours by the power station industry to improve the thermal efficiency of steam power plant by raising the steam temperature and steam pressure in order to reduce the cost of fuel and reduce use of fuel resources. Since roughly 1900, as shown for instance by Fig. 2.1, the heat rate of thermal power plant in Germany has been reduced following a step-by-step increase in the steam parameters from 275°C/12 bar to 620°C/ 300 bar.1,2 A major contribution to the increase in power plant efficiency consisted of the development of heat-resistant steels with a higher creep strength at an acceptable creep ductility level (see for example Kallen).3 The significance of these material properties was not recognised until early damage was suffered by steam turbine bolts in the 1930s, which pointed to the fact that the strength of steels used in power stations operating at higher temperatures depends significantly on the creep behaviour of the material over the full period of operation.4 Based on this experience it was concluded that the strength values should no longer be determined in short-term tests,5,6 for example the ‘durability strength’ according to the DVM (Deutscher Verband fur Material prüfung) creep rate limit test. The procedure to be adopted should be to determine the fracture strength, the creep elongation and creep ductility of the heat-resistant steel in a creep test extending over a period of roughly 100 000 h (see for example Siebel).7 For the DVM creep rate limit test established in Germany in 1930, the ‘durability strength’ was defined to be the stress at the test temperature at which a creep rate of 10 × 10–4 %/h was reached between the 25th and 35th hour.5 Typical results of both tests, which were commenced at the end of the 1930s, are shown by Fig. 2.2. The tests were performed at 500°C on a steel containing 0.30%C–1.61%Cr– 1.28%Mo–0.10%V. The DVM creep rate limit test using smooth specimens 15 WPNL2204
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Creep-resistant steels
40
Specific heat rate (kJ/kWh)
12 bar/275°C 30
15 bar/350°C
20
35 bar/450°C 100 bar/500°C 100 bar/540°C
10
With reheat (540/540°C) 0 1900
1910
1920
1930
1940
1950 1960 Year
1970
280 bar 580/600°C 300 bar 580/600°C
Supercritical (250 bar) 1980
1990
2000
2010
2.1 Heat rate of steam power plants in Germany as a function of steam parameters since the year 1900.
Creep rupture strength (MPa)
1000 500°C
600
Smooth specimens 400 313192 h 200 ‘Durability strength’ of DVM creep rate limit test carried out at about 1936
Notched specimens 294000 h
100 10–1
100
101
102 103 Time to rupture (h)
104
105
106
2.2 Creep rupture strength as a function of time to rupture and ‘durability strength’ of a 1.6%CrMoV steel at 500°C. Test steel: 0.30%C–1.6%Cr–1.3%Mo–0.1%V; heat treatment: 950°C/air + 680°C/air; tensile strength: 893 MPa.
and the creep rupture test using smooth and notched specimens (notch factor 4.3) are compared. Creep rupture tests were even continued up to roughly 300 000 h at the end of the 1970s.8 The ‘durability strength’ in the short-term test was determined at a strength level of 306 MPa. At this stress level, the rupture of the creep rupture tests was reached after about 3000 h, whereas the 100 000 h rupture strength of the smooth specimens lies at 190 MPa. The notched specimens stressed at the same level of 190 MPa failed after roughly 30 000 h, distinctly earlier than the smooth specimens owing to a significant notch-weakening
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The development of creep-resistant steels
17
behaviour. It was also recognised that the tendency to notch-weakening behaviour of the steels was a wrong turn in the development of heat resistant steels on the basis of the DVM creep rate limit test, because the aim of raising the ‘durability strength’ of the heat-resistant steels as high as possible involves the risk of increasing the susceptibility of the steels to embrittlement. To provide a further example of the influence of creep processes on the strength of heat resistant steels, Fig. 2.3 demonstrates the dependence of the creep rupture strength and the strength for 1% creep strain on the test temperature and test period for a carbon steel and a 1%Cr–0.5%Mo steel in comparison with the 0.2% yield limit determined in the short-term tensile test (see for example Wellinger).9 In comparison with the 0.2% yield limit determined in the short-term tensile test, the 100 000 h creep rupture strength is lower for the carbon steel at higher than about 410°C and is also lower for the 1%Cr–0.5%Mo steel higher than about 480°C. The crossover temperatures between the results of the short-term tensile test and the creep strength values are distinctly lower if the 0.2% or the 1% permanent creep strain determined in the 100 000 h test are decisive for the design of power station components. Forming influences marking the development of heat resistant steels over the past 100 years are: long-term operational experience experience gained from long-term creep rupture tests improvements in melting technology systematic investigations into the influence of heat treatment on creep behaviour 300
0.2-Limit
Strength (MPa)
• • • •
1% Cr 0.5% Mo steel
200 100 000 h creep rupture 100
100 000 h 1%-creep strain
C steel
200
300 400 500 Test temperature (°C)
600
700
2.3 0.2-limit, 100 000 h creep rupture strength and 100 000 h 1%creep strength of a carbon steel and 1%Cr–0.5%Mo steel as a function of test temperature.
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• • • • •
• • • •
2.2
Creep-resistant steels
examination of the microstructure of specimens in the virgin condition following long-term thermal and creep loading systematic investigations into the influence of alloying elements computer-aided alloy design methods (e.g. Thermocalc, DICTRA) modelling of creep processes development of metallographic methods and equipment for the identification of precipitates (e.g. transmission electron microscopy (TEM), energy dispersive X-ray spectrometry (EDS), energy filtered transmission electron microscopy (EFTEM), atom probe field ion microscopy (APFIM), field emission Auger electron spectroscopy (FE-AES) and secondary ion mass spectroscopy (SIMS). national and international joint research activities and research projects related to the development of advanced creep resistant steels and longterm tests under creep stress conditions7–16 testing of newly developed heat-resistant steels on the basis of large pilot components and welds fabricated under normal workshop conditions investigations into the oxidation behaviour of advanced heat-resistant steels in the laboratory and in test fields of steam power stations international exchange of experience at conferences and in workshops e.g. EPRI (Electric Power Research Institute USA), EPDC (Electric Power Development Center/Japan), COST (Community of Science and Technology of the European Communities), ECCC (European Creep Collaboration Committee), NIMS (National Institute for Materials Science/ Japan).
Requirements for heat-resistant steels
Heat-resistant steels for use in thermal power stations must be capable of satisfying the specific requirements established for dependable and economic operation. All phases of development and testing must therefore be specifically aligned to the following requirements: • • • • • • • •
high thermal efficiency operational capability in the medium and peak load ranges life expectancy of at least 200 000 h high availability long intervals between overhauls short overhaul periods short manufacturing times competitive production costs for the steam plant and electric power.
These requirements mean that the application of newly developed steels must not involve any additional risks, implying:
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• • • •
19
that long time creep testing up to 100 000 h is needed to predict reliably the creep strength for 200 000 h (which means that the tests must be started with a large number of specimens, because at the outset of the tests a prediction cannot be made about the stress level which will be reached after a test period of 100 000 h. A long test period should also be scheduled if a research project is only due to last for 3–5 years); satisfactory oxidation resistance; high ductility of the steels under conditions of creep stressing; high fracture toughness of the steels in a new condition and following prolonged operational stressing; satisfactory production of the new steels in terms of melting, casting, forging, hot forming and welding.
2.3
Historical development of ferritic steels
2.3.1
Carbon steels
Up to the 1920s it was general practice to use non-alloyed steels for components in the steam admission zone exposed to maximum temperatures of 350°C and pressures of about 15 bar. The components were designed according to the material requirements established in a hot tensile test. In these short-term tests it was not possible to recognise that the elements N, Al and Mn exercised a major influence on the creep strength of carbon steels. Figure 2.3 has already shown the 0.2-limit and the creep rupture strength obtainable with present-day standard non-alloyed steels as a function of the test temperature in comparison with the 1%Cr–1%Mo steel.9
2.3.2
Low alloy steels
At the beginning of the 1920s, operation at steam temperatures of 450°C and pressures of 35 bar called for the development of low-alloyed heat-resistant steels. Developments were limited to individual steel works which at that time were not yet coordinated in joint research programmes. The steels were identified by the trade name of the steel works. The basic test in the development of low-alloyed steels was a hot tensile test which later on was followed by a short-term test, for example the DVM creep rate limit test in Germany.5 In the USA in 1933, a guideline was prepared between the ASME and ASTM for tests covering periods of 500–2000 h to determine the creep strain limits for a permanent creep strain of 0.01%, 0.1%, 1% and the ultimate rupture limit.6 The results were extrapolated at a double logarithmic scale on a straight line pattern up to 104 h and further up to 105 h. Based on the multiplicity of investigations of test steels carried out with different Mo, Cr, Ni, V, CrMo, CrV, MnSi, MoMnSi, CrSiMo, CrNiMo,
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Creep-resistant steels
CrMnV, CrMoV contents, worldwide developments in the manufacture of steam boilers and small forgings for steam turbines produced steels with chemical compositions of 0.15%C–0.3–0.5%Mo, 0.13%C–1%Cr–0.5%Mo17 and 0.10%C–2.25%Cr–1%Mo20 which are still in use today. In addition, in the 1950s, a MoV steel with a composition of 0.14%C–0.5%Mo–0.3%V with an even higher creep strength was developed in Europe for gas turbines and later also qualified in long-term creep tests for steam plants. In the field of turbine manufacturing since the 1950s a steel with a composition of approximately 0.25%C–1.25Cr–1%Mo–0.30%V is in use worldwide for turbine rotors, casings, bolts and small forgings. Systematic investigations into the creep strength of the steels developed in short-term tests between the 1920s and 1940s were followed in the 1950s by long-term creep tests.10–16 In Germany, for instance, a joint research project was established for this purpose in 1949 between steel and power plant manufacturers and plant operators.10 Long-term creep tests of individual melts have actually been performed by research bodies in Germany since the mid-1930s (see for example Diehl and Granacher).8 The activities of the individual national creep groups operating within Europe were coordinated in 1990 and culminated in the establishment of the European Creep Collaborative Committee (ECCC) in December 1991.11 Molybdenum was recognised as an important element for increasing high temperature strength. Mo steels developed in the USA and the UK are alloyed with a Mo-content of about 0.5%. The Mo-content of the steel developed in Germany is roughly 0.3% at a C-content of about 0.15%. Figure 2.4 illustrates the influence of molybdenum on the 100 000 h creep rupture strength at 450°C as opposed to an unalloyed steel with roughly 0.15%C.18 By the addition of approximately 0.5%Mo, the 100 000 h creep strength of the unalloyed steel of roughly 70 MPa is increased to about 260 MPa. The alloying effect of Mo is the result of solution hardening and Mo2C precipitation.9,18 A drawback of Mo alloying to over about 0.35% is a marked decline in ductility under creep stress conditions as well as graphite precipitation. Consequently, steels with an Mo-content of 0.5% should not be used in temperature environments over 400°C. However, the strengthincreasing effect of higher Mo contents, without an unacceptable decrease in ductility, can be utilised by the addition of Cr as in the case of the steels with a composition of 0.13%C–1%Cr–0.5%Mo and 0.10%C–2.25%Cr–1%Mo. Figure 2.5 shows the influence of Mo and Cr on the 100 000 h creep rupture strength of the three steels 0.3%Mo, 1%Cr–0.5%Mo and 2.25%Cr– 1%Mo at 500°C and 550°C.18 The highest creep rupture strength is already achieved by the 0.13%C–1%Cr–0.5%Mo steel at 500°C. At 550°C, subjected to an increase in the Mo and Cr-contents as in the case of the 0.10%C– 2.25%–Cr1%Mo steel, a further increase in the creep rupture strength is obtained. Microstructure investigations on the initial condition of the 0.13%C–
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The development of creep-resistant steels
21
300
200 450°C
100
C-steel
0.1
0.2
0.3 0.4 Mo content (mass%)
0.5
0.6
2.4 100 000 h creep rupture strength of a C-steel as a function of Mo content at 450°C.
1%Cr–0.5%Mo steel revealed M3C, M7C3 and M23C6 precipitations whereas for the 0.10%C–2.25%Cr–1%Mo steel Mo2C and M23C6 precipitations were found (see for example Florin).19 An excellent literature survey about the microstructure formation of the CrMo steels after heat treatment and long term creep is given by Orr et al.20 The 0.14%C–0.6%Mo–0.3%V steel, which in view of its higher creep rupture strength (Fig. 2.8) is given preference for live steam pipes and pipes with superheated steam, features higher strength than the 0.10%C–2.25%Cr– 1%Mo steel owing to finely distributed and thermally very stable V4C3 precipitation and Mo2C.19,21 A drawback for this steel is its tendency to type IV cracking in the intercritical area of the heat affected zone of welds (see for example Schüller et al).22 Amongst the numerous steel versions developed in the 1930s and 1940s for the manufacture of rotors, casings, valves and bolts for steam turbines, a 1%CrMoV steel has found worldwide acceptance which, depending on component size and the location of the site of development, is alloyed with a composition of roughly 0.20–0.30% C, 1–1.5% Cr, 0.70–1.25% Mo, 0.25– 0.35% V and 0.50–0.75% Ni.15,23,24 Figure 2.6 shows schematically the relationship of the 100 000 h creep rupture strength and the fracture toughness FATT50 as a function of the microstructure for the 1%CrMoV steel.26 The highest creep rupture strength of this steel type is achieved with an upper bainite structure.25 The disadvantage of the upper bainite structure is the lower toughness25 so that
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22
200 Steel Steel 1%Cr–0.5%Mo 2.25%Cr–1%Mo 500°C
160
120
80
Steel 0.3Mo
40
0.2
100 000h creep rupture strength (MPa)
500°C
0.4 0.6 0.8 Mo content (mass%)
1.0
200
160
Steel Steel 1%Cr–0.5%Mo 2.25%Cr–1%Mo 500°C
120
80
500°C Steel 0.3%Mo
40
0.5 1.0 1.5 Mo content (mass%)
2.0
2.5 100 000 h creep rupture strength of low alloyed steels as a function of Mo and Cr content at 500°C and 550°C.
the individual alloying elements of the steel as well as heat treatment must be aligned to the specific operational properties of the components.26 In some cases the procedure for turbine rotors is to adapt the heat treatment contour to the operational properties and/or to apply a method of spray hardening with different quenching rates for the specific regions of the components.26 Investigations into the microstructure in the initial state revealed V4C3, Mo2C and M23C6 (see for example Smith).27 With regard to the ductility and toughness of the 1%CrMoV steel, experience of operational stressed components has emphasised the significance of the austenitising and tempering temperature.28–32 Figure 2.7 illustrates the behaviour of smooth and notched specimens at
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1200
Temperature (°C)
TTT-diagram 1%CrMoV steel
800 F 400
Perlite
B M
102 Martensite (M)
104 Time (s)
106
FATT
FATT
Creep strength
Long term creep strength
Upper Ferrite (F) Lower Bainite (B)
2.6 Creep rupture strength and toughness FATT of 1%CrMoV steel as a function of cooling rate after austenisation respectively for martensite, bainite and ferrite microstructures (schematic). TTT is the time temperature transformation.
500°C in a creep rupture test for two different heat treatment temperatures of two 1%CrMoV melts over test periods up to about 120 000 h.28 Austenitising at 1050°C in connection with a tempering temperature of 700°C was found to cause substantial notch weakening and very low ductility of the smooth specimens (case 17c). An acceptable deformation behaviour is obtained with an austenitising temperature of 980°C and tempering treatment at 670°C. A further disadvantage of heat treatment at an excessive austenitising temperature is the initiation of long-term embrittlement which results in a remarkable reduction of toughness at low temperatures. This loss of toughness has been
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24
17a: 0.19%C–1.32%Cr–1.05%Mo–0.54%V 980°C/oil + 2h 670°C/air
17c: 0.17%C–1.10%Cr–1.16%Mo–0.35%V 1050°C/oil + 3h 700°C/air
300 Smooth specimen
Notched specimen
200 500°C
100
Reduction of area (%)
Creep-resistant steels
Creep rupture strength (MPa)
Smooth specimen Notched specimen
400
500°C
102
103
104
105
102
103
104
105
102
103 104 Time to fracture (h)
105
102
103 104 Time to fracture (h)
105
100 80 60 40 20
2.7 Creep rupture strength of smooth and notched specimens of 1%CrMoV steels as a function of heat treatment and time to fracture at 500°C.
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25
105 h creep rupture strength (MPa)
the cause of brittle failure of turbine and valve bolts in the past.32,33 The long-term embrittlement also increases the risk of brittle failure of HP (high pressure) and IP (intermediate pressure) turbine rotors.34,35 In one investigated case, a FATT50 of 340°C was found after long-term service of a turbine rotor with a component temperature of about 380°C.36 However, in this connection, attention must be drawn to the fact that in accordance with technological progress in the early 1950s, the trace elements, owing to the melting process, were still at a relatively high level (e.g. phosphorus up to 0.028%). These high contents of trace elements also contributed to the embrittlement.36 Good long-term experience was gained in Germany with components of 0.20%C– 1%Cr–1%Mo–0.3%V steels in the mid-1950s when the austenitising temperature was limited to a maximum of 950°C and tempering treatment was fixed at 680–740°C. The maximum permissible tensile strength was specified at 835 MPa. Figure 2.8 provides an overview of the 100 000 h creep rupture strength as a function of the test temperature for the non-alloyed and low-alloyed heat-resistant steels currently established for the temperature range below about 565°C. Two new low-alloyed heat-resistant steels have been developed over the past 15–20 years predominantly for the manufacture of water walls for advanced steam power stations. The steels are named HCM2S (0.06%C– 2.25%Cr–2%Mo–1.6%W–0.25%V–0.05%Nb–0.02%N–0.003%B)37 and 7CrMoVTiB (0.07%C–2.4%Cr–1.0%Mo–0.25%V–0.07%Ti–0.01%N– 0.004%B).38 Both steels lend themselves well to welding and do not require post-weld heat treatment. Their creep rupture strengths in comparison with Steel (a) (b) (c) (d) (e) (f)
C 0.18 0.15 0.13 0.10 0.14 0.28
Cr – – 1.0 2.25 0.5 1.0
Mo – 0.3 0.5 1.0 0.6 1.0
V (mass%) – – – – 0.3 0.3
*Upper bainite
200
(f) 1%CrMoV* (e) 0.6%MoV 100
(b) 0.3%Mo (d) 2.25%CrMo
(a) C-steel
(c) 1%CrMo 0 450
500
550
600
Temperature (°C)
2.8 100 000 h creep rupture strength of a C-steel and low alloyed heat-resistant steels as a function of test temperature.
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Creep-resistant steels
the conventional used 0.15%C–0.5%Mo and 0.13%C–1%Cr–0.5%Mo steels are given in Fig. 2.9.
2.3.3
9–12%Cr steels
The development of heat-resistant 9–12%Cr steels was strongly motivated by two major events. During the 1950s it was the development of thermal power stations for public power supply, operating at steam temperatures ranging from 538°C to 566°C and during the 1980s the target was set to develop low-pollution power stations operating at steam admission temperatures of 600–650°C and supercritical pressures up to 350 bar. Figure 2.10 presents a summary of current national and international research projects in progress since the 1980s in Japan, USA and Europe. An overview of the historical development of heat-resistant ferritic– martensitic 9–12%Cr steels from the 1950s to the 1990s is given in the upper part of Fig. 2.11. The lower part shows recent values of the 100 000 h creep rupture strength at 600°C, extrapolated from long-term test data. Table 2.1 illustrates the chemical composition of the steels. As a rule, the steels are an onward development of steels already applied over extended periods of time by using the trial-and-error method. Development of 9–12% Cr steels for steam temperatures up to 620°C
105 h creep rupture strength (MPa)
The steel X22CrMoV 12 1 was developed in the 1950s for thin-walled and thick-walled power station components. Its creep strength is based on solution Steel (a) (b) (c) (d)
C 0.16 0.13 0.06 0.07
Cr – 1.0 2.25 2.4
Mo 0.3 0.5 0.2 1.0
W – – 1.6 –
V – – 0.25 0.25
Nb – – 0.05 –
Ti – – – 0.07
N – – 0.02 0.01
B (mass%) – – 0.003 0.004
200 (d) 7 CrMoVTiB 10 10 (c) HCM 2 S 100 (b) 1%CrMo
(a) 0.3%Mo
0 450
500
550
600
Temperature (°C)
2.9 100 000 h creep rupture strength of heat-resistant low alloyed steels used for water walls of steam plant boilers as a function of test temperature.
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27
International projects of advanced power plants
Japan R & D : DPDC Manufacturers, utilities, EPDC 1981–1991 316 Bar 566/566/566°C 314 Bar 593/593/593°C 343 Bar 649/593/593°C 50 MW pilot power plant
1989–1990 1991–1993 1994–2000 300 bar 630/630°C
USA
Europe
R & D : EPRI
Cost 50/501
Manufacturers Manufacturers. steelworks, Study 1978–1980 utilities and R & D institutes 310 Bar 566/566/566°C 1983–1997 310 Bar 593/593/593°C 300 bar 600/600/600°C 345 Bar 649/549/549°C 300 bar 600/620°C Steels and components for Boiler + Turbine EPRI-RP 1403-15 COST 522 300–900 MW R&D: 1986–1993 Steels and components for boiler + turbine (USA, Japan, Alstom+Man)
NRIM-STX 21Project: USC 650°C/350 bar boiler
EPRI-RP 1403-50 -WO9000-38
1997–2012
1990–1999
Thick wall boiler components
1998–2003 300 bar 620/650°C Steels and components for boiler + turbine
COST 536 2004–2008 300 bar max. 650°C Steels and components for boiler + turbine
Thick wall pipes: P 92 + P 122 (USA, Japan, UK + Denmark)
2.10 International research projects for the development of heatresistant steels for advanced steam power plants since 1978.
hardening and on the precipitation of M23C6 carbides. The steel has been applied successfully in power stations over several decades. The steels H46, FV448 and 56T5 (nos. 2 and 3 in Fig. 2.11) exhibit additional alloying of 0.30–0.45% Nb and roughly 0.05 N. The targeted increase in strength is obtained by secondary MX precipitations of the type VN and Nb (C,N). However, a distinct improvement in creep strength at 600°C, which is of primary interest for components for the aviation industry, is only obtainable in the short-term range. In view of the high Nb-content, these steel grades are only suitable for the manufacture of small-size components because the relatively high Nb-content results in pronounced segregations in ingots used in manufacturing thick-walled components. TAF steel (no. 4) developed in Japan by Fujita39 for small components is an onward development of European Nb-containing steels (no. 2: H46 and FV 448). In addition to an improved balance of the alloying elements–based on a very extensive investigation of the influence of all alloying elements on the creep strength – it also features a high boron contents up to 0.040%, which permits the steel only to be used for small components. According to Fujita’s investigations, boron stabilises the M23C6 carbides by forming M23 (C,B)6. At the end of 1999, Fujita40 gave a report on the actual results of creep tests on specimens of this steel which were carried out at 550°C up to about 70 000 h, at 600°C up to about 20 000 h and at 650°C up to about 125 000 h (Fig. 2.12). The results show that this steel has an extremely high
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Historical development 1 X 22 CrMo(W)V 12 1/rotors, casings, bolts, blades, pipes 2 H 46; FV 448/bolts, blades, gas turbine discs
France:
3 56 T 5/bolts, blades, Japan : 4 TAF/blades, small forgings
(1) Casing EPRI 1403–15 (2) Rotors and casings COST 501–2 Development for fast breeder
USA : 5 11% CrMoVNbN/rotors (GE) USA : 6 X 10 CrMoVNbN 9 1 (P 91)/pipes, pressure vessels, casings Japan : 7 + 8 HCM 12/Tubes; TMK1, TMK2/rotors Improved power plants (600 °C)
Cost 501 : 9 X 18 CrMoVNbB 9 1/rotors Cost 501 : 10 X 12 CrMoWVNbB 10 11/E911 Japan : 11 + 12 NF 616/HCM 12 A/pipes
1950
1960
1970
1980
100 000 h creep strength at 600°C
9
4
120
40
1950
MPa 160
TMK 1 + TMK 2
MPa
80
1990 Year
6 2,3
E911 5
11 12
8 7
10
1
1960
120 80 40
1970
1980
1990 Year
2.11 Overview of the historical development of heat-resistant 9–12% Cr steels within the time range 1950–1995 and the 100 000 h creep rupture strength of these steels at 600°C.
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USA, Germ UK:
Table 2.1 Chemical composition and creep rupture strength at 600°C of the steels in Figure 2.11 (mass%) Country Steel
Chemical composition (weight%) C
X22CrMoV 12 1 H46 FV448 56T5 TAF 11%CrMoVNbN
1. 2.
France Japan USA
3. 4. 5.
USA Japan Japan
6. 7. 8.
Europe Europe
9. 10.
Japan Japan
11. 12.
Advanced steels P 91 HCM 12 TMK 1 TMK 2 X18CrMoVNbB 91 X12CrMoWVNbN E911 P92 P122
Japan Germany
13. 14.
HCM 2S 7CrMoTiB
Cr
Mo
0.22 0.16 0.13 0.19 0.18 0.18
12.0 11.5 10.5 11.0 10.5 10.5
1.0 0.65 0.75 0.80 1.5 1.0
0.50 0.70 0.70 0.40 0.05 0.70
0.10 0.10 0.14 0.14 0.18 0.12 0.11 0.07 0.10
9.0 12.0 10.3 10.5 9.5 10.3 9.0 9.0 11.0
<0.40
0.06 0.07
1.0 1.0 1.5 0.5 1.5 1.0 0.95 0.50 0.40 1.Cu 0.20 1.0
2.25 2.40
Ni
0.60 0.50 0.05 0.80 0.20 0.06 <0.40 – –
W
V
Nb
N
B
104 h
105 h
– – – – – –
0.30 0.30 0.15 0.20 0.20 0.20
– 0.30 0.45 0.45 0.15 0.08
– 0.05 0.05 0.05 0.01 0.06
– – – – 0.035 –
103 118 139 144 216 165
59 62 64 64 (150) (85)
– 1.0 – 1.8 – 0.80 1.0 1.8 2.0
0.22 0.25 0.17 0.17 0.25 0.18 0.20 0.20 0.22
0.08 0.05 0.05 0.05 0.05 0.05 0.08 0.05 0.06
0.05 0.03 0.04 0.04 0.01 0.06 0.06 0.06 0.06
124
94 75 90 90 122 90 98 113 101
1.6 –
0.25 0.25
0.05 –
0.02 0.01
0.01 – – 0.003 0.003 0.003 0.004 0.07Ti
170 185 170 165 139 153 156
80 60
The development of creep-resistant steels
Basic steels Germany UK
Rupture strength at 600°C (MPa)
29
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Creep-resistant steels 500 550°C
400
Creep rupture strength (MPA)
300 600°C 200 650°C
100 80 700°C 60
40 30 100
1000
10 000 Time to fracture (h)
100 000
2.12 Creep rupture strength of the TAF steel (0.18%C–10.5%Cr– 1.5%Mo–0.2%V–0.15%Nb–0.035%B) as a function of temperature and time to fracture.
creep strength. Furthermore they demonstrate the creep strength potentials of ferritic–martensitic 9–11% Cr steels subject to optimum alloying. The rotor steel 11%CrMoVNbN (no. 5), patented in 1964 by the General Electric Company, USA, is also an onward development of the Nb alloyed steels no. 2.41 In particular, the Nb content was greatly reduced (0.08%) in order to prevent harmful segregation at the centre of a rotor. Furthermore, the alloying elements were balanced in order to avoid the formation of delta ferrite. The published creep strength of about 85–90 MPa for 600°C and 100 000 h was extrapolated on the basis of tests at 620°C up to times of 16 195 h duration.42 The steel referred to in the literature as mod. 9Cr1Mo or P91 (no. 6) is a steel of the newer generation. It was developed under a huge American national project in the late 1970s for manufacturing pipes and vessels for a fast breeder reactor. It is tough, readily weldable and, as shown by creep tests at 593°C up to about 80 000 h, has a high creep strength at 600°C and 100 000 h of about 94 MPa.43 In comparison with earlier steels it is characterised, for example, by a lower C content of only about 0.10% and a reduced Cr content of about 9%. This steel has meanwhile found wide application in all new Japanese and European power stations for the manufacture of pipes and small forgings. It is also used for the manufacture of valve chests and turbine casings.44,45 Steel HCM 12 (no. 7) is a newly developed Japanese 12% Cr steel with
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31
0.10%C–1%Mo–1%W–0.25%V–0.05%Nb–0.03%N and a duplex structure of delta-ferrite and tempered martensite with improved weldability and creep strength.37,53 The stability of the creep strength of this steel has been obtained primarily by precipitation strengthening with very fine VN precipitates and high-temperature tempering at over 800°C.37 Experience with this steel has already been accumulated over a period of more than 20 years. The steel has been extensively used for superheater tubes in chemical recovery boilers exposed to severe high-temperature corrosion attack. The good resistance to corrosion is due to the high Cr-content of 12%. The Japanese rotor steels TMK 1 and TMK 2 (no. 8), developed in the 1980s, were based on the known properties of steels no. 1–6.46 Compared with the GE rotor steel (no. 5), the C content in particular was reduced and the sum total (C + N) was selected at around 0.17%. Based on the research work of Fujita,47 the Mo content was raised to 1.5% in TMK 1, whereas TMK 2 is additionally alloyed with about 1.8%W. The Mo content is simultaneously balanced to 0.50% in keeping with Fujita’s result that the highest solution hardening was obtained with (Mo + 0.5W) = 1.5%. Both rotor steels have been used for new power stations in Japan. Rotor steels nos. 9 and 10 are primarily the result of research work performed in the 1980s under the European cooperation programme COST 501.48 Steel no. 9 is a 9%CrMoVNb steel which is additionally alloyed with about 0.01% boron. It is an onward development of the TAF steel for large components with reduced contents of Cr, Nb and B and with an increase in V. Creep tests with specimens from a 900 mm diameter pilot rotor, which have so far reached about 100 000 h, indicate a probable creep rupture strength of about 120 MPa at 100 000 h and 600°C. Creep tests up to 94 000 h with specimens from a 600 mm diameter pilot rotor with a similar composition confirm this value. A larger pilot rotor with a diameter of about 1200 mm is under investigation. The creep strength of the steel X12CrMoWVNbN 10 1 1 (steel no. 10) containing 0.06%N instead of 0.01%B and only 1% Mo instead of 1.5% Mo and, in addition, 0.8% W, is about 20% lower, see Fig. 2.13. So far, the longest test period is about 100 000 h. This rotor steel has primarily been used for the new advanced European steam power plants. This steel type has also successfully been applied for many valve chests and turbine casings in the new plants. The pipe steel, E 911, developed under COST 501, features a somewhat comparable chemical composition. The Cr content was reduced to roughly 9%.49 The Ni content is also distinctly lower since because of the lower Cr content there is no risk of the occurrence of delta-ferrite. Based on results up to roughly 100 000 h, the creep rupture strength of this steel is estimated to be 98 MPa at 600°C and 100 000 h. Steel no. 11 is also based on the research work of Fujita. It is a 9% Cr pipe steel specifically alloyed with 0.10%C–1.8%W–0.5%Mo–0.2%V–0.06%Nb–
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Creep-resistant steels Steel (a) (b) (c) (d) (e)
C 0.28 0.21 0.12 0.12 0.18
Cr 1.0 12.0 10.0 10.0 9.0
Mo 0.9 1.0 1.5 1.0 1.5
W – – – 1.0 –
V 0.30 0.30 0.20 0.20 0.25
Nb – – 0.05 0.05 0.05
N – – 0.05 0.05 0.02
B (weight%) – – – – 0.010
200 (c/d) X12CrMo(W)VNbN101(1)
(a) 1%CrMoV 100
(b) 12%CrMoV
(e) X18CrMoVNbB91 ca. 20°C ca. 70°C
0 500
550
600
650
Temperature (°C)
2.13 100 000 h creep rupture strength of steam turbine rotor steels as a function of test temperature.
0.05%N and about 0.003%B and tempered at about 750°C. It was developed in the second half of the 1980s under the designation NF 616 (P92).50 Based on creep tests up to about 100 000 h, the creep rupture strength is estimated to be 113 MPa at 600°C and 100 000 h.51 A similar pipe steel HCM 12A (P122), steel no. 12, has also been developed in Japan alloyed with a higher Cr content of 11% in order to improve oxidation resistance. About 1% copper has been added to reduce the tendency to deltaferrite formation caused by the higher Cr-content.52 The newest evaluation of the available results of long term creep tests led to an estimated creep rupture strength of 101MPa at 600°C and 100 000 h.53 Figure 2.14 compares the 100 000 h creep rupture strength versus the temperature for the new pipe steels. In addition to the pipe steel P91, all three pipe steels (E 911, NF 616 and HCM 12A) are in successful use for the new advanced steam power plants. Development of 10–11%Cr steels for steam turbines up to 650°C Based on the experience with the steels developed for temperatures up to 620°C, the alloying principles needed to obtain a further increase in the creep strength of 100 MPa for 100 000 h at the application temperature and to improve the resistance to oxidation are very similar in the Japanese and European research programmes: •
increase in Cr-content of 10–11% to improve the resistance to steam oxidation;
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The development of creep-resistant steels Steel (a) (b) (c) (d) (e)
C 0.10 0.20 0.10 0.10 0.10
Cr 2.50 12.0 9.0 9.0 9.0
Mo 1.0 1.0 1.0 1.0 0.45
W – – – 1.0 1.8
V – 0.30 0.20 0.20 0.20
Nb – – 0.05 0.05 0.06
N – – 0.05 0.07 0.05
33
B (mass%) – – – – 0.002
200 (c) P91 (d) E911 (e) P92 (b) 12%CrMoV 100 (a) 2.25%CrMo 25°C 40°C 0 500
550
600
650
Temperature (°C)
2.14 100 000 h creep rupture strength of steam plant piping steels as a function of test temperature.
• • • • •
reduction in Si, Mn- and Ni-content to the lowest possible level to improve the creep strength; addition of 3–6% Co to reduce the diffusion rate and to prevent the formation of delta-ferrite; addition of 0.002–0.018% boron to stabilise the M23C6 carbides by forming M23(C,B)6; addition of 3%W (Japan) or 1.5%Mo (Europe) to improve the solution hardening and to stabilise the M23C6 and M23(C,B)6 carbides; study of the influence of C content on diffusion rate and on M23C6 precipitates.
Table 2.2 provides an overview of the chemical composition of Japanese and European test melts. None of the referenced analysis versions meets the needs for use at 650°C. A direct comparison of the creep strength of the tested versions is only possible in part because, as a rule, the results are only shown in the form of Larson and Miller diagrams. The investigations with the MTR and TOS/JSW test melts resulted in the development of rotor steels MTR 10 A and TOS 110 which were also tested in large pilot rotors.55,57 For these two steels a maximum application temperature of 630°C is given. By way of example, Fig. 2.15 shows the determined creep strength of the weakest version (FB6) and the strongest version (FB8) of the COST programme at 650°C in comparison with the master steel FB2.58,59 At test periods below 10 000 h, the test melts FB6 and FB8 undercut the creep strength curve of the 9% Cr master steel FB2 which is suitable for use up to about 625°C. The
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Table 2.2 Chemical composition of rotor test melts investigated in Europe and Japan (mass%) Type
Japan HR1200/20t ingot HR1200/80t ingot HR1200/C HR1200/D HR1200/E MTR various test melts MTR10A pilot rotor TOS/JSW test melts
TOS110 pilot rotor nominal
Si
Mn
Cr
Mo
W
Co
Ni
V
Nb
N
B
0.13
0.05
0.80
9.3
1.50
–
1.0
0.15
0.20
0.05
0.020
0.010
0.13 0.14 0.16 0.17 0.18 0.12 0.19
0.09 0.10 0.12 0.09 0.08 0.03 0.02
0.09 0.10 0.08 0.09 0.08 0.12 0.11
10.1 11.2 10.2 11.1 11.0 10.3 11.1
1.46 1.45 1.54 1.46 1.48 0.90 0.90
– – – – – 0.42 0.42
2.85 2.93 2.97 2.94 6.03 3.02 3.02
0.15 0.15 0.17 0.20 0.13 0.17 0.17
0.20 0.22 0.20 0.21 0.20 0.21 0.21
0.06 0.08 0.07 0.07 0.06 0.07 0.07
0.017 0.016 0.019 0.023 0.017 0.030 0.020
0.010 0.009 0.009 0.010 0.009 0.010 0.009
0.09 0.10 0.12 0.13 0.14 0.10– 0.12 0.11 0.11 0.11 0.11 0.11 0.11 0.11 0.11 0.11 0.11
0.02 0.06 0.03 0.03 0.03 0.04– 0.10 0.05 0.04 0.03 0.04 0.05 0.03 0.03 0.03 0.03 0.08
0.50 0.46 0.50 0.18 0.15 0.06– 0.50 0.06 0.08 0.08 0.08 0.08 0.08 0.08 0.08 0.08 0.10
11.0 10.2 10.3 10.5 10.1 10.1– 11.6 10.2 10.3 10.2 10.2 10.1 10.1 10.1 10.2 10.2 10.0
0.23 0.14 0.16 0.14 0.18 0.33– 2.50 0.71 0.70 0.70 1.51 1.54 0.69 0.72 0.70 0.70 0.65
2.60 2.50 2.51 2.40 2.41 0– 1.96 1.76 1.80 1.76 – – 1.81 1.78 1.81 1.78 1.80
2.50 2.40 1.18 1.17 1.45 0– 6.0 3.7 3.02 2.98 2.95 2.98 2.95 2.97 2.97 2.97 3.0
0.51 0.25 0.04 0.11 0.18 0.02– 0.30 0.06 0.20 0.21 0.21 0.19 0.20 0.20 0.20 0.20 0.20
0.22 0.21 0.19 0.23 0.20 0.17– 0.21 0.20 0.20 0.20 0.20 0.20 0.20 0.20 – – 0.02
0.07 0.07 0.07 0.07 0.07 0.05– 0.06 0.05 0.06 0.06 0.06 0.06 – – 0.06 0.06 0.05
0.020 0.020 0.018 0.023 0.018 0.021– 0.076 0.029 0.018 0.020 0.018 0.018 0.022 0.019 0.019 0.020 0.020
0.018 0.013 0.011 0.011 0.011 0– 0.008 0.004 – 0.009 – 0.008 – 0.009 – 0.008 0.010
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Creep-resistant steels
COST Programme Master steel: FB2 (625°C) Test Melts: FB5 FB6 FB7 FB8 FB9 FB10 FB11
C
The development of creep-resistant steels
35
Creep rupture strength (MPa)
200
650°C
FB2
100 80 FB6 60 FB8 40
100
Test temperature 650°C
1000 10 000 Time to fracture (h)
100 000
2.15 Creep rupture strength of the COST test steels FB2, FB6 and FB8 as a function of time to rupture at 650°C.
650°C creep strengths of the HR1200/20t ingot, the HR1200/80t ingot and the HR1200/C are shown in Fig. 2.16.60 After 10 000–20 000 h the HR1200/20t and the HR1200/80t versions (variants A and B) reveal a distinct reduction in creep rupture strength. Based on published test progress at 15 000 h, the long-term behaviour of the HR 1200/C version cannot be evaluated. However, the target of 100 MPa at 100 000 h will not be reached even after a further stable pattern of the creep rupture strength curve. According to the results of investigations into the microstructure of versions FB6 and FB8 of the COST programme, the marked decline in creep strength is mainly caused by precipitation of the Z-phase Cr2(Nb,V)2N2 which is formed at the expense of MX precipitations.54,61 A further striking fact is a remarkable coarsening of M23C6 precipitations and the Laves phase. A significant coarsening of the M23C6 precipitations and the Laves phase was also determined in investigations of versions HR1200/20t and HR1200/80t.60 This publication, however, does not provide any information on MX and Z-phase precipitations. It is significant when considering the interpretation of the investigation results into the microstructure of these rotor test melts that the long-term creep-stressed specimens of the thermally very stable steels TAF, B2 and FB2 revealed no Z-phase precipitations, but fine MX precipitations and only minor coarsening of the M23C6 precipitations.61,62 The influence of 8.56– 11.59%Cr on the creep strength at 650°C for a steel which in addition is alloyed with 0.10%C–3%W–3%Co–0.70%Mo–0.15%V–0.06%Nb–0.02%N– 0.01%B has been reported.63 Figure 2.17 presents an overview of the results. The summary includes details of the creep rupture time as a function of the
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36
Creep-resistant steels Test temperature 650°C
300 Ingot sizes: 20 t (A) 80 t (B) small (C) small (E)
Rupture strength (MPa)
200
A
100 80 60
40 30 100
C Si Mn Cr Mo Ni W V Nb N B Co Al
HR1200 variants B C
E new design 0.09 0.09 0.10 0.14 0.02 0.06 0.03 0.03 0.50 0.47 0.50 0.15 11.0 10.2 10.3 10.1 0.23 0.14 0.16 0.18 0.51 0.25 0.04 0.10 2.6 2.59 2.51 2.41 0.22 0.21 0.19 0.20 0.07 0.07 0.07 0.07 0.02 0.018 0.018 0.018 0.018 0.012 0.011 0.011 2.53 2.46 1.18 1.45 0.007 <0.005 0.005 0.001 1000
10 000 Time to fracture (h)
100 000
2.16 Creep rupture strength of the rotor test steels HR1200 (Variants A-B-C) as a function of time to rupture at 650°C.
Cr content and test stresses in the range 300–98 MPa. In the long-term range, that is, at lower test stresses, the alloying procedure with approximately 9%Cr exhibits the highest creep strength whereas in the short-term range the alloy with about 11.5%Cr reveals the highest creep strength. Investigations into the microstructure of the 157 MPa stressed creep specimens of the 9.5%Cr and 11.5%Cr versions showed the following results after creep rupture periods of 5289 h and 1928 h respectively. For the first set result the 11.5%Cr steel, the dislocation density decreased, subgrains were formed and only a minor fine Laves phase could be observed within the lath structure. For the second set result the 9.5%Cr steel on the other hand, the lath width remained narrow and the dislocation density was higher than for the 11.5% Cr steel and in addition, a significant amount of fine Laves phase had still been precipitated within the lath structure. The conclusion drawn was that the coarsening of precipitates, such as the Laves phase, was increased and that recovery of the microstructure could not be suppressed in the higher Cr steels. Development of 11% Cr steels for pipes and headers for 650°C steam plants Based on experience with the steels developed for temperatures up to 620°C, the alloying principles needed to obtain a further increase in the creep strength
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The development of creep-resistant steels
37
Heat treatment 5 h 1070°C +20 570°C +20 h 680°C
98MPa
104 118MPa
137MPa Metallographic investigation
Time to rupture (h)
157MPa
103 177MPa
210MPa
240MPa 102
270MPa
101
300MPa 8
9
10 11 Cr content (%)
12
2.17 Creep rupture strength of 8.5–12%Cr–3.5%W–3%CoVNbB test steels as a function of Cr content, applied stress and time to fracture at 650°C.
to reach 100 MPa for 100 000 h at the application temperature and to improve the resistance to oxidation are similar in the research programmes undertaken in Japan and Europe:64–66 • • • •
increase in Cr-content to about 11% to improve the resistance to steam oxidation, addition of 1–3%Co to reduce the diffusion rate and to prevent the formation of delta-ferrite, addition of 0.003–0.012% boron to stabilise the M23C6 carbides by forming M23(C,B)6, addition of 0.3–3%W and/or 0.15–1.5%Mo (Europe) to improve solution hardening and to stabilise the M23C6 and M23(C,B)6 carbides,
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38
•
Creep-resistant steels
addition of 0.07%Ta and 0.04%Nd to improve the creep strength by fine and stable nitrides.
Table 2.3 provides an overview of the chemical composition of Japanese and European test melts. None of the referenced analysis versions meets the needs for use at 650°C. Figure 2.18 shows a comparison of the published results for steels VM12, NF12 and SAVE 12 with regard to the target creep strength at 650°C.67–69 The behaviour is comparable with the creep strength established for the rotor steels in Table 2.2, that is, with 10–11%Cr and Cocontents of roughly 3%. Results of the investigations into the microstructure have hitherto only been published for version NF12. After creep stressing for 15 000 h at 650°C it was similarly discovered that Z-phase precipitation took place at the expense of MX precipitation.70 Japanese NIMS research project STX 21 for the development of thicksection boiler components The Japanese research project STX 21, started in 1995,71 embraces a most comprehensive 15-year investigation into the influence of the alloying elements on the creep strength and oxidation resistance of the ferritic–martensitic steels in thick-section boiler components for advanced power stations. The systematic investigation into all relevant elements in regard to material property requirements for main steam pipes and headers covers: • • • • •
105 h creep rupture strength at 650°C oxidation resistance in steam weldability and creep rupture strength of welded joints thermal fatigue impact properties and hot workability.
Figure 2.19 demonstrates the design philosophy of a 9%Cr steel72 with the alloying partners W, Mo, Ni, Cu, Co, Si, V, Nb, Ta, C, N and B and their specific property profile. In addition, further analysis versions were tested. Table 2.4 provides an overview of investigated element variations within the project. Reports on test progress were issued on a regular basis at international conferences. The best version was established to be a steel with 0.08%C– 9%Cr–3%W–3%Co–0.20%V–0.05%Nb–0.008%N–0.014%B.73 It was found that the addition of boron at more than 0.01% to the 9%Cr steel remarkably improves the long-term creep strength. Boron stabilises the lath martensitic microstructure of 9%Cr–3%W–3%Co steels during creep deformation at 650°C through the stabilisation of M23C6 in the vicinity of prior austenite grain boundaries by an enrichement of boron in the M23C6 carbides. Further improvement of creep strength in this high boron-containing steel is accomplished by the addition of 0.008% nitrogen enhancing precipitation of fine MX. The steel also exhibits good creep–rupture ductility.
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Table 2.3 Chemical composition of pipes and header test melts investigated in Europe and Japan (mass%) Type
Si
Mn
Cr
Mo
W
Co
Ni
V
Nb
N
B
0.10 0.16 0.16
0.35 0.15 0.30
0.45 0.09 0.22
8.75 11.0 11.1
1.00 1.04 1.04
– 0.32 0.32
– 2.12 2.16
0.20 0.14 0.43
0.21 0.26 0.27
0.080 0.065 0.067
0.050 0.059 0.060
– 0.009 0.011
0.10 0.11 0.15
0.40 0.40 0.30
0.50 0.50 0.22
9.0 11.5 11.2
0.45 0.28 0.50
1.8 1.4 2.1
– 1.30 2.15
0.15 0.23 0.44
0.20 0.24 0.26
0.060 0.065 0.068
0.045 0.056 0.063
0.003 0.003 0.009
0.17 0.19 0.18
0.07 0.10 0.10
0.06 0.52 0.51
9.34 11.6 11.5
1.55 1.48 1.48
– – –
– – –
0.12 0.20 0.25
0.27 0.26 0.26
0.064 0.058 0.058
0.015 0.015 0.047
0.010 0.004 0.004
0.13 0.16 0.16 0.13 0.16 0.12
0.05 0.09 0.10 0.10 0.10 0.50
0.82 0.53 0.50 0.49 0.50 0.35
9.32 11.3 11.4 11.3 11.4 11.5
1.47 1.46 1.48 1.44 1.46 0.30
– – – – – 1.5
0.96 0.90 1.06 3.06 1.49 1.6
0.16 0.26 0.26 0.28 0.27 0.30
0.20 0.25 0.25 0.25 0.27 0.25
0.05 0.047 0.058 0.057 0.044 0.050
0.019 0.063 0.063 0.051 0.050 0.065
0.009 0.006 0.012 0.008 0.008 0.005
0.10 0.10
0.35 0.30
0.50 0.20
10.2 11.0
0.15 –
2.5 3.0
2.0 3.0
<0.10 <0.10
0.22 0.20
0.07 0.07
0.02 0.04
0.005 0.07Ta 0.04Nd
63
Japan Results of investigations of various test melts NF1265 SAVE 1265
39
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The development of creep-resistant steels
COST programme Master steel P91: F1 F2 Master steel P 92: F F3 Master steel B2: A B Master steel FB2: C D E G Pilot pipes VM1264
C
Creep-resistant steels 300
Test temperature 650°C
200
Rupture strength (MPa)
40
SAVE12
NF12
100 80 VM12 60 P92 (mean ECCC) 40 30 100
1000
10 000 Time to fracture (h)
100 000
2.18 Creep rupture strength of piping test steels NF12, SAVE 12 and VM12 as a function of test time at 650°C in comparison to the behaviour of the 9% Cr piping steel P92.
Cr < 8
Cr = 9
Oxidation resistance decrease Addition of Si W + 2Mo < 2
W + 2Mo = 3
Ni = 0.5(Cu, Co)
Ni > 0.5(Cu, Co) Creep strength decrease Ac1 temperature decrease Additon of Ir, Rh, Pd
δ ferrite formation
Si < 0.3
W + 2Mo > 3 Toughness decrease
Creep strength decrease Ni = 0
Cr > 11 δ ferrite formation Addition of austenite stabilising elements
Si = 0.5
Oxidation resistance decrease Addition of Y
Si > 0.6 Carbide agglomeration Acceleration of Fe2W precipitaton
V, Nb, Ta V = 0.2, Nb = 0.05, Ta = 0.1–0.2 C, N, B
Fine dispersion of MC by TMCP Increase in MC volume fraction
C = 0.08–0.1, N = 0.03 –0.06, B = 0.005–0.008 Creep strength decrease
Carbide agglomeration Weldability decrease
2.19 NIMS alloy design philosophy of high Cr ferritic steels for 650°C ultrasupercritical (USC) boilers.
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Table 2.4 Chemical composition of test steels for pipes and boilers of the NIMS STX 21 project (mass%) C
Si
Mn
Cr
Mo
W
9CrW
0.10
0.30
0.50
9.0
–
9CrWMo
0.08
0.30
0.50
9.0
9CrWN
0.10
0.30
0.50
9.3
3 2.8 2.5 1.9 <0.01 –
– 1 2 4 0.003 0.10 0.33 0.62 1.54 3.3
9CrWB 9CrCoTiAl
0.08 0.10
0.30 0.01
0.50 0.01
9.0 9.0
– –
–0.01 0.30 0.50
<0.01
Var. 8.5 10
–
0.14
9–12Cr WSiTiY (Oxidation– tests)
V
Nb
N
Co
B
0.02
0.22
0.05
0.05
0.20
0.05
3.0
0.005
3.0 –
0.20 0.20
0.05 0.05
0.054 0.026 0.001 0.02 0.05
3.0 3.0
0+0.014
2.0
0.20
0.05
0.05
–
9CrWCoPd
0.08
0.80 1.0 0.30
9CrWVTa
0.10
0.30
ODS 10CrTiY2O3
0.065 –0.12
9 –13
–
2
ODS 12CrWY2O3
<0.05
12.3
–
0–3
0.014 (Ti+Al) Var. Ti 0
12 0.50
9.0
–
3.0
0.20
0.05
0.05
0.50
9.0
–
3.0
0.20
0.05
0.05
0.005
Var. Al 0.001
0.10
0.008
Y 2 O3 0.08 –0.35 Y 2 O3 0.24 –0.25
41
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0+3
0.05 0.10 0Pd 1Pd 3Pd Ta 0.02 Ti 0.12 +.22 Ti 0–0.39
Var.Y 0 0.05
The development of creep-resistant steels
Type
42
Creep-resistant steels
Figure 2.20 presents an overview of the influence of B and N on creep strength of 9Cr3W3CoVNbNB steel at 650°C in comparison with the TAF and P92 steels. The creep strength of the best 9%Cr version with 0.014%B and 0.008%N correlates in the long-term range with the creep strength of the TAF steel. Considering the results hitherto available, the authors of this publication conclude that, based on the present chemical composition, a 100 000 h creep strength of 100 MPa at 650°C is achievable. Protective Crrich scale is formed on the surface of 9%Cr by a combination of Si addition and pre-oxidation treatment in argon gas at 700°C for 20 h. This significantly improves oxidation resistance of 9%Cr steels at 650°C.74
2.4
Historical development of austenitic steels
2.4.1
History of austenitic steels
Nickel alloyed austenitic steel originated from 25%Ni–Fe alloy and 25%Ni– 5 to 8%Cr–Fe alloy melted by Krupp in Germany in 1893 and 1894, respectively.75 Krupp also produced 35%Ni–13 to 14%Cr–Fe and 25%Ni–8 to 15%Cr–Fe alloys for use in thermocouple sheaths and molten glass mould material in 1910.75 Krupp continued the development of a series of Ni–Cr– Fe steels and identified a martensitic 10%Cr–2%Ni steel and an austenitic Test temperature 650°C 300
Rupture strength (MPa)
200 TAF (Fujita)
100 80 9Cr3W3CoVNbB(NIMS) 60
40 30 100
0.0139B/0.0034N 0.0092B/0.0016N 0.0048B/0.0011N 0.0135B/0.0079N 0.0100B/0.0200N
1080/800°C 1050/790°C 1050/790°C 1150/770°C 1080/800°C 1000
10000 Time to fracture (h)
P92(ECCC)
100 000
2.20 Creep rupture strength of 9Cr3W3CoVNbB NIMS test steel as a function of B and N content, heat treatment and time to fracture in comparison to the behaviour of the 9% Cr piping steel P92 and the TAF steel.
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The development of creep-resistant steels
43
20%Cr–5%Ni steel in 1912 as stainless steels.76 The latter was designated as V2A in 1922 (V for Versuchstahl, respectively test steel, 2 as the development number, and A for austenitic) and served as an original austenitic steel for use after solution annealing for corrosion resistance in vessel piping, machinery, and so on in acid environments.77 Conventional 18%Cr–8%Ni austenitic steel originated from V2A in the mid-1920s, optimising the content of Cr and Ni to maximise the economical benefit by keeping the structure austenitic.78 Since these austenitic steels suffered severe intergranular corrosion in the weld heat affected zone, extensive investigation into the corrosion mechanism and improvement of the steels was carried out at Krupp around 1930. Consequently, stabilised austenitic steels were developed employing alloying elements Ti, V, Nb and Ta, as well as reducing the carbon content to a maximum of 0.07% and achieving a fine grained steel with finer distribution of chromium carbide.79 These steels are original versions of the conventional Type 321 and Type 347 steels. Type 316 originated from 18%Cr–8%Ni steel containing 3% molybdenum for use in ammonium chloride and sulphuric acid environments as a corrosion-resistant alloy. At that time (early 1930s), addition of the combination of Mo and Cu was found to deliver improved corrosion resistance against sulphuric acid and 18%Cr–8%Ni–2%Mo–2%Cu was thus developed by Krupp. From the viewpoint of better cold workability of austenitic steels, austenite forming elements such as Mn and Cu were increasingly used to develop several new steels, namely 19-9LW(19%Cr–9%Ni–1.25%Mo–1.25%W–NbTi), 199DX(19%Cr–9%Ni–1.5%Mo–1.2%W-Ti) and 17-14CuMo81 modified from Krupp V6A by Armco Steel in USA.80 These steels with stabilised austenitic structures exhibited not only corrosion resistance to sulphuric acid but also resistance to heat and creep. Figure 2.2182 shows the elevation of creep rupture strength of heat-resistant steels for boilers, viewed in terms of change in 100 000 h creep rupture strength at 600°C, for materials developed during the 20th century. After World War II, 18%Cr–8%Ni steels previously developed in Germany before the war were used for heat-resistant purposes as well as in chemical plants worldwide, thereby increasing the steam pressure and temperature of fossilfired power plants. Also, ultra-supercritical pressure power plants constructed in the late 1950s were realised by further applying these austenitic steels to thick walled components. For example, TP316H was used for boiler headers and steam piping and 17-14CuMo and TP321H for superheater and reheater tubes at Eddystone unit No. 1. Esshete 1250 (15%Cr-10%Ni-6%Mn1%MoVNbB)83 austenitic steel was also used for thick section components in the UK. In the 1950s high nickel Alloy800H was developed and put in service in the USA as a high strength and anti-corrosive steel substitutable for Nibased alloys. We must remember that these austenitic steels were basically
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Creep-resistant steels
NF709
105 h Creep rupture strength at 600°C (MPa)
200
RH3C AA-1 XA-704 Spr304H A-3 Next generaton 12Cr-WCoVNb
Stable austenitic 150
17-17CuMo E1250 800H
A-1 TP347HFG
HCM2S T24 (2Cr)
Meta-stable austenitic
50 Ferritic Ni-Cr 0 1900
Cr-Mo
1920
(2Cr) T22
3rd generation
NF616 HCM12A E911 HCM12 2nd generation T91
TP347H TP316H TP321H TP304H
100
SAVE25 HR6W
1st generation F-9 EM12 HCM9M HT9 HT91
410 T9 (9-12Cr)
1940
1960 Year
1980
2000
2020
2.21 Historical improvement in creep rupture strength of boiler steels.
developed for corrosion-resistant materials and unexpectedly showed good creep-resistant properties. From the perspective of the improved creep resistance of the 18%Cr–8%Ni system, austenitic steel alloy development studies started in the 1960s. Tempaloy A-184 was developed through the optimisation of carbon stabilising elements of Ti and Nb in 18%Cr–8%Ni steel in the early 1970s, followed by thermo-mechanically strengthened TP347H with finer grain structure developed in the early 1980s, designating TP347HFG85 for use in superheaters and reheaters as a steam oxidation-resistant and creepresistant material. Subsequently, several other 18%Cr–8%Ni austenitic steels with highly improved creep strength, such as Super304H,86 XA70487 and Tempaloy AA-188 were developed. Within the group of high Cr and high Ni austenitic steels, several new 20– 25%Cr austenitic steels with nitrogen and relatively lower Ni content were developed from the 1980s to the 1990s. As seen in Fig. 2.21 (for reference), the improved creep strength of ferritic steels is notable, along with the improvement of austenitic steels. Regarding ferritic steels, low alloy steels
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The development of creep-resistant steels
45
or 9–12% Cr steels with about 40 MPa of 100 000 h creep rupture strength (half or less than that of 18%Cr–8%Ni steels) had been used for many years and the problem of cost increases existed because, especially in the case of superheater and reheater applications, there was an alloy ‘gap’ between the low alloy steels and the 18%Cr–8%Ni steels, arising as a result of temperature elevation. Accordingly, research work for high strength 9–12%Cr steels was initiated in order to fill this gap and 60 MPa class materials (first generation) were developed over the period from 1950 to 1970. Further efforts have been made and creep strength reached the 100 MPa class (second generation) in the 1980s, whereas the 130 MPa class (third generation) was achieved in the 1990s. Materials for 150 MPa at 600°C or 100 MPa at 650°C, as the next generation, are expected to emerge. In general, the outside diameter and wall thickness of pipes and tubes can be greatly reduced through elevation of creep rupture strength. Thermal stresses can accordingly be lowered and construction of fossil-fired power plants capable of load sliding operation under further elevated pressure and temperature steam conditions will be possible.
2.4.2
Alloy design of creep-resistant austenitic steels
Heat-resistant steels for practical application must be designed by taking their service conditions and environments into consideration and by examining their various properties.89 However when alloy design is performed based on modification of existing steels, both oxidation and corrosion resistance as well as their general material properties are expected to be nearly equivalent to those of the original materials. Hence, chemical compositions and heat treatment conditions are examined with particular consideration of creep strength improvement. Figure 2.2289 shows the concept of alloy design for heat-resistant austenitic steels to improve creep strength through the modification of existing steels. In the case of austenitic steels, chemical composition can be largely classified into the four categories shown in the figure, and solution strengthening and precipitation strengthening are designed specifically for each of these categories. 18%Cr–8%Ni steels based on Type 304 steels include Type 316 steels solution-strengthened through the addition of Mo, as well as Type 321 steels and Type 347 steels precipitation-strengthened through the addition of Ti or Nb. However, these materials were originally developed for chemical equipment as mentioned above, placing emphasis on corrosion resistance, but were not designed from the standpoint of creep strengthening. Accordingly, further enhancement of precipitation strengthening by means of ‘under-stabilising’90 C and/or composition design for improved creep strength is used. 15%Cr– 15%Ni or 21%Cr–30%Ni steels with a full austenitic phase structure are capable of high creep strength in the ‘as-is’ condition, although they are
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46
Austenitic steel
18%Cr-8%Ni
Precipitation strengthening
Type 316 (Mo)
Type 304
Stabilising
Under stabilising
Cr carbide Type 321 (Ti), TiC
Micro alloying of Nb, Ti, B
Type 347 (Nb), NbC
High strength 18%Cr-8%Ni steel
Lower (Ti + Nb)/C
Alloy design for creep
15%Cr-15%Ni (15%Cr-10%Ni)
Type 17-14CuMo
25%Cr-20%Ni
Type 310
21%Cr-32%Ni
Type alloy800H
2.22 General concept of alloy design for austenitic heat-resistant steels.
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Cu, Mo, W addition
Creep resisting steel
N addition
Creep/ corrosionresistant steel
Low-cost creep/ corrosionresistant steel
Creep-resistant steels
Solution strengthening
The development of creep-resistant steels
47
costly because of their high Ni content. Steels containing 20% and more Cr are likely to have excellent oxidation and corrosion resistance, but a costly Ni content of at least 30% is required to maintain a full austenitic structure. Nevertheless, low-cost, high-strength, highly corrosion-resistant austenitic steel can be designed by adding about 0.2% N to reduce the Ni content and by combining the strengthening mechanism as described above. The features of alloy design of austenitic heat-resistant steels are discussed below. As previously noted, ‘under-stabilising’ is a technique for improving the creep strength of 18%Cr–8%Ni steels. This method enhances creep strength through the improvement of precipitation morphology by fixing C in alloys and decreasing carbide forming elements such as Ti and Nb, which hinder Cr carbide formation, to the point where their contents are insufficient for C fixation. Figure 2.2390 shows this, and the peak point of the creep rupture strength against the ratio of (Ti + 0.5Nb)/C is at position far away from the peak point of conventional Type 321 or Type 347 steels, indicating that reducing additions of Ti and Nb relative to the C content can be useful. Figure 2.2486 shows the effect of the Cu and Nb additions on the creep rupture strength of 18%Cr–9%NiNbN steel and 18%Cr–9%NiCuN steel, respectively. Although Cu addition does not show a major change up to about 2%, a substantial enhancement in creep strength is observed owing to the very finely dispersed Cu-rich particle precipitates by means of Cu addition
105 h Creep rupture strength at 650°C (MPa)
130 120 110 100 90 80 70 321H 347H
60 50 40 30 0.01
0.01
0.1
0.1
0.5 1.0 (Ti + Nb)/C Atomic ratio 0.5 1.0 1.5 2.0 3.0 4.0 (Ti + 0.5Nb)/C Mass ratio
1.5
2.0
5.0 6.0 7.0 8.0
2.23 Effect of (Ti + Nb)/C ratio on creep rupture strength of 18Cr10NiNbTi steel.
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48
Creep-resistant steels 18Cr-9NiNbN
18Cr-9NiNuN
103
Time to rupture (h)
Time to rupture (h)
103
102
102 700°C 167MPa 137MPa
700°C/137MPa 750°C/108MPa 10 0
2
4
6
Cu (%)
10 0
0.2
750°C 137MPa 108MPa
0.4 0.6 Nb (%)
0.8
1.0
2.24 Effect of Cu and Nb on creep rupture strength of 18Cr–8Ni steels.
of about 3% or more. However, because the strength tends to be saturated, while a decline in creep rupture ductility can occur when the Cu addition exceeds 3%, the addition of Cu at 3% should be suitable. On the other hand, the effect of Nb addition on the creep rupture strength of 18%Cr–9%NiCu steel is significant when Nb content exceeds 0.2% and is saturated at a concentration of above 0.4%, which is not sufficient to stabilise the 0.08%C usually contained in 18%Cr–8%Ni H grade steels. This was mentioned above as an effect of under-stabilising on the creep strength enhancement. Super304H and other Cu-alloyed and under-stabilised steels were thus developed.
2.4.3
Development of austenitic heat-resistant steels
Austenitic steels for heat exchanger and boiler tube applications Chemical compositions of austenitic heat resistant steels for boiler applications are given in Table 2.5, with development progress in austenitic boiler steels presented in Fig. 2.25.91 Considering steels for boiler applications, various improvements have been made to enhance creep strength in the grade of 18%Cr–8%Ni system steel because this grade is used in the highest temperature region of superheater and reheater sections of boilers with conventional steam parameters. The steam oxidation resistance of some 18%Cr–8%Ni steels is also improved at their inner surface through finer grains. Furthermore, new austenitic steels with a Cr content of 20% or more have been developed for the purpose of improving creep strength and corrosion resistance.
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Table 2.5 Nominal chemical composition of austenitic steels for boilers Steels
18%Cr– 8%Ni
20%– 25%Cr
Si
Mn
Ni
Cr
Mo
W
V
Nb
Ti
B
Others – – – – – – 3.0Cu, 0.1N 0.2N 3.0Cu 3.0Cu 0.1N 0.1N – – – 0.2N 0.4Al 0.15N – 3.0Cu, 0.2N 3.0Cu, 0.2N 0.03Zr –
TP304H TP316H TP321H TP347H TP347HFG Tempaloy A-1 Super304H
18Cr–8Ni 16Cr–12NiMo 18Cr–10NiTi 18Cr–10NiNb 18Cr–10NiNb (FG) 18Cr–10NiNbTi 18Cr–9NiCuNbN
0.08 0.08 0.08 0.08 0.08 0.12 0.10
0.6 0.6 0.6 0.6 0.6 0.6 0.2
1.6 1.6 1.6 1.6 1.6 1.6 0.8
8.0 12.0 10.0 10.0 10.0 10.0 9.0
18.0 16.0 18.0 18.0 18.0 18.0 18.0
– 2.5 – – – – –
– – – – – – –
– – – – – – –
– – – 0.80 0.80 0.10 0.4
– – 0.50 – – 0.08 –
– – – – – – –
XA704 Tempaloy AA-1 17–14CuMo 15–15N AN31 Esshete1250 12R72 TP310 HR3C Alloy 800H Tempaloy A-3 NF709 SAVE25
18Cr–9NiWVNb 18Cr–10NiCuTiNb 17Cr–14NiCuMoNbTi 15Cr–15NiMoWNbN 15Cr–13NiMoNbN 15Cr–10Ni6MnVNbTi 15Cr–10NiMoNbVB 25Cr–20Ni 25Cr–20NiNbN 21Cr–32NiTiAl 22Cr–15NiNbN 20Cr–25NiMoNbTi 22.5Cr–18.5NiWCuNbN
0.03 0.10 0.12 0.12 0.10 0.12 0.10 0.08 0.08 0.08 0.05 0.15 0.10
0.3 0.3 0.5 0.7 0.5 0.5 0.4 0.6 0.4 0.5 0.4 0.5 0.1
1.5 1.5 1.7 1.5 1.5 0.6 2.0 1.6 1.2 1.2 1.5 1.0 1.0
9.0 10.0 14.0 15.0 14.0 10.0 15.0 20.0 20.0 32.0 15.0 25.0 18.0
18.0 18.0 16.0 15.0 16.0 15.0 15.0 25.0 25.0 21.0 22.0 20.0 23.0
– – 2.0 1.5 1.5 1.0 1.0 – – – – 1.5 –
2.5 – – 1.5 – – – – – – – – 1.5
0.3 – – – 0.5 0.2 – – – – – – –
0.3 0.3 0.40 1.0 1.0 1.00 – – 0.45 – 0.70 0.20 0.45
– 0.2 0.30 – – 0.06 0.3 – – 0.50 – 0.10 –
– 0.02 0.008 – – – 0.006 – – – 0.002 – –
Sanicro25
22Cr–25NiWCuNbN
0.08
0.2
0.5
25.0
22.0
–
3.0
–
0.3
–
–
30Cr–50NiMoTiZr 23Cr–43NiWNbTi
0.06 0.08
0.3 0.4
0.2 1.2
50.0 43.0
30.0 23.0
2.0 –
– 6.0
– –
– 0.18
0.20 0.08
– 0.003
High Cr– CR30A High Ni HR6W
WPNL2204
49
C
The development of creep-resistant steels
15%Cr– 15%Ni–
Chemical composition (mass%)
(85)
(74)
20Cr-25NiMoNbTi
18Cr-9NiWVNb
18Cr-8Ni, C<0.08 (40–60)
(AISI 304) +Ti 18Cr-10NiTi (AISI 321) +Nb 18Cr-8Ni (AISI 302)
18Cr-10NiNb (AISI 347) +Mo 16Cr-12NiMo
H Grade 0.04 – 0.10C AISI AISI AISI AISI
304H 321H 347H 316H (75)
NF 709 (SUS310J2TB)
(53) 21Cr-32NiTiAl (Alloy 800H)
(66) 22Cr-15NiNbN
(ASME TP347HFG) Chemistry Optimization (58) 18Cr10NiNbTi Tempaloy A-1 (SUS321J1HTB) Cu Addition
(70)
18Cr9NiCuNbN
17Cr-14Ni CuMoNbTi
Tempaloy A-3 (SUS309J4HTB) Cu Addition (76) 18Cr10NiCuTiNb (89) 23Cr-43NiWNbTi
Tempaloy AA-1 ASME TP 321HCuCb SUS321J2HTB
HR6W
High Cr-High Ni
Super 304H
(AISI 316)
(73) 30Cr-50NiMoTiZr
ASME TP304CuCbN SUS304J1HTB
(91)
Cr30A
22Cr-25NiWCuCoNbN +Cr –Ni
Sanicro 25 (37) 22Cr-12Ni (AISI 309)
(90)
(65)
25Cr-20Ni (AISI 310)
25Cr-20NiNbN
23Cr-18.5NiWCuNbN
HR3C ASME TP310CbN SUS310J1TB
SAVE25 (SUS310J3TB)
( ) Designates 105 h Creep rupture strength (MPa) at 700°C.
2.25 Development progress of austentic boiler steels.
WPNL2204
Creep-resistant steels
–C
XA704 (ASME TP347W) (SUS347 J1TB) Heat Treatment (58) 18Cr10NiNb
50
W Addition
The development of creep-resistant steels
51
18Cr–8Ni steels such as TP304H, TP321H, TP316H and TP347H grades are still used for fossil-fired power plants operating under conventional steam conditions. TP347H, which has the highest allowable stress among these four types, was improved in order to provide a fine-grained structure (grain size number 8 and finer) for steam oxidation resistance and creep strengthening, designated as TP347HFG by ASME. The effect of solution temperature on the creep rupture strength and grain size in thermo-mechanical processes is schematically illustrated in Fig. 2.26.85 This steel is highly effective in improving the reliability of superheater tubes, being applicable to ultrasuper critical pressure power plants up to 600°C class and already fully employed in the superheater tubes of a substantial number of these plants. Since the 15%Cr–15%Ni system steel is stable austenitic and high strengths in creep are likely to be obtained, their allowable stresses are very high. The mechanism of the strengthening of these steels is strong solution hardening by Mo and W, and precipitation hardening effects by inter-metallic compounds as well as by MC and M23C6 carbides make a contribution. Table 2.692 shows the nominal chemical compositions of typical elements forming intermetallic compounds and carbides, Nv – Nc (Nv – Nc is the excess electron vacancy number by PH ACOM which is defined in ASTM A567, v = vacancy, Developed TP347HFG Anneal
Temperature
Grain size, ASTM No.
10 8
NbC Resolved
Precipitated Cold drown
Solution treatments Fine graln
Cold crown
Solution treatments Coarce graln
6 4 2
Developed Conventional
120
105 h creep rupture strength (MPa)
Anneal NbC
650°C × 105 h
100 80
700°C × 105 h ASME A.S./0.67
60 ASME A.S./0.67
40
A.S.: Allowable stress 20 1100
1200 1300 Solution temperature (°C)
1400
2.26 Effect of solution temperature on creep rupture strength and grain size in a thermo-mechanical process.
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52
Creep-resistant steels
Table 2.6 Comparison of major alloying elements and schematic microstructures in 15Cr–15Ni austenitic steels Steel
Mo
W
Nb
Ti
Schematic microstructures NV – NC (700°C) after ageing at 700°C for 104 h
17–14CuMo
2.12
–
0.67
0.28
0.117
15–15N
1.58
1.49
1.07
–
0.096
(Nb, Ti)C NbC
AN31
1.53
–
1.07
–
0.162
Fe2Mo NbC
Esshete 1250
1.04
–
0.96
–
0.197
12R72
1.16
–
–
0.30
0.093
Fe2Mo
TiC M23C6
c = critical) values, and schematic microstructures. The creep rupture properties of five 15%Cr–15%Ni system steels are presented in Fig. 2.27.92 Among these alloys 17-14CuMo exhibits the highest creep rupture strength for the longest time and at the highest temperature. In Fig. 2.28, the allowable stresses for 18%Cr–8%Ni steels and for 17-14CuMo and Esshete1250 representing 15%Cr–15%Ni steels are compared. The allowable stress of Super304H is much higher than that of 17-14CuMo, which is thought to have the highest stress at temperatures up to 670°C among conventional materials. TP347HFG and Tempaloy A-1 also have allowable stresses higher than those of existing conventional steels. Because these steels have been developed on the basis of TP304H or TP321H, their cost effectiveness is excellent. They are also advantageous from the standpoint of resistance to steam oxidation owing to their fine grains. Although newly developed 20%–25%Cr austenitic steels and high Cr– high Ni austenitic steels such as CR30A93 and HR6W94 have excellent resistance to high temperature corrosion and steam oxidation compared to other austenitic steels, their drawback lies in the fact that they are too costly for their allowable stresses. However, as shown in Table 2.5, the materials developed in the 1980s, particularly N-alloyed 20%–25%Cr steel, have excellent high temperature strength as well as being relatively inexpensive. They are practically applied as high strength steels taking high temperature corrosion resistances into account. Allowable stresses for HR3C,95
WPNL2204
The development of creep-resistant steels 700°C
300
Stress (MPa)
53
200
100
17–14CuMo 15–15N AN31 Esshete 1250 12R72
70 102
103 Time to rupture (h)
104
700°C
300
Stress (MPa)
105
200
100 70
102
103 Time to rupture (h)
104
105
2.27 Comparison of creep rupture strength of 15Cr–15Ni austenitic steels.
NF70996 and Tempaloy A-397 are far higher than that of Alloy 800H, as shown in Fig. 2.29 and they can be used in higher steam conditions and in corrosive environments. As an example, Fig. 2.3095 shows the effect of the content of solute N on HR3C, with the changing Nb and Ni content in the alloy. Nitrogen as an alloying element acts as a creep strengthening agent in 25%Cr austenitic steel and the combination of 0.4%Nb and 20%Ni addition is an effective measure of the enhancement of creep strength. Alloy 800H has a stable austenite structure, stabilised by using a large addition of Ni, but high temperature strength is insufficient in relation to cost. Although there is currently no choice but to use austenitic steels for superheater and reheater tubes for ultra-supercritical pressure boilers, certain materials have already been developed that are sufficient to meet the steam conditions of 650°C class boiler superheater and reheater tubes, as noted previously. From the mid-1990s to date, materials taking cost effectiveness into consideration have also been developed. SAVE2598 and Sanicro2599 are
WPNL2204
54
Creep-resistant steels
150
TP
34
100 TP
31
FG 17
Te m 30
4H
o Te m
pa
lo
y
A-
1
1H
600 Temperature (°C)
pa
lo
y
AA
-1
0
550
uM
25
0 500
4C
E1
32
50
–1
6H
TP TP
4
7H
70
7H
A
TP34
X
Allowable stress (MPa)
Super 304H
650
700
2.28 Allowable stresses for 18Cr–8Ni and 15Cr–15Ni austenitic steels.
examples that use 0.25% addition of N to stabilise the austenitic structure based on HR3C, in addition to a small amount of Nb addition aimed at precipitation strengthening by means of ‘under-stabilising’. Furthermore, comprehensive strengthening techniques covering a wide range of temperatures have been employed by introducing the concept of Cu addition in Super304H and W addition in HR6W. Here, cost effectiveness has been secured by stabilising the austenitic structure through addition of N and Cu, while decreasing the Ni addition to 18% and reducing the Cr content to a level slightly below that of HR3C. As shown in Fig. 2.25, austenitic steels that have the highest creep rupture strength at 700°C are HR6W, SAVE25 and Sanicro25, exhibiting around 90 MPa at 100 000 h. Figures 2.3194 and 2.32100 show creep rupture strength in the temperature range 650–800°C. Stable creep rupture properties are confirmed in these figures. From the standpoint of realising 700°C class ultra-super critical pressure power plants, achievement of 100 MPa of 100 000 h creep rupture strength (70 MPa in allowable stress) is considered to be required, so that the limits for application of 18%Cr–8%Ni steels and 20–25%Cr austenitic steels with
WPNL2204
The development of creep-resistant steels
55
NF709 HR
150
loy A -3
3C VE SA
Temp a
25 R6 W
Allowable stress (MPa)
H
Alloy 8 00H
100
CR
TP
50
0 500
550
31
30
A
0H
600 Temperature (°C)
650
700
2.29 Allowable stresses for 20–25Cr and high Cr–high Ni austenitic steels. 200
103 h Creep rupture strength (MPa)
180 160 140 120 100 80 700°C
60
750°C
Nb
Ni 20/23
40
17 0.4/0.5
20 0 0.05
20 23
0.10
0.15 0.20 Soluble N (%)
0.25
0.30
2.30 Effect of content of solute N and Nb and Ni on creep rupture strength of HR3C.
WPNL2204
56
Creep-resistant steels 400
650°C 700°C 750°C 800°C
300
Stress (MPa)
200
100
70 50 40 101
HR6W: 0.07C-23Cr-43Ni-6W-0.1Ti-0.2Nb
102
103 Time to rupture (h)
104
105
2.31 Creep rupture properties of HR6W.
500
650°C 700°C 750°C 800°C
Stress (MPa)
400 300
200 70 50 100
101
102 103 Time to rupture (h)
104
105
2.32 Creep rupture properties of SAVE25 steel.
the highest creep rupture strength in 350 bar plants would be approximately 660°C and 680°C, respectively, as shown in Fig. 2.33.91 In the accompanying figure, allowable stresses at higher temperatures for numerous austenitic steels developed to date are presented together with those for Ni-based alloys. However the allowable stresses in Ni-based alloys are not yet well established owing to insufficient databases and inaccuracy of predicted strength at 100 000 h. Reviewing the austenitic steels developed beyond the strength of conventional steels to date, materials can be classified in the strength levels as shown in Table 2.7 at a temperature of 700°C. Austenitic steels are graded
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The development of creep-resistant steels
57
Table 2.7 Strength classification of austenitic steels 105 h creep rupture strength at 700°C (MPa)
Allowable stress (MPa)
Steels
Category
60
40
65
43
70
50
75 80, 85 90
53 55 60
TP347HFG Tempaloy A-1 HR3C NF709 Tampaloy A-3 Super304H XA704 Tempaloy AA-1 CR30A – HR6W SAVE25 Sanicro25
18Cr 18Cr 20/25Cr 20/25Cr 20/25Cr 18Cr 18Cr 18Cr High Cr – High Cr 20/25Cr 20/25Cr
in five levels except for the 80 MPa and 85 MPa classes in which strengths are separated in increments of 5 MPa. Table 2.891 lists austenitic steels and Ni-based alloys whose maximum application temperature is categorised by 620–660°C, 620–680°C and 680–770°C. The table also shows the ASME Code Case number assigned and applicable product forms, tubes and/or pipes. Austenitic steels for thick-section pipe and turbine component applications Thick section components in boilers consist of headers and steam piping and under conventional steam conditions, CrMo low alloy steel (for example, 2.25%Cr–1%Mo steel, 0.5%Cr–0.5%MoV steel or martensitic 12%Cr stainless steel X20CrMoV121)101 has been used extensively up to 566°C for many years. From 1989 onwards, P91102 has been selected for headers and piping operating up to 600°C and from the mid-1990s P92103 has also been employed for recently constructed 600°C class ultra-super critical power plants. However, under steam conditions above the temperature at which ferritic steels can be applied, generally above 620°C or 630°C, austenitic steels must presently be used from the standpoint of creep strength and oxidation resistance. Among austenitic steels, ASTM TP316 has actual application experience for the header and main steam piping in the Eddystone Station unit No. 1 in the USA under steam conditions of 350 bar and 650°C. Being an austenitic steel, this material has a low thermal conductivity meaning that thermal stresses will be higher than for ferritic steels. Further, as the high temperature properties of TP316 are greatly influenced by its chemical composition, its high temperature strength will be highly unstable if the chemical composition
WPNL2204
58
Trade designation
Nominal composition
ASME Code/ Code Case
Preferred application
Temperature of application (metal)/°C
Austenitic steels (18%Cr)
347HFG Super304H XA704 Temp. AA-1
18Cr–10Ni–Nb 18Cr–8Ni–Cu–Nb–N 18Cr–9Ni–W–V–Nb 18Cr–10Ni–Nb–Ti
2159 2328 2475 2512
T T T T
620–660
Austenitic steels
HR3C NF709 Temp. A-3 800HT SAVE25 Sanicro25 HR120 HR6W
25Cr–20Ni–Nb–N 20Cr–25Ni–Nb–Ti–N 22Cr–15Ni–N6–N 21Cr–32Ni–Al–Ti 23Cr–18Ni–W–Cu–Nb–N 23Cr–18.5Ni–W–Cu–Nb–N 25Cr–42Ni–N 23Cr–43Ni–6W–Nb–Ti
2115 2581 – 1325 – – – –
T T T T T T T P.T
620–680
Haynes230 Inco617 Inco625 Inco740 45TM
22Cr–5Co–3Fe–14W–2Mo–La 22Cr–13Co–9Mo–Al–Ti 22Cr–5Fe–9Mo–Nb–Al–Ti 25Cr–20Co–Mo–Nb–Al–Ti 27Cr–23Fe–2.75Si
2063 1956 1409 – 2188
P.T P.T P.T P.T P.T
680–770
(20–25%Cr)
Ni-base allows
T: superheater/reheater tubes, P: pipes and headers
WPNL2204
Creep-resistant steels
Table 2.8 Candidate alloys for USC power boilers
The development of creep-resistant steels
59
120
one
Inc
7 l 61
on el 74 0
E25
A-3 loy 0H 80 y
40
TP
31
63 c2
50
oni
60
Nim
lo
0S X
Su
Allowable stress (MPa)
Inc
SAV
pa
Al
70
Tem
80
30
25
90
es2
icro
100
Hyn
San
110
A 4
r3
70
pe 04
30
H
20
HR
HR 3C
NF7
10 0 600
Hast
650
700
750
09
elloy
800 850 Temperature (°C)
6W
XR 900
CR30
A Hast
elloy
950
X
1000
2.33 Comparison of maximum allowable stresses at high temperatures.
is not carefully optimised. This must be given special consideration because TP316 is very suitable for use in the heavy sectioned components operating at steam temperatures above the maximum limit for ferritic steels. In the case of TP316, the formation of the sigma phase can be suppressed in theory by having the Ni-valance (= 11.6 + Nieq – 1.35Creq, where Nieq = 0.5Mn + 30(C+N) + Ni, Creq = 1.5Si + Cr + Mo + 0.5 Nb) above 0.104 In practice, however, the sigma phase may precipitate after a long term of service causing a major drop in the creep rupture strength. Figure 2.34105 shows creep rupture test results for the main steam pipes produced at the end of the 1950s and used in the Eddystone No. 1 unit over a period of 130 250 h. It can be seen that a low Ni balance greatly reduces the creep rupture strength. The materials indicated as 17U and 17D in the graph were produced at the start of the 1980s using modern steel-making techniques to keep Ni balance high and were used for only 6000 h under the same conditions, with other samples showing high strength. Among the steels listed in Table 2.5, HR6W is being considered for thick section components of 700°C class ultra supercritical pressure power boilers, taking advantage of its very high creep rupture strength which is on a par with Ni-based alloys.
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60
Creep-resistant steels
400 Sample (Ni-bal.) 25U (2.134) 25D (2.126)
300
20D (–0.370) 20U (–0.367) 17D (1.536) 17U (1.339)
18D (1.195) 18U (2.306) 19D (0.425) 19U (0.045)
8D (–0.134) 8U (–1.167) 7D (–0.017) 7U (–0.961)
Stress (MPa)
200
100 NIMS data (ave.) 650°C
70
50 102
103 Time to rupture (h)
104
105
2.34 Creep rupture properties at 650°C for as-exposed TP316 steel.
A typical thick section component in a steam turbine is the rotor. For quite some time, 1%CrMoV steel and conventional 12%Cr steel have been used for steam temperatures of 566°C and below. The maximum use temperatures for these steels are considered to be 570°C and 590°C, respectively. 12%Cr steel, for which 1.5% W is added in substitution for the equivalent amount of Mo, can be used up to approximately 600°C, but higher temperatures require 12%Cr steel with the addition of W and Co. While the limit has not been clearly defined, safety considerations suggest a temperature of around 630°C. The arrival of ferritic steel that can withstand 650°C is awaited, but advanced materials development is needed since rotors require not only high-temperature strength but also good tensile strength and toughness. Austenitic steels such as Discaloy (13%Cr–25%Ni–3%MoAlTi) and A286 (15%Cr–26%Ni–1.5%MoAlTi) have demonstrated performance in Eddystone No. 1 unit at 650°C and 350 bar, and in a 50 MW demonstration plant in Wakamatsu, Japan also at 650°C. But temperatures over 700°C would require Ni-based alloys such as Inco617 (Ni–22%Cr–13%Co–9%MoAlTi) or Waspaloy (Ni–19%Cr–14%Co–4.5%MoAlTi), or Fe–Ni-based alloys with improved creep strength. From the perspective of practical application, fabricability would need to be improved to enable the manufacture of large components. Figure 2.35 presents the relationship between 100 000 h creep rupture strength and temperature for various materials based on the results of research to date. Along with consideration of Inco617, efforts are also being made worldwide to improve Inco706 (Ni–16%Cr–36%Fe–3%NbAlTi) and to develop alloys modifying the low thermal expansion Inver-type Ni alloy. Of these, LTES700 (Ni–12%Cr–18%MoAlTi)106 is an austenitic alloy featuring a low
WPNL2204
The development of creep-resistant steels
61
105 h Creep rupture strength (MPa)
700 Refractaloy 26
500
300
Waspaloy
Inco617
400
12CrMoW 12CrWCo 12CrMoV
Inco706 LTES
Mod. A286 200 1CrMoV
100 70 450
500
550
600 Temperature (°C)
650
700
750
2.35 Comparison of 105 h creep rupture strength of turbine rotor alloys.
thermal expansion coefficient similar to ferritic steels and having strength equivalent to that of Refractaloy26 (Fe–18%Cr–38%Ni–20%Co–3%MoAlTi). Furthermore FENIX-700 (Ni–16%Cr–36%Fe–2%NbAlTi), 107 modified Inco706 and TOS-1X (Ni–23%Cr–10%Mo–15%CoAlTiB)108 have been developed aiming at 700°C class steam turbine rotor application. For temperatures above around 630°C, austenitic steels must be used because ferritic steels are presently not strong enough. The Discaloy used in Eddystone No. 1 unit weighed 1.6 tons and was made by melting at atmospheric pressure. However, rotors in 1000 MW class ultra supercritical pressure power plants will require a steel ingot weighing nearly 30 tons, and although it is possible to cast ingots of over 800 mm in diameter by the electro slag remelting (ESR) melting process, there will be a greater tendency for freckling to occur and this may impair the mechanical properties.109 Tests on a 1000 mm diameter A286 steel ingot with a rationalised chemical composition cast using large-size ESR equipment have shown that even if freckling occurs, it has no effect on high cycle fatigue or fracture toughness at room temperature, or on creep rupture strength, high or low cycle fatigue at high temperatures, although it may reduce tensile elongation at temperatures below 200°C.110 Ti and C contents influence freckling to the greatest extent and it has been reported that there was no freckling at all in a 1350 mm diameter ESR A286 steel ingot in which the Ti and C contents had been reduced to 1.13% and 0.02%, respectively. Although reducing the Ti content lowered the creep rupture strength of an A286 ingot, as shown in Fig. 2.36,110 for the level of reduction mentioned above, it satisfies the tensile property and creep rupture
WPNL2204
62
Creep-resistant steels
105 h Creep rupture strength at 650°C
160
100 kg ingot 2 ton ingot 16–19 ton ingot
150
140
130
120
110 0.9
1.0
1.2
1.4 Ti (mass%)
1.6
1.8
1.9
2.36 Effect of Ti content on 105 h creep rupture strength of A286 alloy at 650°C.
strength requirements for rotor material to be used at 650°C, with no reduction in creep notch strength. The Ti and C contents in A286 can also be reduced to 1.52–1.81% and 0.023–0.039%, respectively with no loss in toughness at 650°C confirming in theory the excellent suitability of the chemically improved A286 as a 650°C class rotor material. Austenitic steels for chemical plant applications The typical heat-resistant austenitic steels in chemical plants are used for reformer tubes in hydrogen treating equipment and for cracking tubes in ethylene polymerisation equipment in thermal cracking of naphtha at temperatures above about 800°C. In the case of high pressure equipment such as reformer tubes, higher creep rupture strength is required and anticarbonisation is also required for cracking tubes which are used at higher temperatures. Table 2.9111 shows the chemical composition of these austenitic tube materials for chemical plants made by centrifugal casting. The base alloy is HK40, 25%Cr–20%Ni with a high carbon content of 0.4%. The creep rupture strength has been improved through investigation into the effect of chemical elements. Nb and Ti are used to form eutectic carbide which strongly contributes to improved creep rupture strength, and W and Mo are used as solution strengthening elements in the austenitic matrix. In order to prevent the precipitation of the sigma phase and to stabilise the austenitic matrix as well as to improve anti-carbonisation, the Ni content is increased from 20% to 30–50%, and Si is added (also for the anti-carbonisation).
WPNL2204
Table 2.9 Nominal chemical composition of austenitic steels for chemical plants Steels
Cr–Ni– based
Mo addition
Co addition
C
Cr
Ni
Nb
Co
W
Mo
Ti
Others
0.40 0.43 0.45 0.50 0.50 0.50 0.40 0.30 0.50 0.50 0.40 0.45 0.50 0.50 0.15 0.40 0.45 0.50 0.40 0.45 0.40 0.50 0.50
25.0 23.5 25.0 25.0 25.0 27.0 25.0 24.0 25.0 25.0 25.0 25.0 24.0 25.0 22.0 25.0 25.0 28.0 33.0 25.0 28.0 22.0 25.0
20.0 22.5 20.0 35.0 35.0 29.0 20.0 24.0 35.0 35.0 20.0 20.0 23.0 35.0 30.0 35.0 35.0 48.0 50.0 45.0 48.0 30.0 35.0
– – – – – – – 1.5 1.5 1.0 0.2 0.6 0.7 0.7 1.2 – – – – – – – –
– – – – – – – – – – – – – – – – – – – 3.0 5.0 13.0 15.0
– – – – – – – – – 1.5 – – – – – – 2.0 5.0 17.0 3.0 5.0 4.5 5.0
– – – – – – – – – – – – – – – 1.2 – – – 3.0 – – –
– – – – – – 0.20 – – – 0.20 0.15 0.15 0.15 – – – – – – – – –
– – Si:2.0 – Si:2.0 – R.E.:0.3 – – Si:2.0 R.E. :0.3 – – – – – – – – – – – –
The development of creep-resistant steels
Ti addition Nb addition
HK-40 Mod. HK40 High Si-HK40 HP High Si-HP HO HiKa-1A IN-519 Manaurite 36 Manaurite 36XS Hika-1B BST Mod. BST HP-BST Manaurite 900 HOM MoRe-1 NA-22H MoRe-2 HOM-3 Super 22H Tenex Supertherm
Chemical composition (mass%)
63
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High temperature oxidation and corrosion resistance are improved by the addition of rare earth metal material. Figure 2.37112 demonstrates the creep rupture strength of these centrifugally cast austenitic steel tubes. The creep rupture strength of HK40 is lowest, while the other steels are substantially improved by alloy development. HP-BST whose chemistry is a nominal 0.5%C–25%Cr–35%NiNbTi exhibits the highest creep rupture strength. In the case of centrifugally cast tubes, small diameter sizes are difficult to produce and the extrusion process is also applied to produce tubes. However, the eutectic structure is decomposed by the extrusion process, resulting in a weakening of creep rupture strength. In this case, creep strength improvement by heat treatment is used in reforming the eutectic type carbide structures in the austenitic matrix.
2.5
Historical development of steel melting and of the purity of heat-resistant steels
The increasing demand of the steam plant manufacturer for turbine rotor forgings with improved cleanliness, soundness, higher ductility and toughness has mainly led to improvement in steelmaking processes for heat-resistant steels within the last century. Figure 2.38 gives a rough overview of the historical development of steelmaking processes throughout the 20th century and the achievable reduction in the phosphorus and sulphur contents as the main indicators of the effectiveness of these steelmaking processes.15,113–117 At the beginning of the last century only steels of the Thomas and open
50 40
1000°C Mod. BST
Stress (MPa)
30 HP-BST 20
10 IN519
102
103 104 Time to rupture (h)
Super them NA22H MoRe-1 Manaurite 36X HP HK40
105
2.37 Creep rupture strength of centrifugally casted tubes for chemical plants.
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(1) S < 0.075%/P < 0.10% (*) Oxidizing refining in EF
Electric-arc steel EF Thomas steel (1877) Basic open hearth steel (1964) BOH
1900
1920
(1) Max. P and S (%)
reladle
1940
1960 Year of application
1980
2000
P,S
0.02 P
P,S
P
P,S 0.01 P* P* Thomas steel
BOH
EF BOF Steel melting processes
S EF + ESR
S EF + AOD
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2.38 Historical development of steel melting processess and their achievable P and S contents.
S
The development of creep-resistant steels
Ladle refining furnace LRF Vacuum induction melting VIM Vacuum arc remelting VAR Vacuum oxygen decarburization VOD Argon oxygen decarburization AOD Electric slag remelting ESR Vacuum carbon desoxydation VCD Vacuum degassing VAD (H2 <2ppm) Oxygen blown steel BOF
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hearth processes, which possess high P and S contents were available. The electric-arc furnace steel, first introduced in 1908, provided very cost effective and much lower levels of P and S and better overall cleanliness but hydrogen flaking was still a problem of larger forgings. This problem was solved by the introduction of vacuum degassing introduced in the 1950s. The production of oxygen blown steels started in 1952. This basic oxygen steelmaking converter process, which works with a high amount of pig iron (nearly 70%), offers a very economic way to produce clean steels with lower P and S contents and with lower tramp elements, such as As, Sb and Sn, as long as cleaner pig iron is used. A further improvement in vacuum treatment was established in 1961 with the adoption of the vacuum carbon deoxidisation (VCD) process. VCD allows for the removal of oxygen through reaction with carbon to form CO as a gas which is then removed from the pouring stream by a vacuum system. This provided cleaner steels than the earlier practice in which oxygen was tied-up with solid deoxidants (Si, Al) which resulted in undesirable silicate and Al-oxide inclusions. Demands for cleaner and more uniform chemical homogenity were provided by the secondary remelting methods: electroslag-remelting (ESR) in the middle of the 1960s and vacuum-arc remelting (VAR) at the end of the 1960s. The advantage of the ESR process is also a low S content of about 0.002%. The VAR process provides no reduction of the trace elements but very low gas contents. Very effective secondary refining steelmaking processes are also the decarburisation technologies Argon-oxygen decarburisation (AOD) and vacuum-oxgen decarburisation (VOD) were introduced at the end of the 1960s. The advantages of the AOD process are lower S contents (about 0.002%) and lower N2 and H2 levels. Vacuum-induction melting (VIM) also provides very low gas contents. For this process very clean iron and alloy ores, selected scraps and very clean pig iron are used. The risk of shrinkages in the VIM ingot requires a remelting (ESR or VAR) after the VIM process. VIM+ESR is used, for example, for the Ni-alloys Inconel 617 and Inconel 625, but also for special steels, like 17-4PH. For the very high stressed Inconel 706 gas turbine discs, a triple melting of VIM+ESR+VAR is common. In the middle of the 1970s, very large forgings for nuclear and fossil fired plants with higher output demanded larger ingots. This challenge led to the development of ladle refining furnaces and subsequently commited the electricarc furnace only to the role of melting. With the developed ladle refining furnaces, various combinations of heating, degassing, desulphurising, dephosphorising and methods for inclusion shape control are possible in order to provide large forging with high quality.114 This ladle metallurgy allows delivery P and S contents in the range of about 0.002%. Low P and low Mn contents are also often made by oxidising refining processes in the electric furnace (see for example Azuma et al.)118 In addition to phosphorus, antimony, arsenic, tin and copper are elements which lead to embrittlement
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in service and to lower creep strength in the heat-resistant steels. Each of these elements is non-oxidisable and therefore controlled through scrap selection for the electric-arc furnace melting process in order to avoid steel degradation. A further weakness in steelmaking practice in the past was the deoxidisation of the steels with aluminium. In consequence of the discovered negative influence of Al on creep strength and creep ductility in long term investigations, Al is nowadays limited to levels lower than 0.010%.
2.6
Summary
The development of creep-resistant steels is a result of continuous technological progress throughout the 20th century. The urgent need to improve the creep strength of steels was based on endeavours by the power station industry to improve the thermal efficiency of steam power plant by raising the steam temperature and steam pressure in order to reduce the cost of fuel and go easy on fuel resources. The major contribution to the increase in power plant efficiency consisted in the development of heat-resistant steels with a higher creep strength at an acceptable creep ductility level. The significance of these material properties was not recognised until early damage was suffered by steam turbine bolts in the 1930s, which pointed to the fact that the strength of steels used in power stations operating at higher temperatures depends significantly on the creep behaviour of the material over the full period of operation. Based on this experience, it was concluded that the strength values should no longer be determined in short-term tests but that a procedure should be adopted to determine the rupture strength, the creep strain and creep ductility of the heat-resistant steel in a creep test extending over a period of roughly 100 000 h. Based on the multiplicity of investigations of test steels carried out with different Mo, Cr, Ni, V, CrMo, CrV, MnSi, MoMnSi, CrSiMo, CrNiMo, CrMnV and CrMoV contents, worldwide developments in the manufacture of steam boilers and small forgings for steam turbines brought forth low alloyed steels with chemical compositions of 0.15%C–0.3% to 0.5%Mo, 0.13%C–1%Cr–0.5%Mo and 0.10%C–2.25%Cr–1%Mo which are still in use today. In addition, in the 1950s, a MoV steel with 0.14%C–0.5%Mo– 0.3%V with an even higher creep strength was developed in Europe for gas turbines and afterwards also qualified in long term creep tests for steam plants. In the field of turbine manufacture, a steel with approximately 0.25%C– 1.25Cr–1%Mo–0.30%V is in worldwide use for turbine rotors, casings, bolts and small forgings. The development of heat-resistant 9–12%Cr steels was strongly motivated by two major events: during the 1950s by the development of thermal power stations for public power supply operating at steam temperatures ranging from 538–566°C and during the 1980s by the target to develop low-pollution
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power stations operating at steam admission temperatures of 600–650°C with supercritical pressures up to 350 bar. The steel X22CrMoV 12 1 was developed in the 1950s for thin-walled and thick-walled power station components. Its creep strength is based on solution hardening and on the precipitation of M23C6 carbides. The steel has been applied successfully in power stations over several decades up to temperatures of about 566°C. The steel referred to in the literature as mod. 9Cr1Mo or P91, is a steel of the newer generation. It was developed under a huge USA project in the late 1970s for the manufacture of pipes and vessels for a fast breeder. This steel has meanwhile found wide application in all new Japanese and European power stations with steam temperatures up to 600°C for pipes and small forgings. It is also used for the manufacture of castings, like valve chests and turbine casings. The increase of creep strength in comparison to the 12%CrMoV steel is caused by 0.05%Nb and 0.05%N which form thermal stable VN and Nb(C, N) precipitates. A lower Cr content of about 9% also contributes to the higher creep strength. Similar creep strengths are exhibited by new steels developed for rotors, castings and pipes which are in addition alloyed with 1%W. A further increase of about 10% in creep strength reveals the steel type P92 which is in addition alloyed with about 0.003%B and has an increased W content of about 1.8%. Addition of boron gives thermally stabile M23(C,B)6 precipitates whereas the higher W content leads to a higher amount of the Laves phase. Higher B contents in the range of 0.010–0.014% give steels for rotors, casings and pipes which can be used up to temperatures of about 630°C. Further ferritic 9–10%Cr steels are under development for steam temperatures up to 650°C. A great challenge for the development of further improved 9–10%Cr heat-resistant steels is, first, the avoidance of the Z-phase precipitation Cr2(Nb,V)2N2 which precipitates at the expense of the MX particles, and second, the avoidance of BN and metallic borides which precipitate at the expense of M23(C,B)6 and VN. An important contribution to the development of advanced steam power plants was made in the last century with a continuous improvement in the melting technology which improved significantly the purity of the steels and the manufacturability of large heat-resistant steel components. Austenitic steel itself has originally been developed for chemical plant equipments used in various corrosion and oxidation environments. Economically to keep the structure austenitic, Cr and Ni contents were optimised at 18% and 8%, respectively (Type 304 steel) in the 1920s; later completely stabilised austenitic steels such as Types 321 and 347 steels were developed. Type 316 steel was also developed containing 3% Mo for use in ammonium chloride and sulphuric acid environments in 1930s. At the same time, the addition of combination of 2%Mo and 2%Cu to 18%Cr–8%Ni steel was found to provide excellent corrosion resistance against sulphuric acid.
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Fortunately the austenitic steels exhibit not only good corrosion resistance but also very high creep rupture strength, such as above 100 MPa of 100 000 h strength at 600°C for 18%Cr–8Ni% steels. The steam temperature of fossil power plant in the USA and Europe rose to around 538–566°C from the end of the 1940s to the beginning of the 1950s. The same situation reached Japan at the end of the 1950s. This meant that low alloyed Cr–Mo steels could no longer be used in boilers in superheater and reheater tubes whose metal temperature exceeded 580°C, from the view point of oxidation and corrosion as well as creep strength behaviour. Austenitic steels or high Cr ferritic steels needed to be used. Four types of 18%Cr– 8%Ni system austenitic steels could be used for such high temperature steam plants. Also, ultra-supercritical pressure power plants constructed in the latter half of the 1950s were made possible by further application of these austenitic steels to thick walled components. For example, TP316H was used for boiler headers and steam piping, 17%Cr–14%NiCuMoNbTi (categorised in the 15%Cr–15%Ni system austenitic steel), TP321H for superheater and reheater and Discaloy (13%Cr–25%Ni–3%MoAlTi) for a very high pressure turbine rotor at the Eddystone unit No. 1 of Philadelphia Electric operated at a temperature of 649°C in the 1960s. After 1980, numerous studies on the development of austenitic steels for use in power plants operating at temperatures above 650°C were started aiming at the enhancement of creep strength rather than using conventional austenitic steels originally developed for chemical applications. The alloy design concepts for austenitic steels were (1) enhanced precipitation strengthening by means of ‘under-stabilising’ C composition, (2) lowering the material cost by addition of N to reduce Ni content to maintain a full austenitic structure of high Cr above 20% containing steel, (3) utilisation of Cu and W for precipitation hardening and solution strengthening, and (4) precipitation hardening by inter-metallic compounds. Presently, the greatest 100 000 h creep rupture strengths at 700°C available are 70 MPa for 18%Cr– 8%Ni system steels and 90 MPa for 25%Cr austenitic steels. As 100 MPa is thought to be required for high-temperature and high-pressure power plants, the limits for application of 18%Cr–8%Ni system steels and 25%Cr austenitic steels in 350 bar pressure plants would be approximately 660°C and 680°C, respectively. For thick wall components experience of Type 316 steel is already available at the 650°C steam temperature. However, for temperatures above 650°C, development of new austenitic steels or improvement of Nibased alloy is needed. For steam turbine rotor applications, austenitic super alloys such as Dicaloy (13%Cr–25%Ni–3%MoAlTi) and A286 (15%Cr– 26%Ni–1.5%MoAlTi) have demonstrated a restricted performance at the 650°C level, but temperatures of 700°C and greater would require Ni-based alloys such as Inco617 (Ni–22%Cr–13%Co–9%MoAlTi) or Waspaloy (Ni19%Cr–14%Co–4.5%MoAlTi), or Fe–Ni based alloys with improved creep
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strength. The limit for austenitic steels is approximately 680°C for 350 bar pressure plants. On the other hand, Ni-based alloys can be applied at temperatures in excess of 680°C or 700°C, but the material cost is quite high compared with austenitic steels. In future, continued development is needed for austenitic steels that can be used at temperatures above 700°C aiming at the 100 000 h creep rupture strength of 100 MPa and greater at 700°C.
2.7
References
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61 Kauffmann F, Mayer K H, Straub S, Zies G, Scheu C, Willer D, Ruoff H and Maile K, ‘Characterization of the precipitates in modern boron containing 9–12%Cr steels developed in the frame of the COST program’, 8th Liège Conference on Materials for Advanced Power Engineering, 18–20 September, Liège, Belgium, Forschungszentrum Jülich, Germany, 2006. 62 Kager F, Böck N, Spiradek-Hahn K, Höfinger S, Brabetz M and Zeiler G, ‘Superior long-term creep behaviour and microstructure evolution of 9%Cr-steels with boron’, 8th Liège Conference on Materials for Advanced Power Engineering, 18–20 September, Liège, Belgium, Forschungszentrum Jülich, Germany, 2006. 63 Miki K, Azuma T and Ishiguro T, ‘Improvement of long term creep strength in high Cr heat resistant steel’, 15th International Forgemasters Meeting IFM 2003, Kobe, Japan, October 26–29, 2003, Steel Castings and Forgings Association of Japan, 269–275. 64 Scarlin B and Stamatelopoulos G N, ‘New boiler materials for advanced steam conditions’, Proceedings of the 7th Liège COST Conference on Materials for Advanced Power Engineering 2002, Part III, Liège Belgium, Forschungszentrum Jülich, Germany, 2002, 1091–1108. 65 Bendick W, Vaillant J C, Vandenberghe B and Lefebre Bo, ‘VM12 – a new 12%Cr steel for boiler tubes, headers and steam pipes in USC power plants’, ‘Development of a new 12%Cr-steel for tubes and pipes in power plants up to 650°C’, EPRI International Conference on Materials and Corrosion Experiences for Fossil Power Plants, 18–21 November, Wild Dunes Resort, Isle of Palms, SC, USA, Hilton Oceanfront Resort, Hilton Head Island, SC, USA, October 25–28, 2004, Electric Power Research Institute Palo Alto CA, 2003. 66 Masuyama F, ‘Steam plant material development in Japan’, 6th Liège COST conference, Materials for advanced power engineering 1998, Conference proceedings III, Liège, Belgium, September 1998, Forschungszentrum Jülich GmbH, Germany, 1807. 67 Bendick W, ‘Neue Werkstoffentwicklung für moderne Hochleistungskraftwerke’, VGB Tagung Werkstoffe und Qualitätssicherung 2004, 10–11 März, 2004. 68 Blum R, Vanstone R W and Messlier-Gouze C, ‘Materials development for boilers and steam turbines operating at 700°C’, 4th EPRI International Conference on Advanced in Material Technology for Fossil Power Plants, Hilton Oceanfront Resort, Hilton Head Island, SC, USA, October 25–28, 2004, Electric Power Research Institute Palo Alto CA, 118–138. 69 Masuyama F in Landolt–Boernstein, Numerical Data and Functional Relationships in Science and Technology, Group VIII, Vol. 2 Materials Subvol. B, Creep Properties of Heat Resistant Steels and Superalloys, ed by Yagi K, Merckling G, Kern T U, Irie H, Warlimont H, Springer Verlag, Berlin Germany 2004. 70 Danielsen H and Hald J, Z-phase in 9-12%Cr Steels, VÄRMEFORSK Service AB, Stockholm, Sweden, April 2004. 71 Sato A, ‘Research project on innovative steels in Japan’, International workshop on Innovative Structural Materials for Infrastructure in 21th Century, Tsukuba, NIMS Tsukuba, Japan, 12–13 January 2000. 72 Abe F, Igarashi M, Fujitsuna N and Muneki S, ‘Alloy design of advanced ferritic steels for 650°C USC boilers’, Conference on Advanced Heat Resistant Steels for Power Generation, 27–29 April San Sebastian, Spain, Electric Power Research Institute Palo Alto CA, 1998. 73 Semba H and Abe F, ‘Development of creep resistant 9Cr-3W-3Co steel containing
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75 76 77 78 79 80 81
82 83 84
85
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87
88
89 90
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high boron for USC boilers’, First International Conference Super-High Strength Steels, 2–4 November, Centro Sviluppo Metallurgia, Rome, Italy, 2005. Abe F, ‘High performance creep resistant steels for the 21st century power plants’, First International Conference Super-High Strength Steels, 2–4 November, Centro Sviluppo Metallurgia, Rome, Italy, 2005. Krainer H, ‘50 Jahre nichtrostender Stahl’, Stahl und Eisen, 1962, 82, 1527. Strauses B, ‘Non-rusting Cr-Ni steels’, Proceedings American Society Test Materials, 1924, 24, 208. Schottky H, ‘German version of birth of stainless steels’, Iron Age, 1929, December 5, 1512. Mitchel W M, ‘18-8 and related stainless steels’, Metals and Alloys, 1940, January, 14. Houdremont E and Schafmeister P, ‘Verhütung der Korrosion bei Stählen mit 18%Cr und 8%Ni’, Archiv. Eisenhuettenwesen, 1933, 7, 187. Suzuki T, Historical Invent of Stainless Steels (in Japanese), AGNE Technical Center, Tokyo, 2000, 110. Eberle F, Ely E G and Dillon J S, ‘Experimental superheater for steam at 2000Psi and 1250F – Progress report after 12 000 h of Operation’, Trans. ASME, 1956, 76, 665. Masuyama F, Materials for Advanced Power Engineering 1998, Part III, J. LecomteBeckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 1998, 1807. Murray J D, ‘Welding Trails on Esshete 1250’, Welding and Metal Fabrication, 1962, 9, 350. Minami Y, Kimura K and Tanimura M, ‘Creep rupture properties of 18%Cr-8%NiTi-Nb and Type 347H Austenitic Stainless Steels’, ASM International Conference New Development in Stainless Steel Technology, Detroit, Michigan, September 17– 21, ASM International, Materials Park, Ohio, 1984. Yoshikawa K, Fujikawa H, Teranishi H, Yuzawa H and Kubota M, ‘Fabrication and progress of corrosion-resistant TP347H stainless steel’, ASM Conference Coatings and Bimetallic for Aggressive Environments, ed. by R. D. Sission Jr.1985 99, ASM International, Materials Park, Ohio 44073 1984. Sawaragi Y, Ogawa K, Kan S, Natori A and Hirano S, ‘Development of the economical 18-8 Stainless Steel (SUPER 304) having elevated Temperature Strength for Fossil Power Boilers’, Sumitomo Search, 1992, 48, 50. Ishituka T and Mimura H, ‘Development of New 18Cr-9Ni austenitic stainless steel boiler tube’, 7th Liège COST Conference Materials for Advanced Power Engineering 2002, September 2002, Liège Belgium, J. Lecomte-Beckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 2002, 1321. Minami Y, Tohyama A and Hayakawa H, ‘Properties and experiences with a new austenitic stainless steel (Tempaloy AA-1) for Boiler Tube Application’, 7th Liège COST Conference Materials for Advanced Power Engineering 2002, September 2002, Liège Belgium, J. Lecomte-Beckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 2002, 1445. Masuyama F, ‘History of power plants and progress in heat resistant steels’, Transactions of Iron Steel Institute Japan, ISIJ International, 2001, 41, 612. Shinoda T and Tanaka R, ‘Role of carbide precipitations on creep rupture strength of austenitic stainless steels (in Japanese)’, Bulletin Japan Institute Metals, 1972, 11, 180.
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91 Masuyama F, ‘Advanced power plant development and material experiences in Japan’, 8th Liège conference on Materials for Advanced Power Engineering, J. Lecomte-Beckers et al. (eds), Juelich GmbH, Forschungszentrum, Juelich, 2006, Liège, Belgium, 175. 92 Sawaragi Y, ‘Elevated temperature strength and microstructure of several high strength austenitic steels’, JSPS, 123 Committee Report on Heat Resistant Metallic Material, 1985, 26, 369. 93 Tamura M, Yamaguchi N, Tanimura M and Murase S, ‘Processing alloys for the heat exchangers of advanced coal fired boilers’, 1985 Exposition and Symposium, Industrial Heat Exchanger Technology, Pittburgh, PA, 1985. 94 Sawaragi Y and Yoshikawa K, ‘High temperature strength and microstructure of high strength and corrosion resistant austenitic steel for boiler (in Japanese)’, Tetsuto-Hagane, 1986, 72, S672. 95 Sawaragi Y, Teranishi H, Makiura H, Miura M and Kubota M, ‘The development of HR3C steel with high elevated temperature strength and high corrosion resistance for boiler tubes (in Japanese)’, Sumitomo Metals, 1985, 37, 166. 96 Kikuchi M, Sakakibara M, Otoguro M, Mimura H, Araki S and Fujita T, ‘An austenitic heat resisting steel tube developed for advanced fossil-fired steam plants’, International Conference High Temperature Alloys, Petten, Netherlands, 1985. 97 Toyama A, Minami Y and Yamada T, ‘Effect of alloying elements on high temperature properties in an austenitic stainless steel containing high Cr (in Japanese)’, CAMPISIJ, 1988, 1, 928. 98 Semba H, Igarashi M and Sawaragi Y, ‘Development of 23Cr–18Ni–3Cu–1.5W– Nb–N Austenitic Steel for USC Boilers’, Proceedings International Conference Power Engineering–97, Volume 2, JSME, Tokyo, 1997, 125. 99 Rautio R and Bruce S, ‘Sandvik Sanicro 25, a new material for ultrasupercritical coal fired boiler’, Advances in Material Technology for Fossil Power Plants, R. Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH, 274. 100 Masuyama F, in Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology, Group VIII, Volume 2, Materials Subvolume B, Creep Properties of Heat Resistant Steels and Superalloys, Yagi K, Merckling G, Kern T U, Irie H, Warlimont H, Springer Verlag, Berlin Germany 2004. 101 Kalwa G, Haarmann K and Janssen J K, ‘Experiences with ferritic and martensitic steels tubes and piping in nuclear and non-nuclear applications’, Topical Conference Ferritic Alloys for Use in Nuclear Energy Technology, Metallurgical Society AIME, Snowbird, ed. by J W Davis and D J Michel, 1983. 102 Sikka V K, Ward C T and Thomas K C, ‘Modified 9Cr-1Mo Steel – An Improved Alloy for Steam Generator Application’, ASM International Conference Production, Fabrication, Properties and Applications of Ferritic Steels for High-Temperature Applications, Warrendale, PA, ASM International, Materials Park, Ohio, 1981. 103 Sakakibara M, Masumoto H, Ogawa T, Takahashi and Fujita T, ‘High Strength 9Cr0.5Mo-1.8W(NF616) Steel for Boiler Tube (in Japanese)’, Thermal and Nuclear Power, 1987, 38, 841. 104 DeLong W T, Ostrom G A, Szumachowski E R, ‘Mearsurement and Calculation of Ferrite in Stainless Steel Weld Metal’, Weld Journal, 1956, 35, 521S. 105 Masuyama F, Nishimura N, Haneda H, Ellis F V and DeLong J F, ‘Investigation of the Deterioration plus Restoration Behaviour of fourteen Heats of TP316 Stainless Steel removed from Eddystone Unit No 1 Main Steam Lines after 130 520 h
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108 109
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111 112 113 114
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Service’, Properties of Stainless Steels in Elevated Temperature Service, MPC-Vol. 26/PVP-Vol 32, 1987, 173. Yamamoto R, Kadoya Y, Ueta S, Nobe T, Magoshi R, Nishimoto S and Nakano T, ‘Development of wrought Ni-based Superalloy with low Thermal Expansion for 700°C Steam Turbines’, Advances in Material Technology for Fossil Power Plants, R. Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH 623. Imano S, Doi H and Iijima E, ‘Modification of Ni-Fe base Superalloy for Steam Turbine Applications’, Advances in Material Technology for Fossil Power Plants, R. Viswanathan et al. (eds), 2005, ASM International, Metals Park, OH, 575. Imai K, private communication, 2006. Takeda Y and Masuyama F, ‘Heat Resistant Materials for Ultra Super Critical Power Plants (in Japanese)’, JSPS 123 Committee Report on Heat Resistant Metallic Material, 1987, 29, 399. Furuya K, Hizume A, Takeda Y, Fujikawa T, Fujita A, Kinoshita S, Kohono M, Honjo T and Suzuki A, ‘Austenitic Steel Rotor Forging for EPDC’s WAKAMATSU 50MW High Temperature Turbine Project’, Proc. 2nd Int. Conf. Improved CoalFired Power Plants, Palo Alto CA, USA 1989, Electric Power Research Institute, 1989, 59–1. Ohta S, ‘Heat Resistant Steels for Chemical Industries (in Japanese)’ JSPS 123 Committee Report on Heat Resistant Metallic Material, 1977, 18, 383. Koori M, Yosida T and Ohta S, ‘Analysis and Prevention of Failures in Steam Reforming Furnace (in Japanese)’, Kobe Steel Technical Bulletin, 1983, 33, 65. Weißbach W, Werkstoffkunde und Werkstoffprüfung, Vieweg & Sohn, Braunschweig/ Wiesbaden, Germany 1979. Bodnar R L and Cappepellini R F, ‘Effects of residual elements in heavy forgings, past, present, future’, Research Department Publication Bethlehem Steel Corporation, Bethlehem, PA, Philadelphia, USA, 1986. Schmollgruber F, Verfahrenswege zur Herstellung grosser Schmiedestücke und deren qualitative und wirtschaftliche Auswirkungen, Dissertation TH Aachen, February 1974. Nutting J and Viswanathan R, ‘Clean steel: superclean steel’, Conference on Clean Steel: Superclean Steel, Conference Proceedings, 6–7 March 1995, London, UK, 1995. Viswanathan R, ‘Clean steel technology’, Proceedings of the R.I. Jaffee Memorial Symposium, 2–5 November 1992, Chicago, Illinois, USA, ASM International, Materials Park Ohio, 1997. Azuma T, Tanaka Y, Ikeda Y and Yoshida H, ‘Production and properties of superclean 3.5%NiCrMoV rotor forgings for low pressure steam turbine’, Proceedings of the R.I. Jaffee Memorial Symposium, 2–5 November 1992, Chicago, Illinois, USA, 1992, ASM International, Materials Park Ohio, 213–220.
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3 Specifications for creep-resistant steels: Europe G . M E R C K L I N G, RTM BREDA Milano, Italy
3.1
Introduction
Creep problems with metallic materials emerged simultaneously in the USA and in Europe in the 1920s. In both cases, efforts to increase efficiency had led to higher service temperatures in pressure and other load-bearing systems. Unfortunately, on both sides of the Atlantic ignoring creep effects resulted in several casualties and in some cases severe injuries were sustained. As a consequence, researchers in the USA and the UK started to investigate the effect and the first publications describing creep appeared in the mid-1920s.1–3 The search for a simple answer to creep, encouraged by the decreasing strain rate of the primary creep stage, suggested to early researchers that there was a creep threshold stress, that is a stress level beyond which no creep, or at least no creep fracture, would occur. A particular testing type4 was created in Germany in the early 1930s. Criteria that helped identify ‘creep-resistant’ materials at a given temperature were defined by assuming that materials fulfilling these properties (creep strain rate lower than 10–3% h–1 between 25 and 35 h of testing and creep strain smaller than 0.2% in 45 h) would not fail and were applied extensively in parts of Europe. Again, it was casualties that proved this assumption wrong. It thus became apparent that expensive and difficult long-term creep investigations were needed to guarantee operational safety and plant reliability at high temperatures. All the major European countries developed activities to study, understand and predict creep strength and deformation behaviour. Building on efforts in the 1930s, research was re-activated after World War II and has continued ever since. Although aimed at the same target – for the best use of materials at the best safety levels with the available technology – different approaches to the investigation of creep were developed. Essentially three different directions were taken, although other interpretations and sub-fields grew with time: 78 WPNL2204
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1. German industry developed a common effort approach from the outset, which, coordinated by an umbrella organisation, pooled experiences and data from most industrial companies, research and university institutions, including very long-term creep results. As a result, a huge number of steels were characterised, optimised and developed for application in creep regime, and creep strength data containing DIN standards were issued. From this activity, perhaps the only raw long-term creep data collation was published in 19695 (see also Table 3.1), and this formed the basis of the publicly available data pool throughout Europe for a long period. In parallel, technical rules for material properties, welding and component design were issued in the TRD and AD series. 2. In the United Kingdom, creep investigation efforts were more or less confined to big industrial companies operating independently, with less national coordination. Nevertheless, here also a large series of materials were developed on a company-to-company basis for use in creep regimes, accompanied by the related BS and PD standards and rules for creep strength data assessment, welding and component design. 3. A third approach was followed in the 1950s by some of the less industrialised countries, like Italy, who had fewer resources to develop their own materials and/or rules and so relied on external sources. Italian design and material standards in creep regimes were based on a unique mixture of European (German and British) and US research, assessed to suit the local circumstances. In these countries, US material nomenclature, design criteria and codes were used as widely as European specifications. Other countries, like Sweden and France, also developed a significant amount of technical documentation for material applications in creep regimes. Further developments followed the introduction of nuclear power to Europe, mainly concerning austenitic materials, nuclear pollution and the effects of irradiation. As this a very specialised area, the findings will not be discussed here, although it is acknowledged that nuclear research produced valuable data, some of which was made available for the development of new European Standards and joint assessment activities in the European Creep Collaborative Committee (ECCC) (see Section 3.2.5). With the introduction of the Common Market in Europe in the 1960s and the increasing importance of exports within and outside Europe, several panEuropean institutions began attempting to harmonise the various different standards. The European Carbon and Steel Collaboration programmes CECA/ ECSC6 had already begun this process in the 1950s and continued into the 1960s. Since 1971, most major work has been coordinated through the programmes of the organisation for European Cooperation in the field of Scientific and Technical Research (COST)7 and these have contributed
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significantly to innovation and the development and characterisation of new materials for high temperature applications, such as rotor steels, advanced piping materials, welding techniques, and so on. More focused front-line research has been concentrated in the EC funded Brite-Euram-Projects since the 1980s and technical, pre-competitive development of future technological solutions has been through various Thermie-Projects,8 which began in 1993. All these activities have helped European industry and research specialists to exchange their experience and knowledge and have encouraged the formation of a strong core of organisations collaborating while competing in the high temperature field. On the standards development side, right from the beginning the European Community tried to find common denominators for relevant technical issues that could be published as Euronorms (abbreviated as EU). However, the first series of Euronorms were not very well accepted, did not gain general application and made little difference to the use of independent national standards. The situation changed when the now European Union started to issue ‘directives’, which the member countries had agreed would become law and so would formalise voluntary industrial standardisation. For the high temperature sector, the relevant event was the publication in 1997 and mandatory introduction in 2002 of the Pressure Equipment Directive (PED) 97/23/EC,9 which required guarantees for the in-service performance of materials and, as a consequence, required harmonisation of the relevant EN standards. Although not dealing specifically with pressure vessels under a creep regime, the PED required a new series of standards harmonised at all levels for high-temperature materials, for which ‘European’ creep strength values were to be defined. The PED came from the European Committee for Standardisation (CEN) and its related technical committee, the European Committee for Iron and Steel Standards (ECISS) and, under the framework of the European Community funded research programmes, encouraged a large group of industries, research and inspection organisations to form: • •
the European Pressure Equipment Research Council (EPERC), which strongly contributed to implementation and clarification of the PED directive itself; and the European Collaborative Creep Committee (ECCC), which developed creep strength data for some of the materials under the new harmonised EN standards, using new data assessment approaches.
The following sections and tables will try to give a concentrated overview of the most relevant standard publications related to creep-resistant steels suitable for structural applications, beginning with the original national grades up to the current PED-harmonised standards.
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3.2
Specifications and standards
3.2.1
Institutions, publishing rules and standards
81
Each European country publishes norms for voluntary industrial standardisation, which are issued by various organisations, as follows: • •
•
•
•
The national standardisation body (DIN in Germany, BS in the United Kingdom, AFNOR in France, UNI in Italy, etc.) is responsible for preparing and issuing voluntary industrial standardisation documents. Some of the national standardisation bodies, their sub-organisations or umbrella organisations also issue ‘second-level standardisation documents’, such as PD documents in the United Kingdom, SEP documents in Germany, and so on. Ad hoc industrial committees, in some cases affiliated to the national standardisation body, issue design codes for particular areas of technical interest. For example, for pressure application in creep regimes, design rules are issued by TRD and AD in Germany, CODAP in France, CTI with ANCC/ISPESL in Italy, and so on. Often these codes do not just specify a national standard for materials, but add additional prescriptions, controls or requirements. In some countries, inspection authorities have issued additional materialrelated specifications (e.g. the ‘Raccolta M/S’ of the Italian Authority ISPESL) or have required qualifications for suppliers of materials to be used in creep regime (e.g. VdTÜV-Richtlinien in Germany). A few major users or builders of equipment operating under creep regimes have developed design and material qualification guidelines (e.g. National Electricity Boards, refinery and petrochemical groups, turbine and boiler manufacturers, etc).
Since the PED Directive was issued in 1997, standards across Europe have tended towards unification. Standards issued by CEN, the European Standardisation Body, are jointly prepared by specialists from all CEN member countries and some which are not yet members, and are quickly substituted for the corresponding national standards. Under the CEN framework10 for steels ECISS11 is responsible for issuing harmonised standards. Experts from industry and research, delegated by national standards bodies, are recruited into Technical Committees and advisory organisations provide specific information as required, or, as is often the case, via their specialists being members of both the relevant CEN/ECISS Technical Committee issuing the standard and the advisory organisation (e.g. ECCC, EPERC, etc). The European Commission itself is ultimately in charge of defining EUwide regulations in the form of directives, with the objective of standardising approaches and methods, both technical and administrative, within the member countries to simplify and encourage transport, trade, exchange and
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implementation of goods. EU directives normally become law in all EU member countries within 12 to 18 months, superseding previous local legislation. Technical directives often call for common standardisation or refer directly to EN standards, significantly increasing the importance of CEN documents and their influence on industry and administration. For metallic materials used in creep regimes, the issue of the PED in 1997 followed by mandatory application starting in 2002 was a critical step.
3.2.2
The pressure equipment directive
The PED9 deals with all relevant safety issues in the design, manufacture and use of components designed to resist stresses induced mainly by pressure. The directive is aimed at simplifying intra-European trade and exchange of pressure handling equipment and harmonising safety prescriptions in order to allow intra-European exchangeability and serviceability of pressure devices. There are many types of boilers, pipelines, furnaces, heat exchangers, incinerators, chemical reactors, etc within Europe, and all have to conform to the directive. Explicitly excluded are all equipment regulated by other directives (e.g. pressure equipment for transportation on rail or road, ‘simple pressure equipment’, etc), and all types of rotating machines and motors (e.g. combustion engines, turbines and pumps, etc.). The PED is a general law (not a design code) and does not directly consider specific design and safety factors for components operating under creep regimes, although creep is mentioned and regarded as a non-negligible issue in designing. For dimensioning, including equipment operating in creep regime, PED allows the use of any internationally recognised standard and accepts fracture mechanics approaches and finite element simulation, but for materials, the rules are more strict: • • •
All materials must have a guaranteed minimum impact energy at the lowest possible service temperature that the equipment for which they are used may encounter under pressure, e.g. including hydraulic tests. All materials must demonstrate ‘sufficient’ ductility in service conditions to guarantee ‘leak before break’. Materials need to be selected from harmonised standards, i.e. material standards that in their intention and scope consider and include the safety criteria as set forth by the PED. This requirement strongly enhanced and accelerated the production of EN standards incorporating Europewide agreed strength values. If materials cannot be selected from EN standards (typically because they come under non-European design codes, e.g. where ASME Boiler and Pressure Vessel Code or API standards are used), they must either pass a general acceptance procedure conducted by an accredited, notified body, or they must undergo a ‘particular material
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appraisal’ to check their suitability for the application and their conformity to the PED safety criteria, on a case-by-case basis. A final, and in the creep regime very relevant, PED requirement is that all material producers involved in the production chain of a given component must declare and take responsibility that the material supplied (possibly with provisos relating to particular handling, assembly or design operations) is suitable for the intended application and its operating conditions. Alternatively, in a somewhat milder interpretation, producers must certify that the material is indeed fully compliant with the purchase specifications and the design-relevant material properties so that it can be guaranteed for safe service. For materials and welds in the creep regime this requires the material, shape, fitting and so on producer to guarantee: – – – –
long term strength of the base material long term strength of the construction welds long term strength of the assembly welds long-term strength of any repair (if allowed by the design code) of any of the items listed above.
The easiest way to comply is to select materials that conform with the harmonised standards, for which long-term strength values are indicated in the material norm. Materials subjected to the ‘particular material appraisal’ (typically ASTM or API grades, sometimes Japanese materials) have to demonstrate the points listed above, which can be achieved by testing the actual material or by presenting previous data for the same grade and producer. Owing to the expected long service duration of modern plants (generally 200 000–250 000 h, i.e. 25 to 35 years), these demonstrations often become extremely complex for welds, repairs and multiple repairs. Another route that is applied by some notified bodies and/or end user inspectorates is to qualify material producers as PED-compliant using stringent procedures when they first supply materials under the aegis of this notified body. This first-time qualification should include creep strength verification and demonstration through a reduced creep campaign, which, depending on the notified body/inspectorate, the type of material and the scope of supply, may include 2–3 isotherms made or composed from of three to five points each and durations of generally not less than 10 000 h, and often not less than 30 000 h.
3.2.3
Non PED creep applications
Equipment not subjected to PED regulations or the particular EU directives for dangerous goods and fluids, transportation or general safety in the work
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place, is built under the sole responsibility of its manufacturer. For high temperature applications, this is mainly the case for manufacturers of gas and steam turbines, compressors, pumps, geothermics, some types of welldrilling equipment, combustion engines, aircraft engine components and, in some cases, electrical or railway equipment. Creep data and strength determination in Europe in the past were strongly driven by turbine manufacturers. Based on their own large data collections, turbine manufacturers imposed creep testing and, less frequently, creep strength properties on components they bought from forges, steel works and foundries. Typical specifications included quite severe prescriptions and sometimes range limitations in chemical composition, stringent heat treatment details and control and mechanical testing, but generally only specifying ‘creep quality testing’ by means of one or a small group of short-term (24–100 h) stress rupture or time-to-rupture (see ASTM E292) tests. The turbine manufacturers themselves then often carried out large-scale, long-duration tests, so that the assessment of the creep properties of a single material batch was fully under their control.
3.2.4
European national standards reporting creep data
Before the era of CEN and the PED directive, creep rupture strengths were generally: • • •
assessed within the user companies (turbine and boiler manufacturers, or big end-users such as utilities, refineries and the chemical industry); supplied to support sales between companies (typically from steel maker or tube/pipe manufacturer to user); assessed by national associations (often under the control of a national electricity board acting as the main national customer).
Generally such assessments were not made available to the wider public and were confined to the members of the working group collating and preparing the strength values. In some cases, these were one-off activities, that is the collated or produced data were assessed by the working group and then disappeared into the databases of the companies. In other cases, fortunately, centralised databases were created which allowed data to be updated and strength values upgraded as the information available on creep expanded (for instance AGW within VDEh in Germany). National standardisation bodies made several attempts to include creep values into standards and rules, but their success was quite limited because long-term data were expensive to obtain and considered commercially sensitive, so they were often not made available to larger groups for common assessment. Nevertheless some of the older standards include creep strength information, sometimes as an informative appendix, sometimes as a true part of the standard.
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Table 3.1 gives some of the older standards, now mainly discontinued, including creep strength information. Table 3.2 gives examples of creep data for tubes and pipes in the main steel grades included in the older standards. This table also shows the difficulties encountered in achieving common standardisation, because the correlation between steel grades of different countries is not always obvious and their application and heat treatment conditions are sometimes quite different. It is relatively obvious that in spite of the national preferences for using particular steel grades as recommended by national producers, the main national standards also include grades for ‘foreign’ materials, although sometimes with quite significant differences in guaranteed/expected creep strength values.
3.3
The European Creep Collaborative Committee (ECCC)
3.3.1
Motivation and history
With advancing integration within Europe, in 1991 a large number of industrial and research companies from the United Kingdom, Germany, Italy, France, Denmark, Switzerland, Austria, The Netherlands and Belgium founded the European Creep Collaborative Committee (ECCC).12 The common purposes of this organisation were laid down in a Memorandum of Understanding (MoU) that was jointly agreed, signed by all members and is still the foundation for cooperation. The MoU established that ECCC is a voluntary association of organisations representing their nations, led by industry, with the aim of improving, reinforcing and enhancing the position of European industry in the high temperature application market. The companies participating in ECCC cover the whole scope of industries and research organisations dealing with high temperature problems: steel makers, steel product manufacturers, boiler, turbine, plant and equipment builders, utilities and end users, inspection bodies, research institutes, technical universities, testing houses. Currently, the ECCC represent 14 nations, which include, besides those already mentioned above, Sweden, Finland, Portugal, Czech Republic and Slovakia. More than 50 organisations from the 14 countries contribute to ECCC activities. The MoU defines detailed targets to be addressed by ECCC to support European standardisation and research in creep. ECCC and its members: • • • • • •
collate, exchange and jointly assess creep data; support European standardisation; coordinate European creep data generation; develop common creep research and testing programmes; mutually exchange information on HT material development; and define common procedures for data generation and assessment.
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Table 3.1 Former national European Standards including creep strength values (examples) Nation
Title
Date of issue
Status
BS 1501
UK
1988/90
Discontinued
BS 1502
UK
1982
Discontinued
BS 1503
UK
1989
Discontinued
BS 1506
UK
1990
Discontinued
BS 3059-2
UK
1990
Discontinued
BS 3602
UK
1987
Discontinued
BS 3604
UK
1991
BS 3605 DIN 17175 DIN 17176
UK D D
1991 1979 1990
Current (only part 2 – welded tubes) Discontinued Discontinued Discontinued
DIN 17177
D
1979
Discontinued
DIN 17240 DIN 17245
D D
1976 1977
Discontinued Discontinued
DIN 17458
D
Plates, Specification for carbon, alloy and austenitic stainless steel tubes with specified elevated temperature properties Section and bars: Specification for carbon, alloy and austenitic stainless steel tubes with specified elevated temperature properties Forgings: Specification for carbon, alloy and austenitic stainless steel tubes with specified elevated temperature properties Bars for bolting: Specification for alloy and austenitic stainless steel tubes with specified elevated temperature properties Steel boiler and superheater tubes. Specification for carbon, alloy and austenitic stainless steel tubes with specified elevated temperature properties Specification for carbon steel pipes and tubes with specified room temperature properties for pressure purposes Steel pipes and tubes for pressure purposes: ferritic alloy steel with specified elevated temperature properties. Specification for longitudinally arc welded tubes Austenitic Stainless steel pipes and tubes for pressure purposes Nahtlose Rohre aus warmfesten Stählen; Technische Lieferbedingungen, Nahtlose kreisförmige Rohre aus druckwasserstoffbeständigen Stählen; Technische Lieferbedingungen Elektrisch pressgeschweißte Rohre aus warmfesten Stählen; Technische Lieferbedingungen Heat resisting and highly heat resisting materials for bolts and nuts Ferritic Steel Castings Creep Resistant at elevated Temperatures: technical Conditions for Delivery Seamless circular austenitic stainless steel tubes subject to special requirements: Technical delivery conditions
1985
Discontinued
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Creep-resistant steels
Designation
D
DIN 17460
D
DIN 17465 DIN 17470
D D
NF A36-209
F
NF A49-213 NF A49-215 NF A49-219 BS PD 6525 UNI 5462
F F F UK I
UNI 6904
I
UNI 7660
I
VDEh
D
Nahtlose kreisförmige Rohre aus hochwarmfesten austenitischen Stählen: Technische Lieferbedingungen Hochwarmfeste austenitische Stähle. Blech, kalt- und warmgewalztes Band, Stabstahl und Schmiedestücke Heat resistant steel castings: Technical conditions for delivery Heating conductor alloys : Technical delivery conditions for round and flat wire Produits Sidérurgiques : Tôles en aciers inoxydables austénitiques et austéno-ferritiques pour chaudières et appareils à pression
Elevated temperature properties for steels for pressure purposes Tubi di acciaio senza saldatura: Tubi per caldaie, per apparecchi, per tubazioni di impianti termici ad alte temperature ed alte pressioni – qualità, prescrizioni e prove Tubi senza saldatura di acciaio legato speciale inossidabile resistente alla corrosione e al calore. Prodotti finiti di acciaio fucinati, per recipienti a pressione. Qualità, prescrizioni e prove. Ergebnisse deutscher Langzeitversuche
1992
Discontinued
1992
Discontinued
1977 1984
Discontinued Current
1990
Discontinued
1990 1964
Discontinued Discontinued Discontinued Current Discontinued
1971
Discontinued
1977
Discontinued
1961
na
Specifications for creep-resistant steels: Europe
DIN 17459
87
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Table 3.2 Example of creep data availability in some national European Standards for carbon, low and high alloyed ferritic steels for boiler tubes and pipes (all mentioned national standards discontinued in the meantime) Heat
Treat- BS 3604
grade
ment
BS 3059-2
Designation
United Kingdom
NF A49-213,
France
DIN 17175
Germany
Creep rupture strength
215,219
Creep rupture strength
DIN 17176
Creep rupture strength
T range
Duration
(°C)
(h)
Designation
T range
Duration
(°C)
(h)
Designation
T range
Duration
(°C)
(h) 10k, 100k 200k 10k, 100k 200k
P235GH
N
33
350–500
100k
St35.8
380–480
P265GH 8 CrMo 4 5 8 CrMo 5 5 13 CrMo 4 5
N NT NT N
45
350–500
100k
St45.8
380–480
620–460
450–630
13 CrMo 4 5
NT
620–440
450–630
13 CrMo 5 5
NT
621
450–630
625
450–650
10k, 30k, 50k, 100k, 150k, 200k, 250k* 10k, 30k, 50k, 100k, 150k, 200k, 250k* 10k, 30k, 50k, 100k, 150k, 200k, 250k* 10k, 30k, 50k, 100k, 150k, 200k, 250k*
X 11 CrMo 5 A
X 11 CrMo 5 NT QT
TU 10CD5-05
500–575
10k, 100k
TU 13CD4-04
450–560
10k 100k
13 CrMo 4 4
450–570
10k, 100k 200k
TU Z12CD5-05 500–600
10k, 100k
12 CrMo 19 5 480–650
TU Z12CD5-05 500–600
10k, 100k
12 CrMo 19 5 350–650
10k, 100k, 200k 10k, 100k, 200k
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UNI 5462
Italy Ceep rupture strength
Designation
T range
Duration
(°C)
(h)
C14
390-520
100k
C18
390-520
100k
14 CrMo 3
450–600
100k
Creep-resistant steels
EN 10216 Material
16 Mo 3
N
11 CrMo 9 10 NT
243
450–550
622
460–610
TU 15 D3
450–550
10k, 100k
15 Mo 3
450–500
10k, 100k 200k
16Mo5
450–550
100k
TU 10CD9-10
500–625
10k, 100k
10 CrMo 9 10 450–600
10k, 100k, 200k 10k, 100k, 200k
12 CrMo 9 10
470–620
100k
480–580
10k, 100k, 200k
12CrMoV 6 2
11 CrMo 9 10 QT 12 CrMoV 6 2 NT
X 11 CrMo 9 1
X 11 CrMo 9 1
12 CrMo 9 10 400–520 660
A
NT
450–600
629
629–590
X 10 NT CrMoVNb9 1
‘T91’
490–690
X 20 CrMoNiV 11 1
NT
762
10k, 30k, 50k, 100k, 150k, 200k, 250k*
14 MoV 6 3
450–640 50k, 100k, 150k, 200k, 250k*
10k, 30k,
440–670 50k, 100k, 150k, 200k, 250k*
10k, 30k,
10k, 30k, 50k, 100k, 150k, 200k, 250k* 480–680 50k, 100k, 150k, 200k, 250k*
TU Z10CD VNb09-01
TU Z10CD09
TU Z10CD09
550–625
10k, 100k
X 12CrMo9 1 460–600
10k, 100k, 200k
450–675
10k, 100k
X 12CrMo9 1 400–650
10k, 100k, 200k, 10k, 100k
X 10CrMoVNb 91
10k, 30k, 12 1
470–650
X20 CrMoV 470–650 200k
X 10 CrMoNiVNb 91 10k, 100k CrMoNiV 11 1 1
X 20
Specifications for creep-resistant steels: Europe
10k, 30k, 50k, 100k, 150k, 200k, 250k* 10k, 30k, 50k, 100k, 150k, 200k, 250k*
89
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Creep-resistant steels
3.3.2
Organisation
Being a voluntary group, ECCC has no basic resources or funding to rely on. Nevertheless, there are several sources which, owing to the multi-national configuration of the ECCC, provide at least some limited support: • •
•
Over the last ten years, ECCC has run European Community funded projects.13–15 Under certain circumstances, ECCC or ECCC sponsored sub-groups have applied for and obtained European research projects, during which detailed problems identified through ECCC activities could be discussed, tested and satisfactorily solved. In some cases these SMT-Projects16,17 formed the basis for new European Standards (e.g. EN 10291 (creep testing) or 10319 (stress relaxation testing)). Each member nation contributes a balanced mixture of testing activity each year, so that testing resources are available each year for commonly agreed research programmes.
This strategy compensates for the recent reductions in national and EC research budgets and the loss of companies and related specialists dealing with high temperature resulting from globalisation. It has also allowed the ECCC to continue to contribute to debate on safety-related technical issues, which are becoming more and more stringent as requirements for reliability and foreseeability of component behaviour increase. Included in this are the new flexible service conditions, which introduce new hazards by pushing plants into not previously experienced operating conditions. The ECCC is led by a management committee, in which all member nations have a seat and a vote. Technical work is done in: • • • •
Working Group 1: Creep Data Assessment and Generation Procedures, in which specialists from all over Europe can take part Working Group 2: Results Dissemination and Exchange Forum18 Working Group 3: Creep Data Assessment, divided into sub-groups A (ferritic low and high alloyed steels), B (austenitic steels), C (nickel base alloys) and Working Group 4: Creep of Components and Features, which can both only be attended by specialists belonging to organisations that contribute to their country’s fee to the organisation.
3.3.3
ECCC’s contribution to standardisation
The ECCC WG3 subgroups were charged with producing the strength values to be proposed to the CEN Technical Committees working on the new European Standards in 1991/92. To comply with this task, the WG3 groups:
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Specifications for creep-resistant steels: Europe
• • • • • •
91
identified the steel grades which were relevant to industry and should be considered for standardisation; collated raw creep data from all available, reliable sources in Europe (and in some cases, the world); assessed the credibility of the raw data collated; identified gaps within the available data; assessed creep rupture and creep strain strength values for the standardisation body; and promoted and led ECCC harmonised testing programmes relying – at least partially – on common testing resources. Since 1997, more than five million hours of stress rupture and creep rupture testing has been added to the collated data bases of most relevant steel grades (including X10CrMoVNNb 9 1 (grade 91), X 10 CrWMoVNb 9 2 (grade 92), X 19 CrWMoVNb9 1 1 (grade 911), X 6 NiCrMoTiBN 25 20 (grade 709), NiCr22Co12Mo (alloy 617), NiCr20TiAl (alloy 80A), 9%Cr welds, dissimilar welds (2,25%Cr to 9%Cr steels), aged materials (nickel base alloys), etc).
A strategically essential point in this was the need to establish: • • •
a modern system for collating and exchanging creep data in order to guarantee completeness of information and the fair involvement of all contributors; an objective system to judge and to guarantee quality of experimental data for old and future (including those jointly produced) creep tests; and a reliable, commonly agreed procedure on how to derive creep strength values from the large data population collated all over Europe.
A dedicated Working Group (WG1) was established to deal with this central problem and agreed in 1993/9419 on a common approach which was subsequently applied to all ECCC strength value determinations and was enhanced and enlarged several times up to 2005.20 Significant innovations included: • • •
a detailed written procedure by which the whole process, from setting up the detailed specifications of the material to be assessed for creep strength to the presentation of the results, is defined; defined criteria for evaluating statistical significance and material properties of a data set (Table 3; source: ECCC Recommendations, Reference 20; Vol 5 part Ia) and EN 12952-2 Annex B); strict requirements for strength assessment: – strength values for standards to be assessed by two independent assessors — whose results may not diverge by more than 10%,
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Creep-resistant steels
— who have to use at least one of the methods whose procedure is given in (Reference 20: Vol 5 part Ia – app. D); — both assessments to be positively verified by objective post assessment tests (Reference 20: Vol 5, References 21, 22), checking the strength prediction physical credibility, data description quality and repeatability/stability of extrapolation (see Chapter 14, Constitutive equations for creep curves and predicting service life). Owing to the time needed to prepare all these agreed guidelines, the very early data assessments recommended by ECCC did not completely follow the WG1 procedures, but since 1995 all creep strength values published by the ECCC have been validated by this approach. A first collation of creep strength values was presented in 1996,23 and an upgraded and enlarged version in 2005.24 Further upgrades and assessments on other grades are expected over the coming years. ECCC-recommended values have been made available to the relevant CEN Technical Committees and their sub-committees by common members, discussed and sometimes iteratively agreed. ECCC strength values formed the basis for a large part of the creep strength values included in the new EN standards.
3.3.4
ECCC’s future
In spite of the difficult funding situation in industrial research in Europe, ECCC will continue to support CEN in providing reliable creep strength values. In order to avoid duplication, to enhance the interaction with PED topics and to optimise available resources, ECCC will tighten its collaboration with the European Pressure Equipment Research Council (EPERC).
3.4
European Pressure Equipment Research Council (EPERC)
3.4.1
Motivation and history
The European Pressure Equipment Research Council (EPERC)25 was founded in the early 1990s and established officially in 1995 by a group of European industry and research organisations. It was set up to collaborate in harmonising and unifying European laws, regulations, codes and standards on all pressurebearing equipment, including those operating under high temperature and creep regimes. In 2005, EPERC became the EPERC-Technological Platform, in the light of the new EC Framework VII research funding strategy. According to the EPERC Member Agreement,25 the main objectives of EPERC are the establishment of: •
a European Network to support the pressure equipment industry and small and medium enterprises (SMEs), in particular;
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Table 3.3 Recommended data size requirements according to Reference 20, Vol. 5 part I Interim-minimum requirements
•
≥3 tests at each of ≥3 temperatures, at intervals of 50–100°C
Original (TM1)
For ≥6 casts, there should be tu(T, σ0) observations from: • ≥5 tests at each of ≥3 temperatures in the design application range at intervals of 25–50°C
䊐 – ≥3 tests per temperature (different σ0) with tu,max≥ 10 kh
TM2
TM3
For datasets with ≥300 observations, originating from ≥10 casts, at ≥5 temperatures covering the range TMAIN ± ≥50°C For ≥5 casts, there should be tu(T, σ0) observations from:
For datasets with ≥500 observations, originating from ≥20 casts, at ≥5 temperatures covering the range TMAIN ± ≥50°C For ≥5 casts, there should be tu(T, σ0) observations from:
•
•
䊐 – ≥4 tests per temperature (different σ0) with tu,max≥ 40 kh 䊐 – ≥1 test per temperature with tu,max≥ 40 kh
≥5 tests at each of ≥2 temperatures in the design application range at an interval(s) of 25–50°C 䊐 – ≥4 tests per temperature (different σ0) with tu,max ≤ 35 kh 䊐 – ≥1 test per temperature with tu,max≥ 35 kh
Predicted strength values determined from an Interim minimum dataset shall be regarded as tentative until the data requirements defined in one of the target-minimum columns are obtained
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䊐 – ≥4 tests per temperature (different σ0) with tu ≤ 35 kh 䊐 – ≥1 test per temperature with tu,max≥ 35 kh
93
tu,max is the longest test duration available, σ0 is the initial test stress, I is the temperature.
≥5 tests at ≥1 temperature(s) in the design application range (at intervals of 25–50°C)
Specifications for creep-resistant steels: Europe
For ≥3 casts, there should be tu(T, σ0) observations from:
Target-minimum requirements
94
• • • •
Creep-resistant steels
short- and long-term research priorities for the European pressure equipment industry; coordination of cooperative research in the domain of pressure equipment and identification of funding sources for this research; dissemination of research results to the European industry and standardisation bodies; and joint European attitudes to pressure equipment safety and reliability.
EPERC was particularly successful in interconnecting European cooperative research with current design needs, construction, inspection and in-service activities and especially in trying to harmonise contradictory European standards, rules and laws. For this purpose, EPERC was and probably will remain the major discussion forum and investigation coordinator providing technical and scientific depth to support the ‘New approach to Technical Harmonisation and Standards’, which will be the privileged means of complying with the ‘Essential Safety Requirements of the new Pressure Equipment Directive (PED 97/23/EC)’. Another important issue for EPERC is maintaining permanent contacts with standardisation bodies (CEN) and the relevant international organisations, PVRC (USA) and JPVRC (Japan), and so on. This cooperation is supported by a common Memorandum of Understanding (2001) and involves a strong interaction with all the committees and notified bodies’ umbrella organisations dealing with the Pressure Equipment Directive (PED 97/23/EC)9 and its amendments, guidelines and sometimes conflicting interpretations. Since 2002, EPERC has entered several new areas, including support for the new European Design Codes (EN 13445, EN 13480, families etc), also in the area of creep. In this context, EPERC and ECCC have collaborated and will develop an even closer relationship in the future.
3.4.2
Organisation
EPERC has relied on in-kind contributions and financial support via ECfunded projects. It is led by a steering committee in which national representatives from the countries with member organisations have a seat and a vote. The original EPERC structure can be found in Reference 25. The actual organisation of EPERC-TP is structured into five task groups, dealing with (i) Materials & fabrication, (ii) Design, (iii) Damage & rupture analysis, (iv) Operation, inspection, testing & maintenance, (v) Special and emerging issues, and Support to Standardisation.
3.4.3
EPERC’s contribution to creep-related standardisation
EPERC has collaborated on and influenced a huge number of standards and pre-standardisation projects in all areas related to pressure equipment. In the
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Specifications for creep-resistant steels: Europe
95
high temperature area, essential contributions were made to the harmonisation of the non-destructive testing (NDT) codes and practices used for creepserviced component inspection, creep crack growth, minimum invasive testing (small punch) and the new philosophies founded on risk base analysis (RIMAP) and ‘fitness for service’ concepts (FITNET). Further EPERC contributions centred on the design codes and the related CEN TCs, especially for the new EN 13445 standards which include design issues in creep regime in Part 3. This particular project required information on parent material and weld creep strengths, hence the interaction with ECCC.
3.5
The latest generation of CEN standards for creep-resistant steels
3.5.1
General principle
The CEN10 standards are prepared by the relevant CEN Technical Committee and its sub-committees assisted, in some cases, by additional working groups. The new EN standards attempt to harmonise the often contrasting national norms and rules, focusing on those materials which are either present in most of the national standards, are not yet included in national standards but are of immediate technical interest, or are evidently commonly used throughout Europe. In some cases, similar grades codified in Partner National Standards can be grouped into one common type. In the case of standards related to creep, the framework covers revised methods for testing, specifications for steels, extensive use of creep strength data (although this is generally indicated as being informative), welds and – still mainly under development – design rules.
3.5.2
Material specifications
Table 3.4 includes an overview of the latest available CEN standards (up to November 2006, although several were expected in 2007) that relate to creep-resistant steels and reporting strength values. The materials included in these standards are listed in Tables 3.5 to 3.15, along with some information on the available creep strength values. A particular feature is the use of informative annexes reporting mean creep rupture strengths, which are determined from quality experimental data collations representing European and sometimes worldwide research. • •
Values marked * were obtained by extended extrapolation in time, i.e. by more than a factor of 3, beyond the scope of the available experimental data. Values in brackets ( ) were calculated by an extrapolation beyond 80% of the minimum stress of the experimental data set.
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Table 3.4 European specifications for creep-resistant steels Title
Date of issue
Status (Nov/2006)
EN10028-2
Flat products made of steels for pressure purposes – non-alloy and alloy steels with specified elevated temperature properties Flat products made of steels for pressure purposes – stainless steels Heat resisting steels and nickel alloys Technical delivery conditions for steel castings for pressure purposes – steel grades for use at room temperature and elevated temperatures Technical delivery conditions for steel castings for pressure purposes – austenitic and austeno-ferritic steel grades Seamless steel tubes for pressure purposes – Technical delivery conditions – Part 2: Non-alloy and alloy steel tubes with specified elevated temperature properties Seamless steel tubes for pressure purposes – Technical delivery conditions - Part 5: Stainless steel tubes Welded steel tubes for pressure purposes – Technical delivery conditions - Part 2: Electric welded non-alloy and alloy steel tubes with specified elevated temperature properties Welded steel tubes for pressure purposes – Technical delivery conditions - Part 5: Submerged arc welded non-alloy and alloy steel tubes with specified elevated temperature properties Welded steel tubes for pressure purposes – Technical delivery conditions – stainless steel tubes Steel forgings for pressure purposes – ferritic and martensitic steels with specified elevated temperature properties Steel forgings for pressure purposes – martensitic, austenitic and austeno-ferritic stainless steels Butt welding pipe fittings Steels and nickel alloys for fasteners with specified elevated and/or low temperature properties
2004
Current
2002 2001 1998
Current Current Current
1998
Current
2002
Current
2002
Current
2002
Current
2002
Current
2005
Current
2001
Current
2001
Current
1999 2001
Current Current
EN10028-7 EN10095 EN10213-2 EN10213-4 EN10216-2 EN10216-5 EN10217-2*
EN 10217-5*
EN10217-7* EN10222-2 EN10222-5 EN 10253* EN 10269
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Creep-resistant steels
Designation
EN 10295 EN 10302 EN 12952-2* EN 14532-2
Stainless steel bars for pressure purposes Hot rolled weldable steel bars for pressure purposes with specified elevated temperature properties Heat resistant steel castings Creep resistant steels, nickel and cobalt alloys Water tube boilers and auxiliary installations – Part 2: Materials for pressure parts of boilers and accessories Welding consumables: Test methods and quality requirements Part 2: Supplementary methods and conformity assessment of consumables for steel nickel and nickel-alloys
2003 2002
Current Current
2003 2002 2001
Current Current Current
2004
Current
* These standards do not include creep strength values themselves, but the materials included can be referred back to other standards in the table, where strength values are listed.
Specifications for creep-resistant steels: Europe
EN 10272* EN 10273
97
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Table 3.5 Chemical composition of carbon, low and high alloyed ferritic steels in European Standards Material grade
Gross chemical composition
GP240GH GP280GH P195 P235GH P245GH P250GH P265GH P280GH P295GH P305GH P355GH C7–C24 C35E 15 MnCrMoNiV 5 3 15 MnMoV 4 5 18 MnMo 4 5 20 Mn 5 20 MnMoNi 4 5 8 MoB 5 4 12 MoCrV 6 2 2 14 MoV 6 3 16 Mo 3 17 Mo 5 20 MnNb 6
N/ QT N–QT N N A, NT, QT N N N, NT, QT N N, NT, QT N N N, QT NT QT NT QT NT QT N QT N NT QT NT QT N, NT, QT N, NT, QT N
Concentration in mass% of the following elements C
Cr
Mo
Ni
V
0.18–0.23 0.18–0.25 <0.13 <0.16 0.08 –0.2 0.18–0.23 <0.2 0.08 –0.2 0.08–0.2 0.15–0.2 0.1–0.22 <0.24 0.32–0.39 <0.17 <0.18 <0.2 0.17–0.23 0.15–0.23 0.06 –0.1
– – <0.30 <0.30 – <0.30 <0.3 – <0.3
– – <0.08 <0.08 – – <0.08 – <0.08
– – <0.3 <0.3 – <0.3 <0.3 – <0.3
– – <0.02 <0.02 – <0.02 <0.02 – <0.02
<0.3 – <0.4 0.5–1 – <0.3 <0.4 <0.2 <0.2
<0.08 – <0.1 0.2–0.35 0.4–0.6 0.45–0.6 <0.1 0.45–0.6 0.4–0.5
<0.3 – <0.4 0.3–0.7 – <0.3 <0.4 0.4–0.8 –
<0.02 – – 0.05–0.1 0.04–0.08 – – <0.02 –
0.1–0.18 0.12–0.2
0.3–0.6 <0.3
0.5–0.7 0.25–0.35
– <0.3
0.22–0.28 –
<0.22
–
–
–
–
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Other
Mn 1–1.5 Mn 0.9–1.4 Mn 0.9–1.5 Mn 1–1.5 Mn 1–1.5 B 0.002–0.006
Mn 1–1.5; Nb 0.015–0.1 (Cr+Cu+Mo+Ni)>0.7
Creep-resistant steels
Heat treatment
QT NT NT IA NT QT NT NT NT QT NT N NT NT QT QT QT NT QT QT
21 25 40 40 40 42 15
CrMoV 5 7 CrMo 4 CrMo 5 6 CrMoV 4 6 CrMoV 4 6 CrMo 5 6 NiCuMoNb 5 6 4
X 3 CrAlTi 18 2
0.15–0.23 <0.15 0.08–0.14 0.08–0.15 0.08–0.15 0.08–0.15 0.10–0.15 0.10–0.15 0.09–0.17 0.1–0.15 0.08–0.18 0.08–0.18 <0.17 0.11–0.15 0.15–0.2 0.13–0.2 0.15–0.2 0.17–0.23 0.17–0.23 0.17–0.23
– 1–1.5 2.0–2.5 2–2.5 2–2.5 2–2.5 2.0–2.5 2.75–3.25 0.3–0.55 0.3–0.5 0.7–1.15 0.7–1.15 1–1.5 2–2.5 1–1.5 2–2.5 1.2–1.5 3–3.3 3.0–3.5 0.9–1.2
0.4–0.6 0.45–0.65 0.9–1.1 0.9–1.1 0.9–1.1 0.9–1.1 0.9–1.1 0.9–1.1 0.51–0.7 0.4–0.6 0.4–0.6 0.4–0.6 0.45–0.65 0.9–1.1 0.45–0.65 0.9–1.2 0.9–1.1 0.5–0.6 0.5–0.6 0.9–1.1
– <0.3 – – – – <0.3 <0.25 <0.22 – – – <0.3 <0.25 – – – – <0.3 <0.2
– – – – – – – 0.2–0.3 0.21–0.34 0.22–0.3 – – – 0.25–0.35 – – 0.2–0.3 0.45–0.55 0.45–0.55 0.6–0.8
QT QT NT QT NT QT NT QT
0.17–0.25 0.22–0.29 0.39–0.45 0.26–0.44
1.2–1.5 0.9–1.2 1.2–1.5 0.9–1.2
0.55–0.8 0.15–0.3 0.5–0.7 0.5–0.65
<0.6 – – –
0.2–0.35 – – 0.25–0.35
0.39–0.45 <0.17
1.2–1.5 <0.3
0.5–0.7 0.25–0.5
– 0.1–1.3
– –
A
<0.04
17.0–18.0
–
–
–
QT QT (!) QT QT
QT QT QT
QT
N 0.0015–0.016
Si 0.5–0.8
B 0.001–0.010; Ti 0.007–0.15
Cu 0.5–0.8; Nb 0.015–0.045 Si <1; Al 1.7–2.1; Ti 0.2+4*(C+N)– 0.8
Specifications for creep-resistant steels: Europe
G 20 Mo5 10 CrMo 5 5 10 CrMo 9 10 11 CrMo 9 10 11 CrMo 9 10 11 CrMo 9 10 12 CrMo 9 10 12 CrMoV 12 10 12 CrMoV 6 2 2 G 12 MoCrV 5 2 13 CrMo 4 5 13 CrMo 4 5 13 CrMoSi 5 5 13 CrMoV 9 10 G 17 CrMo 5 5 G 17 CrMo 9 10 G 17 CrMoV 5 10 20 CrMoV 13 5 20 CrMoV 13 5 5 20 CrMoVTiB 4 10
99
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Table 3.5 (Cont’d.) Material grade
Gross chemical composition Concentration in mass% of the following elements C
Cr
Mo
Ni
V
Other Nb 0.2–0.5;. W<0.7; B 0.005–0.015; Co 5–7 Si 0.7–1.4; Al 0.7–1.2 Si 0.7–1.4; Al 0.7–1.2 Si 0.7–1.4; Al 1.2–1.7 Si 0.5–1. Al 0.5–1 Nb 0.06–0.1
X 8 CrCoNiMo 10 6
QT
0.05–0.12
9.8 –11.2
0.5 –1
0.2–1.2
0.1–0.4
X X X X X X X X X X X X X
A A A A NT QT A NT1 NT2 IA A1 IA NT NT QT
<0.12 <0.12 <0.12 <0.12 0.08–0.12 0.08–0.15 0.08–0.15 0.08–0.15 0.08–0.15 0.08–0.15 0.08–0.15 0.08–0.15 0.09–0.13
12.0–14.0 17.0–19.0 23.0–26.0 6.0–8.0 8–9.5 4.0–6.0 4.0–6.0 4.0–6.0 4.0–6.0 8.0–10.0 8.0–10.0 8.0–10.0 8.5–9.5
– – – – 0.85–1.05 0.45–0.65 0.45–0.65 0.45–0.65 0.45–0.65 0.9–1.1 0.9–1.1 0.9–1.1 0.9–1.1
– – – – <0.3 – – – – – – – 0.1–0.4
– – – – 0.18–0.25 – – – – – – – 0.18–0.25
X 11 CrWMoVNb 9 1
NT
0.07–0.13
8.5–9.5
0.3–0.6
<0.4
0.15–0.25
GX 12 CrMoVNbN 9 1
NT QT
0.1–0.14
8.0–9.5
0.85–1.05
<0.4
0.18–0.25
X 12 CrMo 5
NT QT
0.10–0.15
<4–6
0.45–0.65
<0.3
–
10 10 10 10 10 11 11 11 11 11 11 11 11
CrAlSi 13 CrAlSi 18 CrAlSi 25 CrAlSi 7 CrMoVNb 9 1 CrMo 5 CrMo 5 CrMo 5 CrMo 5 CrMo 9 1 CrMo 9 1 CrMo 9 1 CrMoWVNb 9 1 1
WPNL2204
– – – W 0.9–1.1; Nb 0.06–0.1; N 0.05–0.09; B0.0005–0.005 W 1.5–2; Nb 0.04–0.09; N 0.03–0.07; B 0.001–0.006 Nb 0.06–0.1; N 0.03–0.07
Creep-resistant steels
Heat treatment
QT QT NT QT NT QT A A QT
0.08–0.15 0.12–0.19 <0.18 <0.18 <0.18 0.15–0.20 0.17–0.23
11–12.5 4–6 4.0–6.0 4–6 4–6 26.0–29.0 10–11.5
1.5 –2 0.45–0.65 0.45–0.65 0.45–0.65 0.45–0.65 – 0.5–0.8
2 –3.0 – – – – – 0.2–0.6
0.25–0.4 – – – – – 0.1–0.3
X 20 CrMoV 11 1 X 20 CrMoWV 12 1 X 20 CrMoWV 12 1 X 22 CrMoV 12 1 GX 23 CrMoV 12 1 GX 30 CrSi7 GX 40 CrSi 13 GX 40 CrSi 17 GX 40 CrSi 24 GX 40 CrSi 28 GX 130 CrSi 29 GX 160 CrSi 18
QT QT 700 QT800 QT QT A A A as cast as cast as cast as cast
0.17–0.23 0.17–0.24 0.17–0.24 0.18–0.24 0.2–0.26 0.2–0.35 0.3–0.5 0.3–0.5 0.3–0.5 0.3–0.5 1.2–1.4 1.4–1.8
10–12.5 11–12.5 11–12.5 11–12.5 11.3–12.2 6.0–8.0 12.0–14.0 16.0–19.0 23.0–26.0 27.0–30.0 27.0–30.0 17.0–19.0
0.8–1.2 0.8–1.2 0.8–1.2 0.8–1.2 1–1.2 <0.15 <0.5 <0.5 <0.5 <0.5 <0.5 <0.5
0.3–0.8 0.3–0.8 0.3–0.8 0.3–0.8 <1.0 <0.5 <1.0 <1.0 <1.0 <1.0 <1.0 <1.0
0.2–0.35 0.2–0.35 0.2–0.35 0.25–0.35 0.25–0.35 – – – – – – –
Si <1. N 0.15–0.25 Nb 0.25–0.55; N 0.05–0.1 W 0.4–0.6 W 0.4–0.6
Si Si Si Si Si Si Si
1–2.5 1–2.5 1–2.5 1–2.5 1–2.5 1–2.5 1–2.5
Specifications for creep-resistant steels: Europe
X 12 CrNiMoV 12 3 GX 15 CrMo 5 X 15 CrMo 5 1 X 16 CrMo 5 1 X 16 CrMo 5 1 X 18 CrN 28 X 19 CrMoNbVN 11 1
101
WPNL2204
Material grade
1% strain creep strength
T range (°C)
Heat-resisting steels and Ni alloys EN 10095
Creep rupture strength
T range (°C)
1% strain creep strength
T range (°C)
Creep rupture strength
T range (°C)
Castings EN 10213-2 Creep rupture strength
T range (°C)
Duration (h)
GP240GH
400–500
GP280GH P195 P235GH
400–500
10k, 100k, 200k 10k, 100k,
380–480
10k–100k
380–480
10k, 100k, 200k
P245GH P250GH P265GH
380–480
10k–100k
380–480
10k, 100k, 200k
P280GH P295GH
380–500
10k–100k
380–500
10k, 100k, 200k
P305GH P355GH
380–500
10k–100k
380–500
10k, 100k, 200k
Duration (h)
Duration (h)
C7–C24 C35E 15 MnCrMoNiV 53 15 MnMoV 4 5
WPNL2204
Duration (h)
Duration (h)
Creep-resistant steels
Plates EN 10028-2
Product form
102
Table 3.6 Available creep data for carbon, low and high alloyed ferritic steels: plates; heat-resistant steels and Ni alloys; castings
18 MnMo 4 5 20 Mn 5 20 MnMoNi 4 5 8 MoB 5 4 12 MoCrV 6 2 2 14 MoV 6 3 16 Mo 3
10k–100k
425–525
10 k, 100k
Na
na
450–490
10 k, 100k
450–530
10k–100k
450–530
10k, 100k 200k
17 Mo 5 20 MnNb 6 G 20 Mo5 10 CrMo 5 5 10 CrMo 9 10
450–600
11 CrMo 9 10 11 CrMo 9 10 11 CrMo 9 10 12 CrMo 9 10 na 12 CrMoV na 12 10 12 CrMoV 6 2 2 G 12 MoCrV 5 2 13 CrMo 4 5 450–570 13 CrMo 4 5 13 CrMoSi 5 5 13 CrMoV 9 10 G 17 CrMo 5 5
450–570 na
10k–100k
450–600
10k, 100k 200k
na na
400–520 400–550
10k, 100k 10k, 100k
10k–100k
450–570
10k, 100k, 200k
100k na
450–570 400–550
100k 10k, 100k
10k, 100k, 200k
450–600
10k, 100k
400–550
10k, 100k, 200k 10k, 100k, 200k
400–600
WPNL2204
103
G 17 CrMo 9 10
400–550
Specifications for creep-resistant steels: Europe
425–525
104
Table 3.6 (Cont’d)
Material grade
1% strain creep strength
T range (°C) G 17 CrMoV 5 10 20 CrMoV 13 5 20 CrMoV 13 5 5 20 CrMoVTiB 4 10 21 CrMoV 5 7 25 CrMo 4 40 CrMo 5 6 40 CrMoV 4 6 40 CrMoV 4 6 42 CrMo 5 6 15 NiCuMoNb 400–500 564 X 3 CrAlTi 18 2
Duration (h)
10k, 100k
Heat-resisting steels and Ni alloys EN 10095
Creep rupture strength
T range (°C)
400–500
1% strain creep strength
Duration (h)
T range (°C)
Duration (h)
Creep rupture strength
T range (°C)
Duration (h)
10k, 100k 500–900
1k, 10k
500–900 1k, 10k, 100k
X 8 CrCoNiMo 10 6 X 10 CrAlSi 13
500–900
1k, 10k
X 10 CrAlSi 18
500–900
1k, 10k
500–900 1k, 10k, 100k 500–900 1k, 10k, 100k
WPNL2204
Castings EN 10213-2 Creep rupture strength
T range (°C)
Duration (h)
400–600
10k 100k 200k
Creep-resistant steels
Plates EN 10028-2
Product form
X 10 CrAlSi 25
500–900
1k, 10k
X 10 CrAlSi 7
500–900
1k, 10k
na
500–670
10k, 100k, 200k
10k
460–625
10k
470–600
500–900
500–900 1k, 10k, 100k
105
X 19 CrMoNbV N 11 1 X 20 CrMoV 11 1
1k, 10k
10k 100k
Specifications for creep-resistant steels: Europe
X 10 CrMoVNb na 91 X 11 CrMo 5 X 11 CrMo 5 X 11 CrMo 5 X 11 CrMo 5 X 11 CrMo 9 1 X 11 CrMo 9 1 X 11 CrMo 9 1 X 11 CrMoWV Nb 9 1 1 X 11 CrWMoVNb 91 GX 12 CrMoVN Nb 9 1 X 12 CrMo 5 460–625 X 12 CrNiMoV 12 3 GX 15 CrMo 5 X 15 CrMo 5 1 X 16 CrMo 5 1 X 16 CrMo 5 1 X 18 CrN 28
500–900 1k, 10k, 100k 500–900 1k, 10k, 100k
WPNL2204
106
Table 3.6 (Cont’d)
Material grade
1% strain creep strength
T range (°C)
Duration (h)
Heat-resisting steels and Ni alloys EN 10095
Creep rupture strength
T range (°C)
1% strain creep strength
Duration (h)
T range (°C)
X 20 CrMoWV 12 1 X 20 CrMoWV 12 1 X 22 CrMoV 12 1 GX 23 CrMoV 12 1 GX 30 CrSi7 GX 40 CrSi 13 GX 40 CrSi 17 GX 40 CrSi 24 GX 40 CrSi 28 GX 130 CrSi 29 GX 160 CrSi 18
WPNL2204
Duration (h)
Creep rupture strength
T range (°C)
Duration (h)
Castings EN 10213-2 Creep rupture strength
T range (°C)
Duration (h)
400–600
10k 100k 200k
Creep-resistant steels
Plates EN 10028-2
Product form
Table 3.7 Available creep data for carbon, low and high alloyed ferritic steels; seamless tubes/pipes; forgings; bars for fasteners and bolts; bars Product form
Seamless tubes/pipes EN 10216-2 Creep rupture strength
Forgings EN 10222-2 1% strain Creep strength
Bars for fasteners and bolting EN 10269
Creep rupture strength
1% strain creep strength
Creep rupture strength
Bars EN 10273 1% strain Creep strength
Creep rupture strength
GP240GH GP280GH P195 P235GH
T range (°C)
Duration (h)
400–500
10k 100k 200k 250k
P245GH P250GH P265GH P280GH P295GH P305GH P355GH C7–C24 C35E 15 MnCr
380–480
400–500
Duration (h)
10k–100k
T range (°C)
380–480
Duration (h)
T range (°C)
Duration T range (h) (°C)
Duration (h)
10k, 100k, 200k
10k 100k 200k 250k
T range (°C)
Duration T range (h) (°C)
Duration (h)
380–480
10k– 100k
380–480
10k, 100k, 200k
380–480
10k– 100k 10k – 100k
380–480
10k, 100k, 200k 10k, 100k, 200k
380–480 380–480
10k–100k
380–480
380–480
10k–100k
380–480
10k, 100k, 200k 10k, 100k, 200k
na
na
400–500
10k, 100k, 200k
450–500
10k–100k
430–500
10k–100k
380–500 380–500 350–500
10k 100k 350–500
10k 100k
na
na
10k 100k 200k
WPNL2204
380–500
10k– 100k 10k– 100k
380–480
380–500 380–500
10k, 100k 200k 10k, 100k, 200k
107
MoNiV 5 3 15 MnMoV 4 5 18 MnMo 4 5 20 Mn 5 20 MnMoNi 4 5 8 MoB 5 4
T range (°C)
Specifications for creep-resistant steels: Europe
Material grade
108
Table 3.7 (Cont’d.) Product form
Creep rupture strength
Forgings EN 10222-2 1% strain Creep strength
Bars for fasteners and bolting EN 10269
Creep rupture strength
1% strain creep strength
Creep rupture strength
Bars EN 10273 1% strain Creep strength
Creep rupture strength
Material grade
T range (°C) 12 MoCrV 6 2 2 14 MoV 6 3 450–600
T range (°C)
Duration (h)
T range (°C)
Duration (h)
10k, 100k, 480–600 200k, 250k 10k, 100k, 450–530 200k, 250k
10k 100k
450–600
10k–100k
450–530
10 k, 100k, 200k 10k, 100k, 200k
Duration (h)
16 Mo 3
450–550
17 Mo 5 20 MnNb 6
400–500
10k, 100k 200k* 250k*
G 20 Mo5 10 CrMo 5 5
450–600
10 CrMo 9 10
450–600
10k, 100k 200k, 250k 10k, 100k 200k, 250k
11 CrMo 9 10 11 CrMo 9 10
400–520
10k, 100k
450–600
10k–100k
450–600
T range (°C)
Duration T range (h) (°C)
Duration (h)
T range (°C)
Duration T range (h) (°C)
Duration (h)
450–530
10k – 100k
450–530
10k, 100k, 200k
450–600
10k, 100k
450–600
10k, 100k, 200k
400–520
100k
10k, 100k, 200k
11 CrMo 9 10 12 CrMo 9 10 12 CrMoV 12 10 12 CrMoV 622 G 12 MoCrV 52
WPNL2204
Creep-resistant steels
Seamless tubes/pipes EN 10216-2
13 CrMo 4 5
450–600
450–570
10k–100k
450–570
10k, 100k, 200k
450–570
10k, 100k na
na
25 CrMo 4 40 CrMo 5 6 40 CrMoV 4 6
420–550
450– 600 10k, 100k, 200k 10k, 100k 420–550 10k, 100k, 200k 10k, 100k 420–550 10k, 100k
na
na
450–550
10k, 100k 200k
40 CrMoV 4 6 42 CrMo 5 6
na
na
450–550
10k, 100k, 200k
400–500
10k, 100k
500–650
10k, 100k 200k*
na
na
500–670
450–570
10k, 100k, 200k
10k, 100k, 200k
109
15 NiCuMoNb 564 X 3 CrAlTi 18 2 X 8 CrCoNiMo 10 6 X 10 CrAlSi 13 X 10 CrAlSi 18 X 10 CrAlSi 25 X 10 CrAlSi 7 X 10 CrMoV Nb 9 1 X 11 CrMo 5
420–550
10k 100k
Specifications for creep-resistant steels: Europe
13 CrMo 4 5 13 CrMoSi 5 5 13 CrMoV 9 10 G 17 CrMo 5 5 G 17 CrMo 9 10 G 17 CrMoV 5 10 20 CrMoV 13 5 20 CrMoV 420–550 13 5 5 20 CrMoVTiB 4 10 21 CrMoV 5 7
10k, 100k, 200k 250k
WPNL2204
110
Table 3.7 (Cont’d.) Product form
Creep rupture strength
Forgings EN 10222-2 1% strain Creep strength
Bars for fasteners and bolting EN 10269
Creep rupture strength
1% strain creep strength
Creep rupture strength
Bars EN 10273 1% strain Creep strength
Creep rupture strength
Material grade
T range (°C)
Duration (h)
X 11 CrMo 5
450–600
X 11 CrMo 5
450–600
X 11 CrMo 5
450–630
10k, 100k 200k, 250k 10k, 100k 200k, 250k 10k, 100k 200k, 250k
X 11 CrMo 9 1 X 11 CrMo 9 1 X 11 CrMo 9 1
460–600 450–650
T range (°C)
Duration (h)
T range (°C)
Duration (h)
Duration T range (h) (°C)
Duration (h)
490–570
10k 100k 450–600
10k, 100k, 200k
10k, 100k 10k, 100k, 200k
X 11 CrMoWV Nb 9 1 1 X 11 CrWMoV Nb 9 1 GX 12 CrMo VNbN 9 1 X 12 CrMo 5 X 12 CrNiMoV 12 3 GX 15 CrMo 5 X 15 CrMo 5 1 X 16 CrMo 5 1
T range (°C)
490–570
10k, 100k
450–600
10k, 100k, 200k
WPNL2204
T range (°C)
Duration T range (h) (°C)
Duration (h)
Creep-resistant steels
Seamless tubes/pipes EN 10216-2
X 16 CrMo 5 1
10k 100k 200k
10k, 100k
450–600
10k, 100k, 200k
470–650
10k 100k
480–630
10k, 100k, 200k
450–600
10k, 100k 450–600
10k, 100k, 200k
450–600
10k, 100k 450–600
10k, 100k
Specifications for creep-resistant steels: Europe
X 18 CrN 28 X 19 CrMoN bVN 11 1 X 20 CrMoV 480–650 11 1 X 20 CrMoWV 12 1 X 20 CrMoWV 12 1 X 22 CrMoV 12 1 GX 23 CrMoV 12 1 GX 30 CrSi7 GX 40 CrSi 13 GX 40 CrSi 17 GX 40 CrSi 24 GX 40 CrSi 28 GX 130 CrSi 29 GX 160 CrSi 18
450–600
111
WPNL2204
112
Table 3.8 Available creep data for carbon, low and high alloyed ferritic steels: castings; creep-resistant steels; consumables; generic products Product form
1% strain creep strength
Creep-resistant steels EN 10302 Creep rupture strength
1% strain creep strength
Consumables EN 14532-2
Creep rupture strength
Creep rupture
Generic ECCC Creep rupture strength
Material grade
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration (h)
10k 100k 200k, 250k, 10k, 100k, 200k, 250k,
GP240GH GP280GH P195
x
400–500
P235GH
x
400–500
P245GH P250GH P265GH
x
400–500
10k, 100k, 200k, 250k,
P280GH P295GH P305GH P355GH
x
400–500
10k, 100k, 200k, 250k,
400–550
10k, 100k,
C7–C24 C35E 15 MnCrMo NiV 5 3 15 MnMoV 4 5 18 MnMo 4 5 20 Mn 5 20 MnMoNi 4 5 8 MoB 5 4
WPNL2204
Creep-resistant steels
Castings EN 10295
12 MoCrV 6 2 2 14 MoV 6 3 16 Mo 3
x
450–550
10k, 100k, 200k, 250k
17 Mo 5 20 MnNb 6 G 20 Mo5 10 CrMo 5 5 10 CrMo 9 10 11 CrMo 9 10 11 CrMo 9 10
x
450–600
x
400–520
10k, 100k, 200k, 250k 10k, 100k
CrMo 9 10 CrMo 9 10 CrMoV 12 10 CrMoV 6 2 2
G 12 MoCrV 5 2 13 CrMo 4 5
x
13 CrMo 4 5 13 CrMoSi 5 5 13 CrMoV 9 10 G 17 CrMo 5 5 G 17 CrMo 9 10 G 17 CrMoV 5 10 20 CrMoV 13 5 20 CrMoV 13 5 5 20 CrMo VTiB 4 10 21 CrMoV 5 7
x xNotch
10k, 100k, 200k, 250k
450–600
10k, 100k, 200k, 250k
420–550
10k, 100k
450– 600
10k, 30k, 100k, 200k 10k, 30k, 100k 200k, (1%) 10k, 100k
420–550
25 CrMo 4
420–550
WPNL2204
113
475–600
Specifications for creep-resistant steels: Europe
11 12 12 12
x
114
Table 3.8 (Cont’d.) Product form
1% strain creep strength
Creep-resistant steels EN 10302 Creep rupture strength
1% strain creep strength
Consumables EN 14532-2
Creep rupture strength
Creep rupture
Generic ECCC Creep rupture strength
Material grade
T range (°C)
T range (°C)
T range (°C)
T range (°C)
T range (°C)
Duration (h)
42 CrMo 5 6
450–550
15 NiCuMo Nb 5 6 4 X 3 CrAlTi 18 2 X 8 CrCo NiMo 10 6 X 10 CrAlSi 13 X 10 CrAlSi 18 X 10 CrAlSi 25 X 10 CrAlSi 7 X 10 Cr
400–500
10k, 30k, 100k, 200k 10k, 100k
500–670
10k, 30k,
450–630
100k, 200k, 10k, 100k, 200k, 250k
Duration (h)
Duration (h)
Duration (h)
Duration (h)
40 CrMo 5 6
x xNotch x x xNotch
40 CrMoV 4 6 40 CrMoV 4 6
470 –650
10k, 100k
MoVNb 9 1 X 11 CrMo 5 X 11 CrMo 5 X 11 CrMo 5 X 11 CrMo 5
500–600
10k, 100k
470 –650
10k, 100k, 200k
x
x x x
WPNL2204
Creep-resistant steels
Castings EN 10295
X 11 CrMo 9 1
450–640
X 11 CrMo 9 1 X 11 CrMo 9 1
x x 480–650
10k, (100k)
10k, 100k
520–650 470–620 450–600
X 12 CrNi
600–800 600–900
10k 100. 1k 10k
600–800
500–600
10k, 100k
450–600
10k, 100k
450–600
470–650
10k, 100k
470–650
470–650
10k, 100k
470–650
450–600
10k, 100k
450–600
10k, 100k, 200k 10k, 100k, 200k 10k, 100k, 200k 10k, 100k
450–600
10k, 100k
450–600
10k, 100k
100, 1k
x
480–650
10k, 30k, 100k, (200k) 10k, 30k, 100k, (200k) 10k, 30k, 100k, 200k, 10k, 100k, 200k 250k
10k, 100k, 200k
115
MoV 12 3 GX 15 CrMo 5 X 15 CrMo 5 1 X 16 CrMo 5 1 X 16 CrMo 5 1 X 18 CrN 28 X 19 CrMoN bVN 11 1 X 20 CrMoV 11 1 X 20 CrMoWV 12 1 X 20 CrMoWV 121 X 22 CrMoV 121 GX 23 Cr MoV 12 1 GX 30 CrSi7 GX 40 CrSi 13
480–650
Specifications for creep-resistant steels: Europe
X 11 CrMo WVNb 9 1 1 X 11 CrWMo VNb 9 1 GX 12 CrMo VNbN 9 1 X 12 CrMo 5
460–600 450–650 200k 520–650
10k, 100k, 200k 10k, 100k, 10k, 100k,
WPNL2204
116
Product form
Castings EN 10295 1% strain creep strength
Creep-resistant steels EN 10302 Creep rupture strength
1% strain creep strength
Consumables EN 14532-2
Creep rupture strength
Creep rupture
Generic ECCC Creep rupture strength
Material grade
GX GX GX GX GX
40 CrSi 17 40 CrSi 24 40 CrSi 28 130 CrSi 29 160 CrSi 18
T range (°C)
Duration (h)
600–900 600 –900 600–900 600–900 600–900
10k 10k 10k 10k 10k
T range (°C)
Duration (h)
600–900 800–900
100, 1k 100, 1k
T range (°C)
Duration (h)
T range (°C)
Duration (h)
Duration or temperature in brackets ( ) represent strength values after extended extrapolation in stress. Duration or temperature with an asterisk * represent strength values after extended extrapolation in time. For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strength assessment is also available. EN 14532-2 does not report strength data but does report assessed equations relating strength, duration and temperature. In the EN 14532-2 column, the remark ‘x notch’ means that an equation for creep notch rupture strength is also provided.
WPNL2204
T range (°C)
Duration (h)
Creep-resistant steels
Table 3.8 (Cont’d.)
Table 3.9 Chemical composition of austenitic steels contained in European Specifications Material grade
Gross chemical composition Concentration in mass% of the following elements Mo
Ni
Ti
Other
X 2 CrNi 18 9 X 5 CrNi 18 10 GX 5 CrNi 19 10 X 6 CrNi 18 10 X 6 CrNi 23 13 X 6 CrNi 25 20 X 8 CrNi 25 21 X 12 CrNi 23 13 X 2 CrNiN 18 10 X 2 CrNiMo 17 12 2 X 3 CrNiMo 17 13 3 X 5 CrNiMo 17 12 2 GX 5 CrNiNb 19 11 GX 5 CrNiMo 19 11 2 X 6 CrNiNb 18 10 A1000 X 6 CrNiNb 18 10 A1100 X 6 CrNiTi 18 10 A1000 X 6 CrNiTi 18 10 A1100 X 6 CrNiMo 17 13 2 X 7 CrNiNb 18 10 X 7 CrNiTi 18 10
<0.03 <0.07 <0.07 0.04–0.08 0.04–0.08 0.04–0.08 <0.1 <0.15 < 0.03 < 0.03 <0.05 <0.07 <0.07 <0.07 <0.08 <0.08 <0.08 <0.08 0.04–0.08 0.04–0.10 0.04–0.08
17.5–19.5 17.5–19.5 18.0–20.0 17.0–19.0 22.0–24.0 24.0–26.0 24.0–26.0 22.0–24.0 17.5–19.5 16.5–18.5 16.5–18.5 16.5–18.5 18.0–20.0 18.0–20.0 17.0–19.0 17.0–19.0 17.0–19.0 17.0–19.0 16.5–18.5 17.0–19.0 17.0–19.0
– – – – – – – – – 2.0–2.5 2.5–3.0 2.0–2.5 – 2.0–2.5 – – – – 2.0–2.5 – –
8.0–10.0 8.0–10.5 8.0–11.0 8.0–11.0 12.0–15.0 19.0–22.0 19.0–22.0 12.0–14.0 8.5–11.5 10.0–13.0 10.5–13.0 10.0–13.0 9.0–12.0 9.0–12.0 9.0–12.0 9.0–12.0 9.0–12.0 9.0–12.0 12.0–14.0 9.0–12.0 9.0–13.0
– – – – – – – – –
N<0.11 N<0.11
X 8 CrNiTi 18 10 X 8 CrNiNb 19 11 X 8 CrNiNb 16 13
<0.1 0.06–0.1 0.04–0.1
17.0–19.0 17.0–20.0 15.0–17.0
– – –
9.0–12.0 9. 0– 13 12.0–14.0
WPNL2204
– – – – – – 5*C–0.7 5*C–0.7 – – 5*(C+N)– 0.8 5*C–0.8 – –
N 0.12–0.22 N <0.11 N <0.11 N <0.11 Nb 8*C–1.0 Nb 10*C–1.0 Nb 10*C–1.0
– Nb 10*C–1.2 N <0.11
Si <0.75; Nb 8*C–1 Nb 10*C–1.2
117
Cr
Specifications for creep-resistant steels: Europe
C
118
Table 3.9 (Cont’d.) Material grade
Gross chemical composition
C
Cr
Mo
Ni
Ti
Other
X 12 CrCoNi 21 20 A
0.08–0.16
20–22.5
2.5–3.5
19.0–21
–
X 12 CrCoNi 21 20 P
0.08–0.16
20–22.5
2.5–3.5
19.0–21
–
X 15 CrNiSi 20 12 X 15 CrNiSi 25 21 X 15 CrNiSi 25 4 A GX 25 CrNiSi 18 9 GX 25 CrNiSi 20 14 GX 40 CrNiSi 27 4 GX 40 CrNiSi 22 10 GX 40 CrNiSi 25 12 GX 40 CrNiSi 25 20 X 2 CrNiMoN 17 11 2 X 2 CrNiMoN 17 13 3 X 3 CrNiMoN 17 13 3 X 5 CrNiNbN 18 10 X 5 CrNiMoB 17 13 3 X 6 CrNiTiB 18 10 X 6 CrNiMoTi 17 12 2 X 6 CrNiMoB 17 12 2 X 8 CrNiMoNb 16 16 X 8 CrNiNbN 25 21 X 10 CrNiCuNb 18 10
<0.2 <0.2 0.1–0.2 0.15–0.35 0.15–0.35 0.3–0.5 0.3–0.5 0.3–0.5 0.3–0.5 <0.03 <0.03 <0.04
19–21 24.0–26.0 24.5–26.5 17–19 19–21 25–28 21–23 24–27 24–27 16.5–18.5 16.5–18.5 16.0–18.0
– – – <0.5 <0.5 <0.5 <0.5 <0.5 <0.5 2.0–2.5 2.5–3.0 2.0–3.0
11.0–13.0 19.0–22.0 3.5–5.5 8.0–10.0 13–15 3.0–6.0 9.0–11.0 11.0–14.0 19.0–22.0 10.0–12.0 11.0–14.0 12.0–14.0
– – – – – – – – – – – –
Nb 0.75–1.25; W 2–3. Co 18.5–21 Nb 0.75–1.25; W 2–3. Co 18.5–21 Si 1.5–2.5; Si 1.5–2.5; Si 0.8–1.5 Si 0.5–2.5 Si 0.5–2.5 Si 1–2.5 Si 1–2.5 Si 1–2.5 Si 1–2.5 N 0.12–0.22 N 0.12–0.22 N 0.1–0.18; B 0.0015–0.005
0.04–0.08 <0.08 0.04–0.08 0.04–0.1 <0.1 0.07–0.13
17.0–19.0 16.5–18.5 16.5–18.5 15.5–17.5 23.0–27.0 17–19
– 2.0–2.5 2.0–2.5 1.6–2.0 – –
9.0–12.0 10.5–13.5 10.0–13.0 15.5–17.5 17.0–23.0 7.5–10.5
5*C–0.8 5*C– 0.7 – – – –
WPNL2204
B 0.0015–0.005 B 0.0015–0.005 Nb 10*C–1.2 Nb 0.2–0.6; N 0.15–0.35 Cu 2.5–3.5; Nb 0.3–0.6; N 0.05–0.12
Creep-resistant steels
Concentration in mass% of the following elements
15.5–17.5 24–26 23–25 16.0–18.0 15.5–17.5
– – <0.5 2.0–3.0 –
12.5–14.5 6.0–8.0 23.0– 25.0 12.0–14.0 15.5–17.5
0.4–0.7 – – – –
X 6 CrNiWNbN 16 16 WW X 6 CrNiMoTiB 17 13 A X 6 CrNiSiNCe 19 10
0.04–0.1
15.5–17.5
–
15.5–17.5
–
0.04–0.08 0.04–0.08
16–18 18–20
2–2.5 –
12.0–14.0 9.0– 11.0
5*C–0.8 –
X 7 CrNiMoBNb 16 16 X 8 CrNiMoVNb 16 13 X 9 CrNiSiNCe 21 11 2
0.04–0.1 0.04–0.1 0.05–0.12
15.5–17.5 15.5–17.5 20–22
1.6–2.0 1.1–1.5 –
15.5–17.5 12.5–14.5 10.0–12.0
– – –
0.07–0.15 0.07–0.13
15.5–17.5 14.0–16.0
– 0.8–1.2
12.5–14.5 9.0–11.0
0.4–0.7 –
X 10 NiCrSi 35 19 X 12 NiCrSi 35 16 GX 35 NiCrSi 25 21 GX 40 NiCrSi 35 17 GX 40 NiCrSi 38 19 GX 40 NiCrSi 35 26 GX 40 NiCrNb 45 35 GX 50 NiCrCo 20 20 20
<0.15 <0.15 0.2–0.5 0.3–0.5 0.3–0.5 0.3–0.5 0.35–0.45 0.35–0.65
17.0–20.0 15.0–17.0 19–23 16–18 18–21 24–27 32.5–37.5 19–22
– – <0.5 <0.5 <0.5 <0.5 – 2.5–3.0
33–37 33–37 23.0–27.0 34–36 36–39 33–36 42–46 18–22
– – – – – – – –
X 5 NiCrAlTi 31 20 X 5 NiCrAlTi 31 20 (RA)
0.03–0.08 0.03–0.08
19.0–22.0 19.0–22.0
– –
30.0–32.5 30.0–32.5
0.2–0.5 0.2–0.5
X 10 CrNiMnNbV 15 10 6 1 X 12 CrNiWTiB 16 13 WW X 10 CrNiMoMnNbV B 15 10 1
WPNL2204
W 2.5–3; B 0.0015–0.006 Mn 8–10; N 0.2–0.4 Si 1–2.5. Nb 0.8–1.8 N 0.1–0.18; B 0.0015–0.005 Nb 10*C–1.2; W 2.5–3.5; N 0.06–0.14 Nb 10*C–1.2; W 2.5–3.5; N 0.06–0.14 B 0.0015–0.006 Si 1–2; N 0.12–0.2; Ce 0.03–0.08 B 0.05–0.1; Nb+Ta 10*C–1.2 Nb 10*C–1.2; V 0.6–0.85 Si 1.4–2.5;. N 0.12–0.2; Ce 0.03–0.08
W 2.5–3; B 0.0015–0.006 Mn 5.5–7; B 0.003–0.009; V 0.15–0.4; Nb 0.75–1.25; N<0.11 Si 1–2; Si 1–2; Si 1–2.0 Si 1–2.5 Si 1–2.5 Si 1–2.5 Si 1.5–2. Nb 1.5–2 Si<1:. Nb 0.75–1.25; Co 18.5–22, W 2–3 Al 0.2–0.5; Cu <0.5; Nb<0.1 Al 0.2–0.5; Cu <0.5; Nb<0.1
119
0.07–0.15 0.2–0.3 0.3–0.5 <0.04 0.04–0.1
Specifications for creep-resistant steels: Europe
X 12 CrNiWTiB 16 13 A X 25 CrMnNiN 25 9 7 GX 40 CrNiSiNb 24 24 X 3 CrNiMoBN 17 13 3 X 6 CrNiWNbN 16 16
120
Table 3.9 (Cont’d.) Material grade
Gross chemical composition
C
Cr
Mo
Ni
Ti
Other
X 6 NiCrNbCe 32 27 X 8 NiCrAlTi 32 21
0.04–0.08 0.05–0.1
26–28 19.0–22.0
– –
31–33 30.0–34.0
– 0.25–0.65
X 8 NiCrAlTi 32 21 (RK)
0.05–0.1
19.0–22.0
–
30.0–34.0
0.25–0.65
X 10 NiCrAlTi 32 21 X 10 NiCrSiNb 35 22 GX 10 NiCrSiNb 32 20 GX 40 NiCrSiNb 35 18 GX 40 NiCrSiNb 38 19 GX 40 NiCrSiNb 35 26 GX 50 NiCrCoW 35 25 15 5 X 6 NiCrSiNCe 35 25
<0.12 <0.15 0.05–0.15 0.3–0.5 0.3–0.5 0.3–0.5 0.45–0.55
19.0–23.0 20–23 19–21 17–20 18–21 24–27 24–26
– – <0.5 <0.5 <0.5 <0.5 –
30–34 33–37 31–33 34–36 36–39 33–36 33–37
0.15–0.6 – – – – – –
Ce 0.05–0.1; Nb 0.6–1 Al 025–0.65; Cu<0.5; Co<0.5; Al+Ti<0.5 Al 0.25–0.65; Cu<0.5; Co<0.5; Al+Ti 0.85–1.2 Al 0.15–0.6 Si 1–2; Nb 1–1.5 Si 0.5–1.5; Nb 0.5–1.5 Si 1–2.5; Nb 1–1.8 Si 1–2.5; Nb 1.2–1.8 Si 1–2.5; Nb 0.8–1.8 Si 1–2. Co 14–16; W 4–6
0.04–0.08
24–26
–
34–36
–
X 2 NiCrMoNbBN 25 22
<0.04
21.5–23
1.0–2.0
22–28
<0.2
X 6 NiCrTiMoVB 25 15 2 X 8 NiCrMoNbBN 25 20
0.03–0.08 0.05–0.1
13.5–16.0 19–21
1.0–1.5 1.0–2.0
24.0–27.0 22–28
1.9–2.3 <0.2
WPNL2204
Si 1.2–2; N 0.12–0.2; Ce 0.03–0.08 Nb 0.1–0.4; B 0.002–0.01; N 0.1–0.25 B 0.003–0.01; V 0.1–0.5 Nb 0.1–0.4; B 0.002–0.01; N 0.1–0.25
Creep-resistant steels
Concentration in mass% of the following elements
Table 3.10 Available creep data for austenitic steels: plates; heat-resistant steels and Ni alloys; castings; seamless tubes/pipes Product form
Plates EN 10028-7 1% strain creep strength
X 2 CrNi 18 9 X 5 CrNi 18 10 GX 5 CrNi 19 10 X 6 CrNi 18 10
X 6 CrNi 23 13 X 6 CrNi 25 20
Creep rupture strength
T range (°C)
Duration (h)
T range (°C)
Duration (h)
500–750
10k 100k
500–750
550–800
10k 100k
550–800 600–910
10k, 30k, 50k, 100k, 200k 10k, 100k 10k, 30k, 50k, 100k, 150k, 200k, 250k
1% strain creep strength
T range (°C)
Duration (h)
Creep rupture strength
T range (°C)
600–900
1k 10k
600–900
X 12 CrNi 23 13 X 2 CrNiN 18 10 X 2 CrNiMo 17 12 2 X 3 CrNiMo 17 13 3
600–900
1k 10k
600–900
1k, 10k, 100k 1k, 10k, 100k
Seamless tubes/pipes EN 10216-5
Creep rupture strength
Creep rupture strength
T range (°C)
Duration (h)
550–700
10k, 100k
T range (°C)
Duration (h)
500–700
10k, 100k,
(750)
200k
121
X 8 CrNi 25 21
Duration (h)
Castings EN 10213-4
Specifications for creep-resistant steels: Europe
Material grade
Heat-resisting steels and Ni-alloys EN 10095
WPNL2204
122
Table 3.10 (Cont’d.) Product form
1% strain creep strength
Heat-resisting steels and Ni-alloys EN 10095
Creep rupture strength
1% strain creep strength
Creep rupture strength
Castings EN 10213-4
Seamless tubes/pipes EN 10216-5
Creep rupture strength
Creep rupture strength
Material grade
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration (h)
X 5 CrNiMo 17 12 2 GX 5 CrNiNb 19 11 GX 5 CrNiMo 19 11 2 X 6 CrNiNb 18 10 A1000 X 6 CrNiNb 18 10 A1100 X 6 CrNiTi 18 10 A1000 X 6 CrNiTi 18 10 A1100 X 6 CrNiMo 17 13 2 X 7 CrNiNb 18 10 X 7 CrNiTi 18 10
T range (°C)
Duration (h)
T range (°C)
Duration (h)
550–700
10k, 100k
550–700
10k, 100k
T range (°C)
550–700
10k, 100k
540–700
10k, 100k, 200k* 10k, 100k
550–800
WPNL2204
Duration (h)
Creep-resistant steels
Plates EN 10028-7
X 8 CrNiTi 18 10
580–750
10k, 100k
580–750
1k 10k
600–800
1k,10k, 100k
10k, 100k, 200k
580–750
600–800
1k 10k
600–900
600–1000 1k 10k
600– 1000 500–900
500–900
1k 10k
1k, 10k, 100k 1k, 10k, 100k 1k 10k
10k, 100k, 200k
Specifications for creep-resistant steels: Europe 123
X 8 CrNiNb 19 11 X 8 CrNiNb 16 13 X 12 CrCoNi 21 20 A X 12 CrCoNi 21 20 P X 15 CrNiSi 20 12 X 15 CrNiSi 25 21 X 15 CrNiSi 25 4 A GX 25 CrNiSi 18 9 GX 25 CrNiSi 20 14 GX 40 CrNiSi 27 4 GX 40 CrNiSi 22 10 GX 40 CrNiSi 25 12 GX 40 CrNiSi 25 20 X 2 CrNiMoN 17 11 2 X 2 CrNiMoN 17 13 3
600–800
WPNL2204
124
Table 3.10 (Cont’d.) Product form
1% strain creep strength
Heat-resisting steels and Ni-alloys EN 10095
Creep rupture strength
1% strain creep strength
Creep rupture strength
Castings EN 10213-4
Seamless tubes/pipes EN 10216-5
Creep rupture strength
Creep rupture strength
Material grade
T range (°C) X 3 CrNiMoN 17 13 3 X 5 CrNiNbN 18 10 X 5 CrNiMoB 17 13 3 X 6 CrNiTiB 18 10 X 6 CrNiMoTi 17 12 2 X 6 CrNiMoB 17 12 2 X 8 CrNiMo Nb 16 16 X 8 CrNiNbN 25 21 X 10 CrNiCuNb 18 10 X 12 CrNiWTiB 16 13 A X 25 CrMnNiN 25 9 7
Duration (h)
T range (°C)
Duration (h)
550–700
10k, 100k, 200k
T range (°C)
700–900
Duration (h)
1k, 10k
WPNL2204
T range (°C)
700–900
Duration (h)
1k, 10k
T range (°C)
Duration (h)
T range (°C)
Duration (h)
550–700
10k, 100k, 200k
580–750
10k, 100k, 200k
Creep-resistant steels
Plates EN 10028-7
550–800
10k, 100k, 200k
600– (1000)
600– (1000)
600– (1000) 600–900
1k, 10k, 100k
1k, 10k, 100k
600– (1000)
600– (1000)
550–800
10k, 100k, (200k)
580–650
10k, 100k 200k
600–780
10k, 100k 200k, 250k*
1k, 10k, 100k
1k, 10k, 100k
1k 10k
600–900
1k, 10k
1k 10k 100k
600– 1000
1k, 10k, 100k
Specifications for creep-resistant steels: Europe 125
GX 40 CrNiSi Nb 24 24 X 3 CrNiMoBN 17 13 3 X 6 CrNiWNbN 16 16 X 6 CrNiWNbN 16 16 WW X 6 CrNiMoTiB 17 13 A X 6 CrNiSiNCe 19 10 X 7 CrNiMo BNb 16 16 X 8 CrNiMoV Nb 16 13 X 9 CrNiSiNCe 21 11 2 X 10 CrNiMn NbV 15 10 6 1 X 12 CrNiWTiB 16 13 WW X 10 CrNiMoMn NbVB 15 10 1 X 10 NiCr Si 35 19 X 12 NiCr Si 35 16 GX 35 NiCr Si 25 21 GX 40 NiCr Si 35 17
WPNL2204
126
Table 3.10 (Cont’d.) Product form
1% strain creep strength
Heat-resisting steels and Ni-alloys EN 10095
Creep rupture strength
1% strain creep strength
Creep rupture strength
Castings EN 10213-4
Seamless tubes/pipes EN 10216-5
Creep rupture strength
Creep rupture strength
Material grade
T range (°C) GX 40 NiCr Si 38 19 GX 40 NiCr Si 35 26 GX 40 NiCr Nb 45 35 GX 50 NiCrCo 20 20 20 X 5 NiCr AlTi 31 20 X 5 NiCrAlTi 31 20 (RA) X 6 NiCrNb Ce 32 27 X 8 NiCr AlTi 32 21 X 8 NiCrAlTi 32 21 (RK) X 10 NiCr AlTi 32 21 X 10 NiCr SiNb 35 22
600–700 550–700
Duration (h)
T range (°C)
Duration (h)
10k (100k) 10k (100k)
500–700
10k, 100k, (200k) 10k, 100k, (200k)
500–700
T range (°C)
Duration (h)
T range (°C)
10k 100k
700– 1000
T range (°C)
Duration (h)
T range (°C)
500–700 500–700 800– (1000)
700– 1000
Duration (h)
700– 1000
1k 10k
600 –900
600– (1000)
1k 10k
600–900
WPNL2204
10k, 100k (200k) 10k, 100k (200k)
10k, 100k
10k, 30k 100k, 200k
600–900
Duration (h)
1k, 10k, 100k 1k, 10k
10k, 100k (200k)
Creep-resistant steels
Plates EN 10028-7
600– 1000
1k, 10k 100k
600– 1000
1k, 10k 100k
Specifications for creep-resistant steels: Europe
GX 10 NiCr SiNb 32 20 GX 40 NiCr SiNb 35 18 GX 40 NiCr SiNb 38 19 GX 40 NiCr SiNb 35 26 GX 50 NiCrCo W 35 25 15 5 X 6 NiCrS iNCe 35 25 X 2 NiCrMo NbBN 25 22 X 6 NiCrTiMoV B 25 15 2 X 8 NiCrMoNb BN 25 20
127
WPNL2204
128
Table 3.11 Available creep data for austenitic steels: forgings; bars for fasteners and bolts; castings; creep-resistant steels Product form
Creep rupture strength
Bars for fasteners and bolting EN 10269 1% strain creep strength
Castings EN 10295
Creep rupture strength
1% strain creep strength
Creep-resistant steels EN 10302
Creep rupture strength
1% strain creep strength
Creep rupture strength
Material grade
X 2 CrNi 18 9 X 5 CrNi 18 10 GX 5 CrNi 19 1 X 6 CrNi 18 10 X 6 CrNi 23 13 X 6 CrNi 25 20 X 8 CrNi 25 21 X 12 CrNi 23 13 X 2 CrNiN 18 10 X 2 CrNiMo 17 12 2 X 3 CrNiMo 17 13 3 X 5 CrNiMo 17 12 2 GX 5 CrNiNb 19 11 GX 5 CrNiMo 19 11 2 X 6 CrNiNb 18 10 A1000 X 6 CrNiNb 18 10 A1100
T range (°C)
Duration (h)
550–700 550–700
10k, 100k 10k, 100k
550–700
10k, 100k, 200k
540–700
10k, 100k, 200k 10k, 100k, 200k
540–700
540–700 540–700
T range (°C)
Duration (h)
T range (°C)
Duration (h)
550–700
10k, 100k
550–700
10k, 100k, 200k
T range (°C)
10k, 100k, 200k 10k, 100k, 200k
WPNL2204
Duration T range (h) (°C)
Duration (h)
T range (°C)
Duration T range (h) (°C)
Duration (h)
Creep-resistant steels
Forgings EN 10222-5
10k 100k 200k
10k 100k 200k
580–750 550–850 550–850
600–900 10k
10k
600–900
600– 1000 600–900
10k
600–900
100, 1k
700– 1000 700– 1100
10k
700– 1000 700– 1100
100, 1k
600–900
10k
100, 1k
580– 750 550– 900 550– 900
10k, 100k, 200k 10k, 100k, 200k 10k, 100k, 200k
129
10k 100k 200k 10k 100k 200k
10k
100, 1k
10k 100k 10k 100k 10k 100k
Specifications for creep-resistant steels: Europe
X 6 CrNiTi 18 10 A1000 X 6 CrNiTi 540–700 18 10 A1100 X 6 CrNiMo 17 13 2 X 7 CrNiNb 540–700 18 10 X 7 CrNiTi 18 10 X 8 CrNiTi 18 10 X 8 CrNiNb 19 11 X 8 CrNiNb 16 13 X 12 CrCoNi 21 20 A X 12 CrCoNi 21 20 P X 15 CrNiSi 20 12 X 15 CrNiSi 25 21 X 15 CrNiSi 25 4 A GX 25 CrNiSi 18 9 GX 25 CrNiSi 20 14 GX 40 CrNiSi 27 4 GX 40 CrNiSi 22 10 GX 40 CrNiSi 25 12 GX 40 CrNiSi 25 20 X 2 CrNiMoN 550–700 17 11 2 X 2 CrNiMoN 550–700 17 13 3
WPNL2204
130
Table 3.11 (Cont’d.) Product form
Creep rupture strength
Bars for fasteners and bolting EN 10269 1% strain creep strength
Castings EN 10295
Creep rupture strength
1% strain creep strength
Creep-resistant steels EN 10302
Creep rupture strength
1% strain creep strength
Creep rupture strength
Material grade
X 3 CrNiMoN 17 13 3 X 5 CrNiNbN 18 10 X 5 CrNiMoB 17 13 3 X 6 CrNiTiB 18 10 X 6 CrNiMo Ti 17 12 2 X 6 CrNiMo B 17 12 2 X 8 CrNiMo Nb 16 16 X 8 CrNiNb N 25 21 X 10 CrNiCu Nb 18 10 X 12 CrNiW TiB 16 13 A X 25 CrMn NiN 25 9 7 GX 40 Cr NiSiNb 24 24
T range (°C)
Duration (h)
550–800
10k, 100k, 200k
550–700
10k, 100k, 200k 10k, 100k, 200k
540–700
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
700– 1000
WPNL2204
Duration T range (h) (°C)
10k
700– 1100
Duration (h)
100, 1k
T range (°C)
Duration T range (h) (°C)
Duration (h)
550–700
10k, 100k, 200k
550–850
10k, 100k, 200k 10k, 100k, 200k
580–750
10k, 100k
580–750
600–750
10k, 100k
600–750
10k, 100k
Creep-resistant steels
Forgings EN 10222-5
X 10 CrNiMo
10k, 100k, 200k
550–800 580–750 (550)– 650 600–700
580–670
10k, 100k
580–670
550–650
10k, 100k,
10k, 580–750 100k 10k, 100k (550)– 650 10k, 100k 600–700
200k 10k, 100k, 200k 10k, 100k 10k, 100k
10k, 100k
10k, 100k, 200k
700– 1000 700– 1000 700– 1100 700– 1000 1000
WPNL2204
10k
800–1100 1k
10k
700–1100 100 1k
10k
700–1100 100 1k
580–650
10k, 100k 580– 650
10k, 100k, 200k
500–700
10k, 100k, 200k
10k, 100k, 200k
500– 700
10k 10k
900–1100 100 1k
131
MnNbVB 15 10 1 X 10 NiCrSi 35 19 X 12 NiCrSi 35 16 GX 35 NiCr Si 25 21 GX 40 NiCr Si 35 17 GX 40 NiCr Si 38 19 GX 40 NiCr Si 35 26 GX 40 NiCr Nb 45 35
550–800
Specifications for creep-resistant steels: Europe
X 3 CrNiMo BN 17 13 3 X 6 CrNiWN bN 16 16 X 6 CrNiWNbN 16 16 WW X 6 CrNiMoTiB 17 13 A X 6 CrNiSi NCe 19 10 X 7 CrNiMo BNb 16 16 X 8 CrNiMo VNb 16 13 X 9 CrNiSi NCe 21 11 2 X 10 CrNiMn NbV 15 10 6 1 X 12 CrNiWTiB 16 13 WW
132
Table 3.11 (Cont’d.) Product form
Creep rupture strength
Bars for fasteners and bolting EN 10269 1% strain creep strength
Castings EN 10295
Creep rupture strength
1% strain creep strength
Creep-resistant steels EN 10302
Creep rupture strength
1% strain creep strength
Creep rupture strength
Material grade
T range (°C) GX 50 NiCr Co 20 20 20 X 5 NiCr AlTi 31 20 X 5 NiCr AlTi 31 20 (RA) X 6 NiCr NbCe 32 27 X 8 NiCr AlTi 32 21 X 8 NiCr AlTi 32 21 (RK) X 10 NiCr AlTi 32 21 X 10 NiCr SiNb 35 22 GX 10 NiCr SiNb 32 20 GX 40 NiCr8 SiNb 35 1 GX 40 NiCr SiNb 38 19 GX 40 NiCr SiNb 35 26
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration (h)
T range (°C)
Duration T range (h) (°C)
900– 1100
10k
Duration (h)
T range (°C)
Duration T range (h) (°C)
Duration (h)
600–700
10k, (100k) 10k, (100k)
10k, 100k, (200k) 10k, 100k, (200k)
800–1000 100 1k
550–700
700– 1000
700– 1000
700– 1000 700– 1100
WPNL2204
10k
10k 10k
700– 1000 700– 1100 700–900 700–1 1100
100 1k 100 1k 100 1k 100 1k
500–700 500–700
10k, 100k 700– 1000
10k, 100k, (200k)
Creep-resistant steels
Forgings EN 10222-5
1000– 1100
500–650
10k, 100k
500–650
10k, 100k
10k
500–650
10k, 100k 500– 650
10k, 100k
Specifications for creep-resistant steels: Europe
GX 50 NiCrCo W 35 25 15 5 X 6 NiCrSi NCe 35 25 X 2 NiCrMo NbBN 25 22 X 6 NiCrTiMo VB 25 15 2 X 8 NiCrMo NbBN 25 20
133
WPNL2204
134
Creep-resistant steels
Table 3.12 Available creep data for austenitic steels: consumables; generic products Product form
Consumables
Generic
EN 14532-2
ECCC
Material grade Creep rupture
X 2 CrNi 18 9 X 5 CrNi 18 10 GX 5 CrNi 19 10 X 6 CrNi 18 10 X 6 CrNi 23 13 X 6 CrNi 25 20 X 8 CrNi 25 21 X 12 CrNi 23 13 X 2 CrNiN 18 10 X 2 CrNiMo 17 12 2 X 3 CrNiMo 17 13 3 X 5 CrNiMo 17 12 2 GX 5 CrNiNb 19 11 GX 5 CrNiMo 19 11 2 X 6 CrNiNb 18 10 A1000 X 6 CrNiNb 18 10 A1100
T range (°C)
Duration (h)
550–720
(10k)–(100k)
500–(750)
10k, 100k, 200k
600–910
10k, 30k, 50k, 100k, 150k, 200k, 250k
580–800 500–700
10k, 100k, 200k 10k, 30k, 100k, 200k
x
500–850
10k, 30k, 100k, 200k
x
540–710
x
540–730
10k 30k, 50k, 100k, 150k, 200k, 250k 10k 30k, 50k, 100k, 150k, 200k, 250k 10k 100k 200k 10k 30k, 50k, 100k, 150k, 200k, 250k
x
X 6 CrNiTi 18 10 A1000 X 6 CrNiTi 18 10 A1100 X 6 CrNiMo 17 13 2 X 7 CrNiNb 18 10 X 7 CrNiTi 18 10 X 8 CrNiTi 18 10 X 8 CrNiNb 19 11 X 8 CrNiNb 16 13 X 12 CrCoNi 21 20 A X 12 CrCoNi 21 20 P X 15 CrNiSi 20 12 X 15 CrNiSi 25 21 X 15 CrNiSi 25 4 A GX 25 CrNiSi 18 9 GX 25 CrNiSi 20 14 GX 40 CrNiSi 27 4 GX 40 CrNiSi 22 10 GX 40 CrNiSi 25 12 GX 40 CrNiSi 25 20 X 2 CrNiMoN 17 11 2 X 2 CrNiMoN 17 13 3 X 3 CrNiMoN 17 13 3 X 5 CrNiNbN 18 10 X 5 CrNiMoB 17 13 3 X 6 CrNiTiB 18 10
Creep rupture strength
550–700 540–720
x (A1000) x (A1100)
x x x
WPNL2204
600–750 580–750
10k, 100k 10k, 100k, 200k (1%)
550–750
10k, 100k, 200k
Specifications for creep-resistant steels: Europe X 6 CrNiMoTi 17 12 2 X 6 CrNiMoB 17 12 2 X 8 CrNiMoNb 16 16 X 8 CrNiNbN 25 21 X 10 CrNiCuNb 18 10 X 12 CrNiWTiB 16 13 A X 25 CrMnNiN 25 9 7 GX 40 CrNiSiNb 24 24 X 3 CrNiMoBN 17 13 3 X 6 CrNiWNbN 16 16 X 6 CrNiWNbN 16 16 WW X 6 CrNiMoTiB 17 13 A X 6 CrNiSiNCe 19 10 X 7 CrNiMoBNb 16 16 X 8 CrNiMoVNb 16 13 X 9 CrNiSiNCe 21 11 2 X 10 CrNiMnNbV 15 10 6 1 x X 12 CrNiWTiB 16 13 WW X 10 CrNiMoMnNbVB 15 10 1 x X 10 NiCrSi 35 19 X 12 NiCrSi 35 16 GX 35 NiCrSi 25 21 GX 40 NiCrSi 35 17 GX 40 NiCrSi 38 19 GX 40 NiCrSi 35 26 GX 40 NiCrNb 45 35 GX 50 NiCrCo 20 20 20 X 5 NiCrAlTi 31 20 x X 5 NiCrAlTi 31 20 (RA) X 6 NiCrNbCe 32 27 X 8 NiCrAlTi 32 21
X 8 NiCrAlTi 32 21 (RK)
X 10 NiCrAlTi 32 21 X 10 NiCrSiNb 35 22 GX 10 NiCrSiNb 32 20 GX 40 NiCrSiNb 35 18 GX 40 NiCrSiNb 38 19 GX 40 NiCrSiNb 35 26 GX 50 NiCrCoW 35 25 15 5 X 6 NiCrSiNCe 35 25 X 2 NiCrMoNbBN 25 22 X 6 NiCrTiMoVB 25 15 2 X 8 NiCrMoNbBN 25 20
135
580–820
10k, 30k, 100k, 200k
580–770 600–750
10k, 30k, 100k 10k, 100k
550–800
10k, 100k, (200k)
550–1100
10k, 20k, 100k, 200k
600–790
10k, 30k, 100k, 200k
500–700 500–700 580–950 550–1000
10k, 100k, (200k) 10k, 100k (200k) (10k–100k) 10k, 20k, 30k, 50k, 70k, 100k, 150k, 200k, 250k, 300k 10k, 20k, 30k, 50k, 70k, 100k, 150k, 200k, 250k, 300k (1%)
600–1000
600–820
10k, 30k, 100k
580–770
10k, 30k, 100k
Duration or temperature in brackets ( ) represent strength values after extended extrapolation in stress. Duration or temperature with an asterisk * represent strength values after extended extrapolation in time. For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strength assessment is also available. EN 14532-2 does not report strength data but does report assessed equations relating strength, duration and temperature. In the EN 14532-2 column, the remark ‘x notch’, means that an equation for creep notch rupture strength is also provided.
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Table 3.13 Chemical analysis of nickel and cobalt base alloys contained in European Standards Gross chemical composition Heat
Concentration in mass% of the following elements
treatment
C
Cr
Co
Fe
Mo
Ni
G NiCr15 G NiCr28W G NiCr50Nb NiCo20Cr20MoTi
As cast As cast As cast P
0.35–0.65 0.35–0.55 <0.1 0.04–0.08
12.0–18.0 27–30 48–52 19–21
– – – 19–21
Balance Balance <1 <0.7
<1 <0.5 <0.5 5.6–6.1
58–66 47–50 Balance Balance
NiCr15Fe NiCr15Fe7TiAl
A AT P
0.05–0.1 <0.08
14–17 14–17
<1.5 <1
6.0–10.0 5.0–9.0
– –
Balance Balance
NiCr15Fe7TiAl
P 980
<0.08
14–17
<1
5.0–9.0
–
Balance
NiCr15Fe7TiAl
P 1170
<0.08
14–17
<1
5.0–9.0
–
Balance
NiCr19Fe19Nb5Mo3
P
0.02–0.08
17–21
<1
Balance
2.8–3.3
50–55
NiCr20Co13Mo4 Ti3Al
P
0.02–0.1
18–21
12.0–15
<2
3.5–5
Balance
NiCr20Co18Ti NiCr20Ti NiCr20TiAl
P AT AT P 2 stage
<0.13 0.08–0.15 0.04–0.1
18–21 18–21 18–21
15–21 <5 <1
<1.5 <5 <1.5
– – –
Balance Balance Balance
WPNL2204
Other
W 4–6 Nb 1–1.8. N<0.16 Al 0.3–0.6; Ti 1.9–2.4; B Al. Ti Al 0.4–1; Ti 2.25–2.75; Nb+Ta 0.7–1.2; Al 0.4–1; Ti 2.25–2.75; Nb+Ta 0.7–1.2; Al 0.4–1; Ti 2.25–2.75; Nb+Ta 0.7–1.2; Al 0.3–0.7; Nb+Ta 4.7–5.5; Ti 0.6–1.2; B 0.002–0.006 Al 1.2–1.6; Ti 2.8–3.3; B; Zr; Al 1–2; Ti 2–3; Zr; B Al; Ti 0.2–0.6; Al 1–1.8; B <0.008. Ti 1.8–2.7
Creep-resistant steels
Material grade
P 3 stage
0.04–0.1
18–21
<1
<1.5
–
Balance
NiCr22Fe18Mo NiCr22Mo9Nb
AT A
0.05–0.15 0.03–0.1
20.5–23 20–23
0.5–2.5 <1
17–20 <5
8.0– 10 8.0–10.0
Balance Balance
NiCr23Co12Mo
AT
0.05–0.1
20–23
11.0–14
<2
8.5–10
Balance
NiCr23Fe
AT
0.03–0.1
21–25
<1.5
<18
–
Balance
NiCr25Co20TiMo
P 1100
0.03–0.07
23–25
19–21
<1
1–2.0
Balance
NiCr25FeAlY
AT
0.15–0.25
24–26
–
8.0–11
–
Balance
NiCr26MoW NiCr28FeSiCe
AT AT
0.03–0.08 0.05–0.12
24–26 26–29
2.5–4 <1.5
Balance 21–25
2.5–4 –
44–47 Balance
NiCr29Fe CoCr20W15Ni G CoCr28
AT AT As cast
<0.05 0.05–0.15 0.05–0.25
27–31 19–21 27–30
– Balance 48–52
7.0–11 <3 Balance
– – <0.5
Balance 9.0–11 <4
Al 1–1.8; B <0.008. Ti 1.8–2.7 Al <0.5; B; W 0.2–1 Al <0.4; Nb+Ta 3.15–4.15 Al 0.7–1.4; Ti 02–0.6; B Al 1–1.7; Ti <0.4; B<0.006 Al 1.2–1.6; Nb+Ta 0.7–1.2; Ti 2.8–3.2; B 0.01–0.015; Zr 0.03–0.07; Ta Al 1.8–2.4; Ti 0.1–0.2; Y 0.05–0.12; Zr W 2.5–4 Ti <0.5; Si 2.5–3; Ce 0.03–0.09 Al <0.5 W 14–16 Nb <0.5
Specifications for creep-resistant steels: Europe
NiCr20TiAl
137
WPNL2204
138
Table 3.14 Available creep data for nickel and cobalt base alloys: heat-resistant steels and Ni alloys; bars for fasteners and bolts; castings Product form
Heat treatment
1% strain creep strength
Creep rupture strength
Bars for fasteners and bolting EN 10269 1% strain creep strength
Creep rupture strength
Castings EN 10295 Creep rupture strength
Creep rupture strength
Material grade
T range (°C) G NiCr15 G NiCr28W G NiCr50Nb NiCo20Cr 20MoTi NiCr15Fe
As cast As cast As cast P
NiCr15Fe7TiAl NiCr15Fe7TiAl NiCr15Fe7TiAl NiCr19Fe19 Nb5Mo3 NiCr20Co 13Mo4Ti3Al NiCr20Co18Ti NiCr20Ti
AT P P 980 P 1170 P
NiCr20TiAl
AT P 2 stage P 3 stage
NiCr20TiAl
A
Duration T range (h) (°C)
Duration T range (h) (°C)
Duration T range (h) (°C)
Duration T range (h) (°C)
700–1100 10k 700–1000 10k
500–900
10k, 100k 600–1000 1k, 10k, 100k 500–800
10k 100k 500–800
10k 100k
500–800
10k 100k 500–800
10k 100k
P P AT
Duration T range (h) (°C)
600–1000 1k, 10k, 100k
WPNL2204
900–1000 800–1100 700–1100
Duration (h) 100. 1k 1k 100, 1k
Creep-resistant steels
Heat-resisting steels and Ni-alloys EN 10095
AT A AT AT
NiCr25Co 20TiMo NiCr25FeAlY NiCr26MoW NiCr28FeSiCe NiCr29Fe CoCr20W15Ni G CoCr28
P 1100 AT AT AT AT AT As cast
700–900
1k, 10k
600–1000 1k, 10k, 100k
700–1000 10k, 100k 700–1000 10k, 100k
700–1100 10k
900–1000
100, 1k
Specifications for creep-resistant steels: Europe
NiCr22Fe18Mo NiCr22Mo9Nb NiCr23Co12Mo NiCr23Fe
139
WPNL2204
140
Table 3.15 Available creep data for nickel and cobalt base alloys: creep-resistant steels; generic products Product form
Heat treatment
G NiCr15 G NiCr28W G NiCr50Nb NiCo20Cr20MoTi NiCr15Fe NiCr15Fe7TiAl NiCr15Fe7TiAl NiCr15Fe7TiAl NiCr19Fe19Nb5Mo3 NiCr20Co13Mo4Ti3Al NiCr20Co18Ti NiCr20Ti NiCr20TiAl NiCr20TiAl NiCr22Fe18Mo NiCr22Mo9Nb NiCr23Co12Mo NiCr23Fe NiCr25Co20TiMo NiCr25FeAlY NiCr26MoW NiCr28FeSiCe
As cast As cast As cast P A AT P P 980 P 1170 P P P AT AT P 2 stage P 3 stage AT A AT AT P 1100 AT AT AT
1% strain creep strength
Generic ECCC
Creep rupture strength
T range (°C)
Duration (h)
T range (°C)
Duration (h)
500–900
10k, 100k
500–900
10k, 100k
650–800 500–800 500–800
10k, 100k 10k, 100k 10k, 100k
700–800
10k, 100k
650–800 500–800 500–800 650–800 550–800
10k, 10k, 10k, 10k, 10k,
500–800
10k, 100k
500–800
10k, 100k
600–1000
10k, 100k
550–1000
10k, 100k
650–1000
10k, 100k
580–1000
10k, 100k
600–(900) 650–1200 600–1050
10k, 100k 10k, 100k 10k, (100k)
550–900 650–1200 600–1000
10k, 100k 10k, 100k 10k, (100k)
WPNL2204
Creep rupture strength
T range (°C)
Duration (h)
600–900
10k, 30k, 100k, 200k
450–670
10k, 30k, 100k, 200k
580–950
10k 30k 100k
100k 100k 100k 100k 100k
Creep-resistant steels
Material grade
Creep-resistant steels EN 10302
AT AT As cast
700–1050 700–1000
10k, 100k 10k, (100k)
700–1050 700–1000
10k, 100k 10k, (100k)
Duration or temperature in brackets ( ) represent strength values after extended extrapolation in stress. Duration or temperature with an asterisk * represent strength values after extended extrapolation in time. For steels assessed by ECCC the symbol (1%) means that a 1% strain creep strength assessment is also available. EN 14532-2 does not report strength data but does report assessed equations relating strength, duration and temperature. In the EN 14532-2 column, the remark ‘x notch’, means that an equation for creep notch rupture strength is also provided.
Specifications for creep-resistant steels: Europe
NiCr29Fe CoCr20W15Ni G CoCr28
141
WPNL2204
142
Creep-resistant steels
Although these marks provide obviously significant information, it is not clear how design should consider them with regard to safety factors, or how and if the marked strengths can be further extrapolated. The challenge to state guidelines on the handling of ‘different reliability’ strength values is a target of new design standards currently under development. In some cases, the standard also provides details about the scatter band of the data set used and the assessment method. If the standard data were assessed by ECCC, details of the data set, the collation specification (mechanical data, chemical composition, quantity and duration of tests, etc) and a straight equation for the chosen assessment (relating rupture time, temperature and strength) can be found in the ECCC Data Sheets.23,24 For bolting materials, stress relaxation may also become a design issue as well as the creep properties themselves. Therefore some relaxation properties have been collated and published, in EN 10269 and the ECCC Data Sheets.23,24 Table 3.16 gives an overview of available materials and the scope of relaxation strengths included.
3.5.3
Welding consumables and qualification
Welding consumables for creep-resistant steels are now classified, like all others, according to: • • • • •
EN 12070 (Welding consumables – Wire electrodes and wires and rods for arc welding of creep resisting steels – classification); EN 12072 (Welding consumables – Wire electrodes and wires and rods for arc welding of stainless and heat resisting steels – classification); EN 12536 (Welding consumables – Rods for gas welding of non-alloy and creep resisting steels – classification); EN 1599 (Welding consumables – Covered electrodes for manual arc welding of creep resisting steels – classification); or EN 1600 (Welding consumables – Covered electrodes for manual arc welding of stainless and heat resisting steels – classification).
The general welding procedure qualification comes under EN ISO 15613 or EN ISO 15614. The qualification of the consumables themselves is regulated by EN 14532 (Welding Consumables – Test methods and quality requirements) and a particular, creep-related ‘supplementary prescription’ is included in: EN 14532-2 (Welding Consumables – Test methods and quality requirements – Part 2: Supplementary methods and conformity assessment of consumables for steel, nickel and nickel alloys). Consumables intended for use in the creep regime, defined numerically by a temperature range for ferritic– martensitic, austenitic and nickel-base materials, shall be qualified for the maximum service temperature by comparison tests with a ‘sufficiently similar’ parent material: a stress rupture all weld material creep test series of at least
WPNL2204
Table 3.16 Relaxation strength in European Standards EN 10269
ECCC
Relaxation strength
Creep rupture strength
Heat treatment
C
Cr
Mo
Ni
V
42 CrMo 5 6 40 CrMoV 4 6 21CrMoV 5 7 20 CrMoVTiB 4 10
QT QT QT QT
0.39–0.45 0.26–0.44 0.17–0.25 0.17–0.23
1.2–1.5 0.9–1.2 1.2–1.5 0.9–1.2
0.5–0.7 0.5–0.65 0.55–0.8 0.9–1.1
– – <0.6 <0.2
– 0.25–0.35 0.2–0.35 0.6–0.8
X 22 CrMoV 12 1 X 19 CrMo NbVN 11 1
QT
0.18–0.24
11–12.5
0.8–1.2
0.3–0.8
0.25–0.35
QT
0.17–0.23
10–11.5
0.5–0.8
0.2–0.6
0.1–0.3
0.07–0.13
14.0–16.0 0.8–1.2
9.0–11.0
–
0.04–0.1
18–21
<1.5
–
X 10 CrNiMoMn NbVB 15 10 1
NiCr20TiAl
AT P 2 stage
<1
Other
Nb 0.25– 0.55; N 0.05–0.1 Mn 5.5–7; B 0.003– 0.009; V 0.15–0.4; Nb 0.75–1.25; N<0.11 Al 1–1.8; B <0.008. Ti 1.8–2.7
T range (°C)
Strain range (%)
Duration (h)
T range (°C)
350–500 400–500 300–540 B 0.001– 0.010; Ti 0.007– 0.15 400–580
0.15 0.15 0.2 400–600
1k, 10k, 30k 1k, 10k, 30k 1k, 10k, 30k 0.15
1k, 10k, 30k
0.2
1k, 10k
400–600
0.2
1k, 10k, 30k
550–700
0.15
10k
450–640 650–750
0.15 0.15– 0.2
1k, 10k, 30k 30k 1k, 10k
30k, 430–590
Strain Duration range (%) (h)
0.15
10k, 30k., 100k, 200k
Specifications for creep-resistant steels: Europe
Material grade
Gross chemical composition Concentration in mass% of the following elements
30k
143
WPNL2204
144
Creep-resistant steels
four tests has to be produced with the longest test not shorter than 10 000 h, one test above 1000 h and one between 50 and 250 h; tests shorter than 50 h are not considered. The regression line of this test series is then compared with the ±20% scatter band around the mean rupture strength lines of the reference parent material, for which EN 14532-2 Appendix E gives the equations based on Manson–Brown-parameter approaches following the discontinued ISO 6303 Appendix. The source of the coefficients to the equations is the British Guideline PD 6525. All weld material specimens – according to EN 14532-2 – are expected to fall within the scatter band of the matching parent material. If this is not the case, as is likely for longer testing times, detailed warnings about the longer or shorter duration have to be stated on the consumables qualification certificate. To date, EN 14532-2 is the only standard that includes creep strength equations. Some materials are marked with an ‘x’ in Tables 3.5 to 3.15 because the equation applicability is not just related to the temperature–duration plane.
3.5.4
Testing and testing standards
A side product of the ECCC WG1 Creep Data Generation and Assessment Procedures and of the EC-funded Standard Measurement and Testing (SMT) Projects conducted in the 1990s16,17 produced a group of testing procedures, which: • • •
harmonised the rules of the main National Standards (BS 3500, DIN 50118, NF A03-355, ASTM E139/E292, UNI 5111, etc); collated the results from the relevant SMT projects; and critically overviewed the creep testing procedures of 15 leading European labs, as laid down in Reference 19 and Reference 20: Vol 3 part I.
These recommendations were taken up by CEN/ECISS TC1 and became, with slight modifications, the standards: • • •
EN 10291:2000 (Metallic materials – Uniaxial creep testing in tension), which included stress-rupture, creep-rupture and creep testing with both interrupted and non-interrupted strain measurement; EN 10319:2003 (part 1) – 2006 (part 2) (Metallic materials – Tensile stress relaxation testing – Part 1: Procedure for testing machines and Part 2: Procedure for bolted joint models); CWA 15261-3:2005 (Measurement uncertainties in mechanical tests on metallic materials – Part 3: The evaluation of uncertainties in creep testing). This standard has the support and is based on the ground rules as established by the EC funded projects UNCERT and UNCERT-AM.26
EN 10291 is currently part of the discussions regarding revision of international standard ISO 204 (Metallic materials – Uninterrupted uniaxial
WPNL2204
Specifications for creep-resistant steels: Europe
145
creep testing in tension – method of test). Creep testing on notched specimens has been proposed for inclusion according to the new ECCC-recommended notch geometry (Reference 19 and Reference 20: Vol. 3 part V; References 27, 28) and methods following the High Temperature Mechanical Testing Committee HTMTC Working Group on ‘Notch Creep Behaviour’.29,30 A particular testing technique (miniature specimen testing with the small punch method) is also under discussion in CEN WorkShop Agreement 21 (Small Punch Test Method for Metallic Materials) with the support of EPERC; the issuance of CWA 15627 concluded WS21’s work). Testing technique Compared to standards prepared elsewhere, the new European creep testing norms are based on the ECCC experience in refining testing requirements and some additional testing options. EN 10291 and EN 10319 were designed to best suit long duration tests, while realising that although most relevant to design and dimensioning, quality control laboratories would also use them for the more common short-duration stress rupture testing. •
Test methods and strain measurement: in Europe a large amount of the available long-term creep data was obtained in stress rupture tests, i.e. tests in which the specimen is held at constant temperature and load until fracture occurs. Such tests are often performed in multi-specimen testing machines, which may include up to ten strings each including ten or more specimens. Creep rupture testing, where creep strain is continuously measured during the test, is considered more expensive owing to the need to maintain thermal and electrical stability of the extensometer, or because ambient temperature must be controlled to reduce thermal drift in the transducers. It must be taken into account that during the initial stages of long- and medium-term creep rupture tests, elongations of a few micrometres must be measured, to keep uncertainty to a reasonable level. As a less expensive alternative,16 interrupted creep rupture tests are performed using techniques originally developed in the 1940s31 and improved in Germany in the 1980s.32 During these tests, the specimen is cooled, unloaded, its elongation measured, re-assembled, reloaded and re-heated at regular, generally logarithmically equi-spaced, time intervals. Although the ‘off line’ strain measurement, requires particular care and should be performed in a suitable metrological laboratory, the practical comparison of results achieved in a round robin test on BCR reference materials ‘Alloy 75’ (CRM 425, NiCr20Ti EN 10095) and ‘Durhete 1055’ (20 CrMoVTiB 4 10 EN 10269, 1%CrMoVTiB) showed,16 that interrupted and non-interrupted tests provide creep curves contained in the same scatter-band of approximately ±25% in strain
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•
•
Creep-resistant steels
around the mean line. Owing to these premises, EN 10291 and probably the future ISO 204 will regulate this alternative technique. Temperature: temperature stability is an essential parameter during longterm tests, and the main consideration here is distinction and consideration of the various factors contributing to thermal stability: EN 10291 defines temperature stability as the minimisation of the difference between true test temperature and specified test temperature, where these definitions account for all errors contained in the temperature measurement and regulation chain, i.e. regulation deviation plus random errors. Because of this, the allowable temperature differences are somewhat higher than, for instance, those in ASTM E139, where thermal stability has been considered equivalent to indicated temperature minus specified temperature, that is, to the regulation-induced deviations only. Another topic concerning thermal stability is the calibration of thermocouples, generally expensive PtRh–Pt thermocouples of type R, S or B, which unavoidably deteriorate during high temperature exposure owing to metallurgical instability.33 For long-duration tests, several calibration methods were compared and discussed34 and recommendations for both off-line and in situ calibration were stated (Reference 20: Vol 3 part I). Both of these consider the temperature range and gradient to which the thermocouple is exposed during its use, i.e. the depth to which the thermocouple is inserted into the testing furnace, which could be very high in multi-specimen machines where discarding the exposed thermocouple part may be very expensive. Data assessment is subdivided into two distinct steps: (i) Single test raw data assessment consists of transforming elongation time data into strain–time data by dividing elongation by the reference length, which, as a function of the test piece geometry, may be different from the cylindrical length because collars used for the extensometer fixation will contribute to the measured deformation. The reference length is therefore computed according to expected creep behaviour and the Norton stress exponent and may have to be re-assessed when the true results are obtained (Reference 20: Vol 3 part I Appendix 1). From the creep curve (strain versus time), initial plastic strain (if present) and times to specified plastic strains are collated, for which linear interpolation in a bi-logarithmic strain-time diagram appears to be best [20 Vol 3 Part I Appendix 1]. Small differences between interrupted and non-interrupted testing due to inelastic recovery may be negligible in most technical applications. (ii) Data assessment of an entire test programme is described in Chapter 6 by Holdsworth.
WPNL2204
Specifications for creep-resistant steels: Europe
147
Testing programmes Setting up a testing programme for parent materials depends on what the data are needed for. In material quality assurance, single specimen tests with smooth or combined smooth-notched specimens are required. Production qualification tests generally require an isothermal approach including four to six specimens per temperature and a duration not shorter than 10 000 h, sometimes 30 000 h to meet the requirement to have an extrapolation factor below 3 when computing 100 000 h strengths. The ECCC approach to common testing programmes, intended to fill gaps identified in the currently available creep data population of a given material, also follows isothermal programme set-ups, for which durations of 1000, 3000, 10 000, 30 000 and 70 000 h are requested. To fully meet significance criteria suitable for standardisation, the data additionally have to meet the requirements listed in Table 3.3. In situations where creep properties must be verified in a relatively short period of time, such as during residual life assessment or, in some cases, particular material appraisal according to PED for non-harmonised materials, alternative methods are adopted, although these are generally not standardised and strongly based on the individual assessor’s experience. In such cases ‘iso-stress’, ‘MPC-Ω’ approaches35,36 or reduced test programmes which produce an estimate of a Larson–Miller master curve are used. The effectiveness and credibility of these tests, although in some cases clearly demonstrated, has proven to be strongly dependent on the assessor’s skill and experience with the testing procedure and the material tested (Reference 20: Vol 5 part III). Component creep tests In the past, a considerable number of near- and full-scale tests for components have been developed in Europe. As a consequence, procedures to test closeto-component specimens, generally tubes and pipes, have also been developed. Owing to the complexity of the testing devices and procedures, few testing facilities are available, although an impressive number of tests has been collated (Reference 20: Vol. 9 part III). In addition, codes of practice for tubular specimens under internal pressure and combined internal pressure and tension or torsion have been published.37,38 Relaxation testing The use of bolting at high temperature has to take into account the loss in pre-loading with time (relaxation) owing to creep effects under constant total strain. Various testing techniques have been proposed:
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•
Creep-resistant steels
Uniaxial relaxation testing, keeping a standard specimen under total strain control, i.e. keeping its elasto–plastic strain constant over time (generally to a value comparable to the pre-load of a bolt, ca. 0.1–0.5% strain). Two testing approaches are used: –
–
•
Isothermal relaxation: the specimen is heated to test temperature, and loaded to the foreseen strain, which is then kept constant, while the loss in stress over time is recorded. Severe technical testing issues need to be addressed during such tests,17 for instance closedloop testing systems proved to be very difficult to keep stable over long periods of time, when the stress decrease becomes small. Non-isothermal relaxation: in this case the specimen is loaded at room temperature with the required total strain and is then heated to the testing temperature.
Model bolt relaxation testing, developed and already standardised in Germany,39 consists of preloading a series of bolt-flange assemblies at room temperature and then tightening the screws to a known and controlled pre-load (generally measuring the bolt shaft strain using strain gauges). The assemblies are then put into a furnace and heated. One by one the bolt-flange assemblies are extracted at specified time intervals, cooled and loosened, using strain gauges to measure the remnant bolt pre-load.
A comparison of the two methods was conducted on a typical bolting steel (“X19” = X19CrMoNbVN 11 1 EN 10269, 11%CrMoNbNV) under the SMT project.17 Results showed that for long periods of time (ca. 10 000 h), the remnant stress measured using the two different methods are comparable. For both methods, detailed guidelines are found in ECCC Recommendations 2005, Vol 320 and in EN 10319, parts 1 and 2. Design standards At present there are no generally available European standards for design that directly address components which operate under European creep regimes. Nevertheless, considerable work is underway for the new standards EN 13445 part 3 (Unfired pressure vessels – Part 3: Design) and EN 13480 part 3 (Metallic industrial piping – Part 3 Design) and their various appendices. To be in line with PED, the new standards will, besides traditional creep design methods, also have to cover the dimensioning of welds, the use of finite element simulation techniques, the applicability of fracture mechanics and the consideration of statistics and/or risk-based methods (CEN Workshop 22 on FITNET and CEN Workshop 24 on Risk Based Inspection and Maintenance). The related CEN-TCs, EPERC-TP and ECCC are working towards compiling collations of reliable creep strength equations, weld creep strength reduction factors and other material related design features.
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3.5.5
149
Specifications of creep-resistant steel users
The market for plant building in refineries, chemical plants and power units has changed since the publication of the PED with the breakup and privatisation of monolithic electricity boards and mergers/closures of big boilermakers. Plant building now tends to be in the hands of general engineering companies, which have significantly less skill and experience in handling creep problems and the materials suitable for such operating conditions. In addition, plant owners are asking for longer and longer service periods, so that creep strength values of 200 000 h or 250 000 h are no longer the exception. As a consequence, the plant builders’ specifications for creep-resistant steels tend to refer as far as possible to any available standards and they ask material, component fitting or spool suppliers to guarantee the values in accordance with PED and under all possible configurations during assembly and operation. Since long-term creep test results have up to now been collected and stored by specialist companies (boiler makers, turbine manufacturers, large furnace and reactor builders) rather than the material producers, negotiations to find a PED-conforming compromise are often very difficult. Even the specialist companies are often unable to provide experimental results for materials after complex forming and heat-treating operations (shapes, fittings, repaired castings, etc) or on welds after unusual heat treatments (for instance welds subjected to more than three post-weld heat treatments owing to repair or close assembly to manufacturing welds). As a consequence, the creep strength values of EN standards tend to become mandatory values in engineering companies’ specifications, sometimes lowered by a percentage apparently accounting for the usual ±20% scatter band in stress on rupture strength values. In most typical situations, where the material, component, fitting, spool or vessel manufacturer cannot comply with pre-existing data, ‘qualification’ creep test rows are requested, which, owing to the short time frame and the belief that ‘time–temperature parameter based’ extrapolations are safe, are generally reduced to a few tests below 1000 h, then grandly extrapolated to the required 200 000 h strengths. Validation procedures, like ECCC’s PATs (Reference 20: Vol 5), are generally unknown and not applied.
3.5.6
Residual life assessment
Residual life assessment, that is, activities to decide whether a given plant can be used beyond its original design life, is generally led by in-service inspection procedures, dictated by national rules and/or by more advanced systems like RBI (risk based inspection). Nevertheless, some national regulations require that design reviews are also conducted, during which past as well as future service is checked against the most credible available
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creep strength values. This requires on the one side sound assessment procedures, because residual life assessment tends to ask for very long durations, longer than available and possible tests, and on the other side the refined strength values published in the various documentation for the standards. Some guidelines on how to handle these particular needs were laid down in Reference 20, Vol 5 and are assessed in the CEN WorkShop Agreements 22 (FITNET – European Procedure for Fitness for Service) and 24 (Risk Based Inspection and Maintenance Procedures – RIMAP) and in the new Italian standard for residual life assessment, UNI 11096: 2004.
3.6
Future trends
The new concept of CEN standards, including creep strength values, and the new ECCC strength assessment procedure were a first step towards improving the reliability of creep strength and design values. Nevertheless, a huge amount of work is still to be done: •
• •
• •
•
New materials are continuously offered by leading companies with metallurgical skills, for instance the new 9–12%Cr steels and the new low-alloyed steels of 2.25%Cr type, as well as many new austenitic and nickel base grades. These materials need big duration tests and thorough assessment, but, since industrial users never have the time to wait for these results to come through, credible creep strength reference values at least are needed as soon as possible. The prescriptions of EN 12952-2 and ECCC (Reference 20: Vol 5 part I), along with the data requirements according to Table 3.3, make it clear that testing has to be taken further. Reliable assessments are only possible with very large, or at least strategically planned, Europe-wide coordinated test campaigns. ECCC and other bodies have also identified large gaps in the understanding of creep behaviour and data availability in consolidated materials, which urgently need to be filled. Some materials, obsolete or no longer used for their original applications, may suddenly find new users or applications (for instance alloy 617 for piping and forgings), so that data previously disregarded become relevant again and may need to be up-dated. Cast steel grades in general, with the exception of the 9–12%Cr steels, have not been a priority for assessment. A very particular area, mainly the concern of refineries and petrochemical plants, are the cast austenitic high-alloyed steels (i.e. G-X40CrNi25 20 or G-X40NiCr35 25, partially included in EN 10295) and their modifications based on mixtures of Co, W and Nb. These have not yet been considered in detail. CEN will review all standards on a five-to-ten year basis, so that creep strength values will be updated and upgraded in the light of on-going
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tests and new high-priority materials (for instance grades 23 (10 CrWMoNbB 9 6), 24 (10 CrMoVTiB 10 10), 92 (X 10 CrWMoVNb 9 2), 911 (X 19 CrWMoVNb9 1 1), 709 (X 6 NiCrMoTiBN 25 20), etc). Some changes and refinements, particularly for those grades assessed without fully representative data sets or under reduced testing times (for instance grade 91, X 9 CrMoVNNb 9 1) will need to be upgraded in the published strength values. Besides the specifically material-related issues, there are other topics that will require strong collaboration between material specialists and designers, such as how to handle strength values of extended extrapolation, how to include time-dependent weld creep strength reduction factors, and so on. There is also an urgent need to harmonise the PED interpretation of creep handling in pressure vessel design, which, owing partly to a somewhat vague formulation, has led to radically different interpretations among the notified bodies. Agreement and reciprocal understanding is also desirable between the mandatory European PED and the forthcoming very advanced voluntary design codes (EN 13445, etc) and the well-established ASME, API and Japanese codes. Finally a few basic issues should also be addressed, including how to extrapolate from data sets of limited size or duration in the most reliable way and how to consider scatter or probabilistic effects for creep strength data assessments. These should be addressed by more basic research projects. In the past, Europe has shown that, despite the inherent difficulties, collaboration between different technical schools of thinking is possible and can lead to outstanding results. The next step is to include non-Europeans and to enlarge the multi-disciplinarity of the interacting specialists.
3.7
References
1 French, Tucker ‘Flow in a low-carbon steel at various temperatures’ NIST report T296, 1925; (re-published by Geil and Carwile, National Institute of Standards and Technology (NIST) RP2329, 1952). 2 Pomp A. and Dahmen A., ‘Entwicklung eines abgekürzten Prüfverfahrens zur Ermittlung der Dauerstandfestigkeit von Stahl bei erhöhten Temperaturen’, Mitteilung a. d. Kaiser-Wilhelm-Institut für Eisenforschung zu Düsseldorf, Abhandlung 72–95 IX Band (1927), 33–52. 3 Parravano N. and Guzzoni G., ‘Prove statiche delle leghe ultraleggere’, La Metallurgia Italiana, 1930, XXII, 367 ff. 4 DIN 50 112 (DVM-Creep Test), ca. 1930. 5 German Iron and Steel Society (VDEh), Ergebnisse deutscher Standzeitversuche langer Dauer, in collaboration with the German Creep Resistant Steel Committee (AGW) and the Committee for High Temperature Engineering Materials, Verlag Eisen und Stahl, Düsseldorf, 1969.
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6 Former CECA/ECSC (European Carbon and Steel Collaboration): http:// cordis.europa.eu/ecsc-coal/home.html, and actual http://cordis.europa.eu/coal-steelrtd 7 COST website: www.cost.esf.org 8 Information on EC programme Joule–Thermie: http://erg.ucd.ie/ttp.html 9 Official PED-site http://ec.europa.eu/enterprise/pressure_equipment/ped/index_en.html 10 European Standardisation Body: www.cenorm.be 11 European Committee for Iron and Steel Standardization ECISS: http://www.cenorm.be/ cenorm/aboutus/structure+/relations/workprogrammeeciss.asp 12 European Collaborative Creep Committee website: www.etd1.co.uk/eccc/ advancedcreep 13 EC funded Brite Euram Concerted Action, Creep –BE-5524, 1992–1996. 14 EC funded Thematic Network, BET2-0509 - Weldcreep, 1997–2001. 15 EC funded Framework V Thematic Network GTC2-2000-33051 Advanced Creep, 2001–2005. 16 EC funded Standard–Measurement and Testing Project MAT1-CT-940065, Development and Validation of a Code of Practice for a Low Cost Method of Strain Measurement from Interrupted Creep Testing, 1994–1998. 17 EC funded Standard–Measurement and Testing Project MAT1-CT-940078, Development of Standard European Methodology for Stress Relaxation Testing of Metals, 1994–1998. 18 International Conference on Creep & Fracture in High Temperature Components – Design and Life Assessment Issues, organized by ECCC, September 12–14, London, 2005. 19 ECCC Recommendations 1994, ERA Technology Ltd, Leatherhead. Volume 1 (issue 1) 1994: Creep Data Validation and Assessment Procedures – Overview, Holdsworth S.R. (ed.). Volume 2 (issue 2) 1994: Terms and Terminology for use with Stress Rupture, Creep and Stress Relaxation: Testing, Data Collation and Assessment, Orr J. (ed.). Volume 3 (issue 2) 1994: Recommendations for Data Acceptability Criteria for Creep, Creep Rupture, Stress Rupture and Stress Relaxation Data, Holdsworth S.R. and Granacher J. (eds). Volume 4 (issue 1) 1994: Guidance for the Exchange and Collation of Creep Rupture, Creep Strain–Time and Stress Relaxation Data for Assessment Purposes, Merckling G. & Bullough C.K. (eds). Volume 5 (issue 1) 1995: Guidance for the Assessment of Creep Rupture, Creep Strain snd Relaxation Data, Holdsworth S.R. (ed.). 20 ECCC Recommendations 2005, ETD Ltd, Leatherhead. Volume 1 (issue 6) 2005: Creep Data Validation and Assessment Procedures – Overview, Holdsworth S.R. (ed.). Volume 2, 2005: Terms and Terminology for use with Stress Rupture, Creep, Stress Relaxation, Creep Crack Initiation and Multi-Axial Creep: Testing, Data Collation and Assessment, Morris P., Orr J., Servetto C., Seliger P., Holdsworth S.R. and Brown T.B. (eds). (a) part I (issue 8): Parent Material, (b) part Iia (issue 2): Welding Processes and Weld Configurations, (c) part Iib (issue 2): Weld Creep Testing, (d) part III (issue 4): Post Exposure Material, (e) part IV (issue 2): Generation and assessment of creep crack initiation data, (f) part V (issue 1): Generation and assessment of multi-axial feature specimen and component data. Volume 3, 2005: Recommendations for Data Acceptability Criteria and Generation for Creep, Creep Rupture, Stress Rupture, Stress Relaxation, Creep Crack Initiation and Multi-Axial Creep Data, Holdsworth S.R., Granacher J., Klenk A., Buchmayr B., Gariboldi E., Brett S., Merckling G., Müller F., Gengenbach T., Dean D. and Brown T.B. (eds). (a)
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22
23 24 25 26 27 28
29
30
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part I (Issue 5): Generic recommendations , (b) part II (issue 3): Creep data for welds, (c) part III (issue 4): Creep testing of post exposure (ex-service) material, (d) part IV (issue 2): Creep crack initiation, (e) part V (issue 1): Multi-axial feature specimen and component test data. Volume 4, 2005: Guidance for the Exchange and Collation of Creep Rupture, Creep Strain–Time, Stress Relaxation, Creep Crack Initiation and Multi-Axial Creep Test Data for Assessment Purposes, Merckling G., Calvano F., Bullough C.K., Holdsworth S.R. and Tonti A. (eds). (a) part I (issue 6): Creep rupture, strain–time and relaxation data; (b) part II (issue 1) Creep crack initiation, (c) part III (tba); (d) part IV (issue 1): Creesty user manual. Volume 5, 2005, Guidance for the Assessment of Creep Rupture, Creep Strain and Stress Relaxation Data, Holdsworth S.R. and Merckling G., (eds). (a) Part Ia (issue 5): Full-size creep rupture data sets datasets, (b) Part Ib (issue 2): Creep strain and creep strength data, (c) Part Ic (issue 2): Full-size stress relaxation datasets, (d) Part Iia (issue 1): Sub-size datasets, (e) Part Iib (issue 1): Weld creep rupture datasets, (f) Part III (issue 2): Post exposure (ex-service) creep data. Volume 6, 2005 (issue 1): Residual Life Assessment and Microstructure, Concari S. (ed.). Volume 9, 2005: High Temperature Component Analysis, Patel R. (ed.). (a) part II (issue 1) Overview of assessment & design procedures, (b) part III (issue 1): Database of component tests and assessments. Holdsworth S.R., Orr J., Granacher J., Merckling G. and Bullogh C.K. on behalf of the ECCC-WG1, ‘European Creep Collaborative Committee activities on creep data generation and assessment methodologies’ in Materials for Advanced Power Engineering, D. Coutsouradis, F. Schubert, D.V. Thornton and J.H. Davidson (eds), Liege, 3–6 October 1994, Kluwer Academic Publishers, 1994, 591–600. Merckling G. and Holdsworth S.R., ‘Long term creep rupture strength assessment: the development of the European Collaborative Creep Committee Post assessment tests’ , International Conference on Creep & Fracture in High Temperature Components – Design and Life Assessment Issues, September 12–14, London, 2005. ECCC Data Sheets, issue 1996. ERA Technology Publishers, 1996. ECCC Data Sheets, issue 2005. European Technology Development Publishers, 2005 European Pressure Equipment Research Council Web Site: www.mpa-lifetech.de/ eperctp/ EC-funded Thematic Network and CEN Workshop 11 “UNCERT-AM”: http:// www.mpa-lifetech.de/UNCERT-AM/HTML_Files/Main/UNCERTDefault.htm Morris P. and Granacher J., ECCC-Information Day, European Technology Development Ltd, Prague, 2001. Scholz A., Schwienheer M. and Morris P.F., ‘European notched testpiece for creep rupture testing’, in: Proceedings of the 21st Symposium of the German Iron & Steel Institute (VDEh) & German Society for Material Testing (DVM), Herausforderung durch den Industriellen Fortschritt, 4–5 December 2003, Buchholz W.O. and Geisler S. (eds), Bad Neuenahr, Verlag Stahleisen, Düsseldorf, 2003, 308–314. Webster G.A., Holdsworth S.R., Loveday M.S., Perrin I.J. and Purper H., A Code of Practice for Conducting Notched Bar Creep Rupture Test and the Interpretation of the Data, ESIS TC11, 1999. Webster G.A., Holdsworth S.R., Loveday M.S., Nikbin K., Perrin I.J., Purper H., Skelton R.P. and Spindler M.W., ‘A code of practice for conducting notched bar creep tests and for interpreting the data, issue 3’, Fatigue and Fracture of Engineering Materials & Structures, 2004, 27 (4), 319–342.
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31 Matteoli L. and Andreini, B., ‘Le prove di scorimento interrotte. Influenza delle interruzioni di sollecitazione e di riscaldamento sulle proprietà di scorrimento’, La Metallurgia Italiana, 1947, XXXIX, 41. 32 Granacher J., Oehl M. and Preußler T., ‘Comparison of interrupted and uninterupted creep rupture test’, Steel Research, 1992, 63, 39–45. 33 Granacher J. and Scholz J., ‘Über die langzeitige Temperaturgenauigkeit von Zeitstandsprüfanlagen’, Materialprüf., 1973, 15, 116–123. 34 McCarthy P. and Loveday M.S. (eds), Proceeding of the Seminar on the Practicalities of Thermocouple Calibration and the Usage for Materials Testing, High Temperature Mechanical Testing Committee HTMTC, 16th December 1996. 35 Regis V., Livraghi M. and Di Pasquantonio F. Invecchiamento dei materiali per scorrimento viscoso, ENEL-CRTN, 1972. 36 Prager M., ‘Development of the MPC omega method for life assessment in the creep range’, J. Pressure Vessel. Technology, 1995, 117, 95. 37 How I.M., Browne R.J., Coleman M.C., Craig I.H., Ham M.W., Hurst R.C. and Meecham P.C., ‘A code of practice for internal pressure testing of tubular components at elevated temperatures’, in Proceedings of the HTMTC Symposium on Harmonised Testing Practice for High Temperature Materials, Loveday M.S. and Gibbons T.B. (eds), Ispra, Italy 18–19 October 1990, 363–400. 38 Rees D.W.A., Brown M.W., Hyde T., Lohn R.D., Morrison C.J. and Shammas M., ‘A code of practice for torsional creep testing of tubular testpieces at elevated temperatures’, in Proceedings of the HTMTC Symposium on Harmonised Testing Practice for High Temperature Materials, Loveday M.S. and Gibbons T.B. (eds), Ispra, Italy, 18–19 Octobers 1990, 331–361. 39 SEP 1260, Relaxationsversuch bei erhöhter Temperatur mit Schraubenverbindungsmodellen, VDEh, 1996.
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4 Specifications for creep-resistant steels: Japan F . M A S U YA M A, Kyushu Institute of Technology, Japan
4.1
Introduction
Creep-resistant steels have been widely used in structural components and machinery operated under high temperature and high pressure. There are many different types of heat-resistant steels, as the properties required vary according to application and these properties can generally be categorized into physical, chemical, mechanical, manufacturability and economical aspects. Chemical composition influences physical and chemical properties such as thermal conductivity, thermal expansion and high temperature corrosion/ oxidation resistance. Mechanical properties, including hardness, tensile, impact, creep and fatigue properties, depend on microstructures formed by heattreatment and chemical composition. The general specifications for heatresistant steels in Japan established JIS (Japanese Industrial Standards)1 and in METI (Ministry of Economy, Trade and Industry)2 Codes and JSME (Japan Society of Mechanical Engineers)3 Codes for power applications. This section reviews the status of specifications for heat-resistant steels in Japanese codes and standards.
4.2
Types of heat-resistant steels in Japan
Table 4.1 lists heat-resistant steels (SUH) and stainless steels (SUS) in JIS G4311 and G4312 for use under high temperature conditions, indicating their applications as well. In the case of ferritic steels, cooling in air from the annealing temperature of 780–880°C is necessary, since slow cooling at around 600°C causes material embrittlement owing to sigma phase precipitation. These steels can be used for heater boxes, burners and heater equipment operated up to 1100°C in air and gas containing sulphur. Martensitic steels are annealed in the temperature range of about 800–900°C, slow cooled at about 1000–1100°C, oil quenched and tempered at about 650–800°C, and then subjected to rapid cooling or air cooling. These steels are used for high temperature intake valves or turbine blades, applications that utilize their 155 WPNL2204
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Table 4.1 Heat-resistant steels in JIS
Martensitic
Precipitation hardening Austenitic
Nominal composition (equivalent)
Applications
SUH 21 SUH 409 SUH 446 SUS 405 SUS 430 SUH 1 SUH 3 SUH 4 SUH 11 SUH 600 SUH 616 SUS 403 SUS 410 SUS 410J1 SUS 420J1 SUS 420J2 SUS 431 SUS 440A SUS 440B SUS 440C SUS 630 SUS 631 SUH 31 SUH 35 SUH 36 SUH 37
19Cr–3Al (18SR, JIS FCH2) 11Cr–Ti 25Cr–N–0.4C 13Cr–0.2Al–0.06C 18Cr-0.1C 9Cr–3Si–0.4C 11Cr–2Si–1Mo–0.4C 20Cr–1.5Ni–2Si–0.8C 9Cr–1.5Si–0.5 12Cr–0.6Mo–0.3V–0.5Nb–N–0.15C (H46) 12Cr–1Mo–1W–0.25C AlSl 422) 13Cr–low Si–0.1C 13Cr–0.5Mo–0.1C 13Cr–Mo–0.15C 13Cr–0.2C 13Cr–0.4C 16Cr–2Ni–0.15C 18Cr–0.5Mo–0.7C 18Cr–0.5Mo–0.85C 18Cr–0.5Mo–1C 17Cr–4Ni–4Cu–Nb–0.05 (17–4PH) 17Cr–7Ni–1Al–0.07C (17–7PH) 15Cr–14Ni–2Si–2.5W–0.4C 21Cr–4Ni–9Mn–0.45N–0.5C (21–4N) 21Cr–4Ni–9Mn–0.45N–0.5C–high S (21–4N) 21Cr–11Ni–0.25N–0.2N (21–11N)
SUH 38
20Cr–11Ni–2.3Mo–0.2P–0.3C–B (20–11P)
Heater, automobile exhaust gas cleaner Automobile exhaust gas cleaner, muffler Combustion chamber Combustion turbine compressor blade Furnace, bumer parts up to 900°C Diesel engine intake valve Valve, pre-combustion charier Inhale/exhaust valve, abrasion resistant parts Gas/diesel engine intake valve up to 750°C Steam turbine blade, disk, rotor, bolt Steam turbine blade, disk, rotor, bolt Steam turbine blade, nozzle High temperature oxidation resistant parts up to 800°C Steam turbine blade, high temp./press. components Steam turbine blade, valve stem, nozzle, pump shaft Piston ring, fuel injection nozzle, seat ring Shaft, bolt, spring Valve parts, ball valve Bearing, roller, valve parts Bearing, abrasion resistant parts Steam/combustion turbine blade High temp. spring, bellows Gas/diesel engine exhaust valve up to 1150°C Gas/diesel engine exhaust valve with high strength Gas/diesel engine exhaust valve with high strength Gas/diesel engine exhaust valve with oxidation resistance Gas/diesel engine exhaust valve, bolt
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Ferritic
Designation
SUH SUH SUH SUH SUH
309 310 330 660 661
SUS 304
22Cr–12Ni–0.2C 25Cr–20Ni–0.2C 15Cr–35Ni–0.1C 15Cr–25Ni–1.5Mo–0.5V–2Ti–Al–B (A286) 22Cr–20Ni–20Co–3Mo–2.5W–Nb–0.15N (LCN155) 18Cr–8Ni
SUS 309S SUS 310S SUS 316 SUS 317 SUS 321 SUS 347 SUSXM15J1
22Cr–12Ni–0.06C 25Cr–20Ni–0.06C 18Cr–12Ni–2.5Mo–0.06C 18Cr–12Ni–3.5Mo–0.06C 18Cr–8Ni–Ti–0.06C 18Cr–8Ni–Nb–0.06C 18Cr–13Ni–4Si (ASTM XM15)
Furnace, burner cyclically operated up to 980°C Furnace, burner cyclically operated up to 1035°C Furnace, oil cracking equipment Turbine rotor, bolt, shaft up to 700°C Turbine rotor, bolt, blade, shaft up to 750°C Cyclic heating equipment with oxidation resistance up to 870°C Furnace parts cyclically operated up to 980°C Furnace, automobile exhaust gas cleaner up to 1035°C Heat exchanger with high creep strength heat exchanger with high creep strength Welded parts for corrosion-resistant use at 400–900°C Parts for corrosion resistant use at 400–900°C Automobile exhaust gas cleaner
Specifications for creep-resistant steels: Japan
Austenitic
157
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superior oxidation-resistant and/or creep-resistant properties. Austenitic heatresistant steels listed in JIS G4311 and G4312 are for furnace equipment, heat exchangers and chemical plants taking advantage of their superior high temperature corrosion/oxidation resistance and creep resistance. Except for SUH35 to 38 and SUH660/661, which need to be age-hardened, austenitic steels are used as-solution annealed at about 1030–1200°C, followed by rapid cooling. For the aforementioned exceptional steels, aging treatments are applied at about 700–800°C after solution annealing.
4.3
Specifications for high temperature tubing and piping steels
Table 4.2 shows JIS and METI Code Specifications for high temperature tubing and piping steels. These tables include steel designation in the JIS and METI Code, minimum tensile and yield strengths, heat treatment conditions and application components in power boilers. However, several of these steels are not actually used in power boilers despite being listed in the codes and standards. The maximum usage temperature is determined by the designer, but the allowable stresses are listed in the METI Code up to the temperature indicated in the code for individual steels. The criteria for allowable stresses in the creep temperature range in the METI Code are stated as follows. The allowable tensile stress in the creep temperature range shall not exceed the minimum values given below: 1 2 3
average value of the stress by which creep rate of 0.01%/1000 h occurs at the relevant temperature; 67% of the average value of the stress by which rupture takes place in 100 000 h at the relevant temperature; 80% of the minimum value of the stress by which rupture takes place in 100 000 h at the relevant temperature.
On the other hand, the JSME Code recently changed the above criteria (2), substituting 67% with 100Favg% in accordance with the ASME criteria, where Favg is a multiplier applied to average stress for rupture in 100 000 h. At 815°C and below, Favg is 0.67. Above 815°C, it is determined from the slope of the log time-to-rupture versus log stress plot at 100 000 h, such that log Favg = 1/n, but it may not exceed 0.67. n is a negative number equal to δ log time-to-rupture divided by δ log stress at 100 000 h. METI Code materials are defined as KA-XXXX, with XXXX, taking on a JIS-like designation that has not yet been integrated into JIS and indicating power applications only. However, new steels recently developed in Japan for power boiler applications, such as high strength ferritic steels and austenitic steels, are included in the METI Code with the symbol ‘J’ followed by the development number in the designation. Tables 4.3 and 4.4 show the chemical compositions
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Table 4.2 JIS and METI Code specification for high temperature tubing and piping materials
KA-STB480
C-Steel
480
275
Ann. or nor.
Alloy steel tube for boiler & heat exchanger
STBA12
0.5Mo
380
205
STBA13 STBA20
0.5Mo 0.5Cr, 0.5Mo
410 410
205 205
STBA22
1Cr, 0.5Mo
410
205
STBA23
1.25Cr, 0.5Mo
410
205
Low temp. ann. isothermal ann., full ann. Nor. or nor. + temper Low Temp. ann. Iso-thermal ann. Full ann. or nor. + temper iso-thermal ann. full ann. or nor. + temper
STBA24 STBA25 STBA26
2.25Cr, 1Mo 5Cr, 0.5Mo 9Cr, 1Mo
410 410 410
205 205 205
JIS G 3462 (1988)
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✓
✓
✓
✓
✓
✓ ✓
✓ ✓
✓ ✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
Steam pipe
✓ ✓
159
C-steel tube for power METI Code
✓ ✓ ✓
SH/Rh header
No specification
Reheater
175 255 295
Superheater
340 410 510
Connect pipe
C-Steel C-Steel C-Steel
Applications
Header
Min. Heat treatment yield (MPa) Water wall
Min. tensile (MPa)
Specifications for creep-resistant steels: Japan
Part I C-steel tube for STB340 boiler & heat STB410 exchanger STB510 JIS G 3461 (1988)
Nominal composition (mass%)
Economizer
Designation
160
Table 4.2 Cont’d
2.25Cr, 1.6W 9Cr, 2Mo
510 510
400 295
KA-STBA28
9Cr, 1Mo, Nb, V
590
410
KA-STBA29
9Cr, 1.8W
620
440
SUS304TB SUS304HTB SUS304LTB SUS309TB SUS310TB SUS316TB SUS316HTB SUS316LTB SUS321TB
18Cr, 18Cr, 18Cr, 23Cr, 25Cr, 16Cr, 16Cr, 16Cr, 18Cr,
520 520 480 520 520 520 520 480 520
205 205 175 205 205 205 205 175 205
Part II Stainless steel tube for boiler & heat exchanger JIS G 3463 (1994)
8Ni 8Ni 8Ni, Low C 12Ni 20Ni 12Ni, 2Mo 12Ni, 2Mo 12Ni, 2Mo, Low C 10Ni, Ti
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1010°C 1040°C 1010°C 1030°C 1030°C 1010°C 1040°C 1010°C 920°C
Steam pipe
KA-STBA24J1 KA-STBA27
Nor. Ann. or nor. + temper Ann. or nor. + temper at min. 650°C Nor. + temper Nor. at min. 900°C+ temper at min. 700°C Nor. at min. 1040°C+ temper at min. 730°C Nor. at min. 1040°C+ temper et min. 730°C
SH/Rh header
255 205 205
Reheater
410 410 410
Superheater
1.25Cr, 0.3Cu 1Cr, 0.3Mo 2.25Cr, 1Mo
✓
✓ ✓
✓ ✓
✓
✓ ✓
✓ ✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
✓
Connect pipe
Alloy steel tube KA-STBA10 for power boiler KA-STBA21 METI Code KA-STBA24EG
Applications
Header
Min. Heat treatment yield (MPa) Economizer
Min. tensile (MPa)
Creep-resistant steels
Nominal composition (mass%)
Water wall
Designation
✓
✓
590 690 590 690
235 345 245 345
1040°C 1050°C 1050°C 1050°C
✓ ✓ ✓ ✓
✓ ✓ ✓ ✓
590 660 640 650
235 295 270 295
1120°C 1030°C 1100°C 1030°C
✓ ✓ ✓ ✓
✓ ✓ ✓ ✓
520
205
1100°C
✓
✓
520
205
SUS347TB SUS347HTB
18Cr, 10Ni, Nb 18Cr, 10Ni, Nb
520 520
205 205
SUS410TB SUS430TB
13Cr 16Cr
410 410
KA-SUS304J1HTB 18Cr, 9Ni, 3Cu, Nb, N KA-SUS309J1TB 24Cr, 15Ni, 1Mo, N KA-SUS309J2TB 22Cr, 14Ni, 1.5Mo, N KA-SUS309J3LTB 25Cr, 14Ni, 0.8Mo, N, 0.2Si KA-SUS309J4HTB 22Cr, 15Ni, Nb KA-SUS310J1TB 25Cr, 20Ni, Nb, V KA-SUS310J2TB 20Cr, 25Ni, 1.5Mo KA-SUS310J3TB 22.5Cr, 18.5Ni, 1.8W, 3Cu, 0.45Nb, 0.2N KA-SUS321J1HTB 18Cr, 10Ni, Ti, Nb
WPNL2204
161
✓
18Cr, 10Ni, Ti
Specifications for creep-resistant steels: Japan
Part III Stainless steel tube for power boiler METI Code
✓
205 245
1095°C for cold finished, 1050°C for hot finished 980°C 1095°C for cold finished, 1050°C for hot finished 700°C AC or FC 700°C AC or FC
SUS321HTB
162
Table 4.2 Continued Min. tensile (MPa)
Min. Heat treatment yield (MPa)
Superheater
Reheater
10Ni, 3Cu, Nb,
500
205
1160°C
✓
✓
10Ni, Nb 9Ni, V, W, Nb, N 1Mo, W, V, Nb 2W, 0.4Mo, V, Nb 2W, 0.4Mo, V, Nb
520 650 590 620 620
205 270 390 400 400
1150°C 1100°C Nor. + temper Nor. + temper Nor. + temper
✓ ✓ ✓ ✓ ✓
✓ ✓ ✓ ✓ ✓
C-steel C-steel C-steel
370 410 480
215 245 275
As rolled for hot finished low temp. ann. or nor. for cold finished
✓ ✓ ✓
✓ ✓ ✓
Part IV Alloy steel piping
STPA12
0.5Mo
380
205
✓
✓
STPA20
0.5Cr, 0.5Mo
410
205
✓
✓
STPA22 STPA23 STPA24
1Cr, 0.5Mo 1.25Cr, 0.5Mo 2.25Cr, 1Mo
410 410 410
205 205 205
Low temp. ann. isothermal ann. full ann. nor. or nor. + temper low temp. ann. isothermal ann. full ann. or nor. + temper Iso-thermal ann. full ann. or nor. + temper
✓ ✓
✓ ✓
JIS G 3458 (1988)
WPNL2204
Steam pipe
Connect pipe
Header
C-steel for high STPT370 temp. pipe JIS STPT410 G3456 (1988) STPT480
SH/Rh Header
KA-SUS321J2HTB 18Cr, Ti, B, N KA-SUSTP347HTB 18Cr, KA-SUS347J1TB 18Cr, KA-SUS410J2TB 12Cr, KA-SUS410J3TB 11Cr, KA-SUS410J3DTB 12Cr,
Water wall
Applications
✓ ✓
✓ ✓
Creep-resistant steels
Nominal composition (mass%)
Economizer
Designation
Part V Stainless steel piping JIS G 3459(1997)
5Cr, 0.5Mo 9Cr, 1Mo
410 410
205 205
KA-STPA21 KA-STPA24J1 KA-STPA27
1Cr, 0.3Mo 2.25Cr, 1.6W 9Cr, 2Mo
410 510 510
205 400 295
KA-STPA28
9Cr, 1Mo, Nb, V
590
410
KA-STPA29
9Cr, 1.8W
620
410
SUS304TP SUS304HTP SUS304LTP SUS309TP SUS310TP SUS316TP SUS316HTP SUS316LTP SUS321TP SUS321HTP
18Cr, 18Cr, 18Cr, 23Cr, 25Cr, 16Cr, 16Cr, 16Cr, 18Cr, 18Cr,
520 520 480 520 520 520 520 480 520 520
205 205 175 205 205 205 205 175 205 205
8Ni 8Ni 8Ni, Low C 12Ni 20Ni 12Ni, 2Mo 12Ni, 2Mo 12Ni, 2Mo, Low C 10Ni, Ti 10Ni, Nb
1010°C 1040°C 1010°C 1030°C 1030°C 1010°C 1040°C 1010°C 920°C 1095°C for cold finished, 1050°C for hot finished
✓
✓
✓ ✓
✓ ✓
✓
✓
✓
✓
163
WPNL2204
Ann. or Nor. + Temper Nor. + Temper Nor. at min. 900°C + temper at min. 700°C Nor. at min. 1040°C + temper at min. 730°C Nor. at min. 1040°C + temper at min. 730°C
Specifications for creep-resistant steels: Japan
Alloy steel power piping METI code
STPA25 STPA26
164
980°C 1095°C for cold finished, 1050°C for hot finished
Stainless steel power piping
KA-SUS410J3TP
11Cr, 2W, 0.4Mo, V, Nb
620
400
Nor. + temper
Ann = annealing, Nor = normalizing, Temper = tempering, AC = air cooling, FC = furnace cooling.
WPNL2204
Steam pipe
205 205
SH/Rh header
520 520
Reheater
18Cr, 10Ni, Nb 18Cr, 10Ni, Nb
Superheater
SUS347TP SUS347HTP
Applications Connect pipe
Min. Heat treatment yield (MPa) Header
Min. tensile (MPa)
Water wall
Nominal composition (mass%)
Economizer
Designation
✓
✓
Creep-resistant steels
Table 4.2 Continued
Table 4.3 Chemical compositions for ferritic boiler tubing steels Steels
Chemical composition (mass%) C
KA-STB510
≤ 0.25
≤ 0.35
KA-STB480
≤ 0.30
≤ 0.10
KA-STBA10
≤ 0.10
STBA12
STBA23
0.10 ~0.20 0.10 ~0.20 ≤ 0.15
0.20 ~0.80 0.10 ~0.50 ≤ 0.50
STBA24
≤ 0.15
0.50 ~1.00 ≤ 0.50
KA-STBA24E-G
≤ 0.15
≤ 0.50
0.04 ~0.10
≤ 0.50
STBA26
≤ 0.15
KA-STBA27
≤ 0.08
0.25 ~ 1.00 ≤ 0.50
KA-STBA28
0.08 ~0.12
KA-STBA21
KA-STBA24J1
9Cr
0.20~ 0.50
P
S
Ni
Cr
Mo
V
Nb
Al
N
W
B
0.30 ~0.80 1.00 ~1.50 0.29 ~1.06
≤ 0.035
≤ 0.035
–
–
–
–
–
–
–
–
–
≤ 0.035
≤ 0.035
–
–
–
–
–
–
–
–
–
≤ 0.048
≤ 0.058
–
–
–
–
–
–
–
–
–
≤ 0.80
≤ 0.025
–
–
–
–
–
–
–
–
0.30 ~0.80 0.30 ~0.60 0.30 ~0.60 0.30 ~0.60 0.30 ~0.60 0.10 ~0.60
≤ 0.035
0.015 ~0.030 ≤ 0.035
–
–
–
–
–
–
≤ 0.035
≤ 0.035
–
–
–
–
–
–
≤ 0.030
≤ 0.030
–
–
–
–
–
–
≤ 0.030
≤ 0.030
–
–
–
–
–
–
≤ 0.030
≤ 0.030
–
–
–
–
–
–
≤ 0.030
≤ 0.010
–
≤ 0.030
≤ 0.030
–
≤ 0.030
≤ 0.030
–
≤ 0.020
≤ 0.010
≤ 0.40
0.30 ~0.60 0.30 ~0.70 0.30 ~0.60
1.00 ~1.50 – 1.90 ~2.60 1.00 ~1.50 1.90 ~2.60 1.90 ~2.60 1.90 ~2.60
0.45 – ~0.65 0.87 ~ – 1.13 0.45 – ~0.65 0.87 – ~1.18 0.87 – ~1.18 0.05 0.20 ~0.30 ~0.30
8.00 ~10.00 8.00 ~10.00 8.00 ~9.50
0.90 – ~1.10 1.80 ~ – 2.20 0.85 ~ 0.18 1.05 ~0.25
WPNL2204
0.02 ~0.08
≤0.03
≤0.03
1.45 ~ 1.75
0.0006
–
–
–
–
–
–
–
–
–
–
–
–
0.06 ~0.10
≤0.04
0.030 ~0.070
165
≤ 0.35
Mn
Specifications for creep-resistant steels: Japan
≤ 0.32
C–Steel STB410
Low alloy steel
Si
166
Steels
Chemical composition (mass%) C
Mn
P
S
Ni
Cr
Mo
V
Nb
Al
N
W
B
0.030 ~0.070
1.50 ~2.00
0.001 ~0.005
0.07 ~0.13
≤ 0.50
0.30 ~0.60
≤ 0.020
≤ 0.010
≤ 0.40
8.00 ~9.50
0.03 ~ 0.60
0.15 ~0.25
0.04 ~0.09
≤0.04
KA-SUS410J2TB
≤ 0.14
≤ 0.50
≤ 0.030
≤ 0.030
–
≤ 0.50
≤ 0.020
≤ 0.010
≤ 0.50
≤ 0.70
≤ 0.020
≤ 0.010
≤ 0.50
0.20 ~0.30 0.15 ~0.30 0.15 ~0.30
0.04 ~0.10 0.04 ~0.10
≤0.04
≤ 0.50
0.80 ~1.20 0.25 ~0.60 0.25 ~0.60
≤0.04
0.07 ~0.14 0.07 ~0.14
11.00 ~13.00 10.00 ~11.50 11.51 ~ 12.50
≤0.20
KA-SUS410J3TB
0.30 ~0.70 ≤ 0.70
KA-STBA29
12Cr
Si
KASUS410J3DTB
WPNL2204
≤0.04
– 0.040 ~0.100 0.040 0.100
0.80 ~1.20 1.50 ~2.50 ~1.50 ~2.50
– 0.0005 ~0.005 0.0005 ~0.005
Creep-resistant steels
Table 4.3 Continued
Table 4.4 Chemical compositions for austenitic boiler tubing steels Steels
Chemical composition (mass%) C
SUS304HTB
~0.10 0.04 ~0.10 SUS347HTB 0.04 ~0.10 KA–SUS304J1HTB 0.07 ~0.13 KA–SUS309J1TB ≤ 0.06 SUS316HTB
Mn
P
S
KA–SUS309J4HTB
8.00 18.00 ~11.00 ~20.00 ≤0.75 ≤ 2.00 ≤0.030 ≤ 0.030 9.00 17.00
≤0.75 ≤ 2.00 ≤0.030 ≤ 0.030 ≤1.00 ≤ 2.00 ≤0.030 ≤ 0.030 ≤0.30 ≤ 1.00 ≤0.040 ≤ 0.010 ≤1.50 ≤ 2.00 ≤0.040 ≤ 0.030
≤ 0.04
≤1.00 ≤ 2.00 ≤0.040 ≤ 0.030 ≤1.50 ≤ 2.00 ≤0.030 ≤ 0.030
KA–SUS310J2TB
≤ 0.10
≤1.00 ≤ 1.50 ≤0.030 ≤ 0.010
0.05 ~0.12 KA–SUS321J1HTB 0.07 ~0.14
≤1.50 ≤ 2.00 ≤0.030 ≤ 0.010 ≤1.00 ≤ 2.00 ≤0.040 ≤ 0.030
~13.00 11.00 ~14.00 9.00 ~13.00 7.50 ~10.50 12.00 ~16.00 12.50 ~15.50 13.00 ~16.00 14.50 ~16.50 17.00 ~23.00 22.00 ~28.00 15.00 ~22.00 9.00 ~12.00
~20.00 16.00 ~18.00 17.00 ~20.00 17.00 ~19.00 23.00 ~26.00 21.00 ~23.00 23.00 ~26.00 21.00 ~23.00 23.00 ~27.00 19.00 ~23.00 21.00 ~24.00 17.50 ~19.50
Mo
V
Nb Al
N
Cu
W
B
Others
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
Ta: 4 × C ~ 0.60
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
Nb+Ta: 8 × C ~1.00 –
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
–
≤ 0.20
–
–
–
–
–
≤ 0.20
–
0.50 ~0.80 0.20 ~0.60 0.10 ~0.40 0.30 ~0.60 ≤ 0.40
2.00 ~3.00 – – 0.50 ~1.20 1.00 ~2.00 0.50 ~1.20 – – 1.00 ~2.00 – –
WPNL2204
0.30 ~0.60 –
0.05 ~0.12 0.25 ~0.40 0.10 ~0.25 0.25 ~0.40 0.10 ~0.20 0.15 ~0.35 0.10 ~0.25 0.15 ~0.30 –
2.50 ~3.50 –
≤ 0.006 –
0.002 ~0.010 2.00 0.80 – ~4.00 ~2.80 – – –
– – – – (Ti+Nb/2)C 2.0~4.0
167
KA–SUS310J1TB
0.03 ~0.10 ≤ 0.10
KA–SUS310J3TB
Cr
≤0.75 ≤ 2.00 ≤0.040 ≤ 0.030
≤1.00 2.50 ≤0.030 ≤ 0.030 ~3.50 KA–SUS309J3LTB ≤ 0.025 ≤0.70 ≤ 2.00 ≤0.040 ≤ 0.030 KA–SUS309J2TB
Ni
Specifications for creep-resistant steels: Japan
SUS321HTB
0.04 ~0.10 0.04
Si
168
Steels
Chemical composition (mass%) C
KA–SUS321J2HTB
0.07 ~0.14 KA–SUS3TP347HTB 0.04 ~0.10 KA–SUS347J1TB ≤ 0.05
Si
Mn
P
≤ 1.00
≤ 2.00
≤0.040 ≤ 0.010
≤ 0.75
≤ 2.00
≤0.030 ≤ 0.030
≤ 1.00
≤ 2.00
≤0.040 ≤ 0.030
S
Ni
Cr
9.00 17.50 ~12.00 ~19.50 9.00 17.00 ~13.00 ~20.00 8.00 17.00 ~11.00 ~20.00
Mo V
Nb
Al
–
–
0.10 – 2.50 – ~0.45 ~3.50 8 X C% – – – ~1.00 0.25 0.10 – 1.50 ~0.50 ~0.25 – ~2.60
– –
WPNL2204
0.10 ~0.25 – –
– 0.20 ~0.50
N
Cu
W
B
Others
0.0010 (Ti+Nb/2)C ~0.0040 2.0~4.0 – – –
–
Creep-resistant steels
Table 4.4 Continued
Specifications for creep-resistant steels: Japan
169
of JIS and METI Code steel tubes used extensively in high temperature components such as economizers, water walls, superheaters and reheaters in Japanese power boilers. The heat treatment conditions and tensile requirements for these steels are indicated in Tables 4.2–4.4. The specific designations for product forms other than boiler tubing and piping are shown in comparison with ASME Code in Table 4.5 for recently developed high strength ferritic steels.
4.4
Specifications for steam turbine steels
Table 4.64 shows the chemical compositions for steam turbine steels used in Japanese power plants. Steels are classified for specific components such as blades and discs, rotors and casings. The rotor steels shown here are recently developed 12%Cr high strength materials, although CrMo/CrMoV rotor steel and casings are still generally used for conventional steam turbines. Specifications for these turbine steels are controlled by turbine manufactures individually and not by the regulatory specifications. The technical control of melting and refining is also very important for turbine steels, as well as heat treatment control in order to maintain appropriate mechanical (strength/ toughness) and creep/fatigue properties. Thus, the listed nominal chemical compositions and heat treatment conditions in the table are suggested as guidelines.
4.5
Heat-resistant super alloys
Table 4.7 shows the chemical compositions of heat-resistant alloys (known as super alloys) listed in JIS G4901 for bars and G4902 for plates. These alloys are designated similarly to NCF XXX, with the lower three columns of the unified numbering system (UNS) number or ASTM designation. These alloys are annealed or solution annealed, except several grades that are aged after solution annealing. JIS does not specify the recommended annealing or solution temperature and aging temperature, although tensile/yield strength, rupture elongation and hardness are specified. These alloys exhibit extremely high creep strength and high temperature corrosion/oxidation resistance compared with other steels.
4.6
Summary
Specifications for creep-resistant steels in Japan are well established in JIS for ordinary use including high temperature applications, with METI Codes and JSME Codes only for boiler and nuclear applications. These specifications show general recommendations and regulations for chemical compositions, heat treatment conditions, mechanical properties, and so on. Mechanical
WPNL2204
170
Steels (trade name) 2.25Cr
HCM2S HCM9M
9Cr
Mod.9Cr NF616 HCM12
12Cr
HCM12A
Code
Tube
Pipe
Plate
ASME METI ASME METI
SA-213 T23 KA-STBA24J1 – KA-STBA27
SA-335 P23 KA-STPA24J1 – KA-STPA27
SA-1017 Gr.23 KA-SCMV4J1 – –
SA-182 F23 KA-SFVAF22AJ1 – KA-SFVAF27
ASME METI ASME METI ASME METI
SA-213T91 KA-STBA28 SA-213T92 KA-STBA29 – KA-SUS410J2TB
SA-335P91 KA-STPA28 SA-336P92 KA-STPA29 – –
SA-387Gr.91 KA-SCMV28 – – – –
SA-182F91 KA-SFVAF28 SA-182F98 KA-SFVAF29 – –
ASME METI
SA-213T122 KA-SUS410J3TB KA-SUSF410J3DTB
SA-335P122 KA-SUS410J3DTP –
SA-1017Gr.122 KA-SUS410J3 –
SA-182F122 KA-SUS410J3 –
WPNL2204
Forgings
Creep-resistant steels
Table 4.5 Comparative code designations in METI Code and ASME for new ferritic boiler steels
Table 4.6 Chemical compositions and heat treatment for turbine steels Steels
Nominal chemical composition (mass%) —————————————————————————————————————— Ni
Cr
Mo
W
Co
V
Nb
B
N
Small H46 (12Cr–0.5MoVNbN) 0.15 components TAF (10.5Cr–1.5MoVNbB) 0.18 TAF650 0.11 (11Cr –2.6W–3CoVNbB) TOS203 0.11 (10.5Cr –2.5W–1CoVNbBRe)
0.40 0.30 0.01
0.60 0.50 0.50
– – –
12.0 10.5 11.0
0.5 1.5 0.2
– – 2.6
– – 3.0
0.30 0.20 0.20
0.25 0.15 0.08
– 0.030 0.015
0.050 0.015 0.020
1,150 –650 1,150 –700 1,100 –750
0.05
0.50
0.6
10.5
0.1
2.5
1.0
0.20
0.10
0.010
0.030
1,120 –680
Rotor
GE (10.5Cr–1MoVNbN) 0.18 TMK1 (10.3Cr–1.5MoVNbN) 0.14 HR1100 0.15 (10.3Cr–1.2Mo–0.3WNbN) TOS107 0.14 (10Cr–1Mo–1MVNbN) TMK2 0.14 (10.2Cr–0.5Mo 1.8WVNbN) TR1200 0.13 (11Cr–0.2W–2.5WVNbB) HT1200 0.11 (11Cr–2.6W–3CoNiVNbB) TOS110 (10Cr–0.7Mo– 0.11 1.8W –3CoVNbB)
0.30 0.05 0.03
0.60 0.50 0.60
0.6 0.6 0.6
10.5 10.3 10.3
1.0 1.5 1.2
– – 0.3
– – –
0.20 0.17 0.20
0.06 0.06 0.05
– – –
0.060 0.040 0.050
1,050 –620 1,100 –680 1,075 –660
0.05
0.60
0.7
10.0
1.0
1.0
–
0.20
0.07
–
0.050
1,050 –660
0.05
0.50
0.5
10.2
0.5
1.8
–
0.17
0.06
–
0.040
1,050 –700
0.05
0.50
0.8
11.0
0.2
2.5
–
0.20
0.08
–
0.050
1,050 –710
0.05
0.50
0.5
11.0
0.2
2.6
3.0
0.20
0.08
0.015
0.025
1,050 –720
0.08
0.10
0.2
10.0
0.7
1.8
3.0
0.20
0.05
0.010
0.020
1,070 –680
MJC12 (9.5Cr –1MoVNbN) 0.10 TOS302 0.12 (10Cr –1Mo–0.8WVNbN) TOS303 0.12 (10Cr–1.8W–3CoVNbB)
0.70 0.25
0.70 0.50
0.5 1.0
9.5 10.0
1.0 1.0
– 0.8
– –
0.15 0.20
0.06 0.10
– –
0.040 0.050
1,050 –710 1,050 –710
0.15
0.50
0.2
10.0
0.7
1.8
3.0
0.20
0.05
0.006
0.020
1,100 –730
Casing
WPNL2204
171
Mn
C
Specifications for creep-resistant steels: Japan
Si
Heat treatment Nor.-temper (°C)
172
Table 4.7 Chemical compositions of heat-resistant alloys (JIS G4901 and G 4902) Chemical composition (mass%) C
Si
Mn
P
S
NCF600
≤ 0.15
≤ 0.50
≤ 1.00
≤ 0.030
≤ 0.015
72.00~
NCF601
≤ 0.10
≤ 0.50
≤ 1.00
≤ 0.030
≤ 0.015
NCF625
≤ 0.10
≤ 0.50
≤ 0.50
≤ 0.015
≤ 0.015
58.00 ~63.00 58.00~
NCF690
≤ 0.05
≤ 0.50
≤ 0.50
≤ 0.030
≤ 0.015
58.00~
NCF718
≤ 0.08
≤ 0.35
≤ 0.35
≤ 0.015
≤ 0.015
NCF750
≤ 0.08
≤ 0.50
≤ 1.00
≤ 0.030
≤ 0.015
50.00 ~55.00 70.00~
NCF751
≤ 0.10
≤ 0.50
≤ 1.00
≤ 0.030
≤ 0.015
70.00~
NCF800
≤ 0.10
≤ 1.00
≤ 1.50
≤ 0.030
≤ 0.015
NCF800H
0.05 ~0.10 ≤ 0.05
≤ 1.00
≤ 1.50
≤ 0.030
≤ 0.015
≤ 0.50
≤ 1.00
≤ 0.030
≤ 0.015
≤ 1.00
≤ 1.00
≤ 0.030
≤ 0.015
NCF825 NCF80A
0.04 ~0.10
Ni
30.00 ~35.00 30.00 ~35.00 38.00 ~46.00 Bal.
Cr
Fe
14.00 ~17.00 21.00 ~25.00 20.00 ~23.00 27.00 ~31.00 17.00 ~21.00 14.00 ~17.00 14.00 ~17.00 19.00 ~23.00 19.00 ~23.00 19.50 ~23.50 19.50 ~23.50
6.00 ~10.00 Bal.
–
≤ 0.50
–
–
–
–
–
≤ 1.00
–
–
–
≤ 5.00
8.00 –10.00 –
–
1.00 ~1.70 ≤ 0.40
≤ 0.40
≤ 0.50
–
–
≤ 0.30
0.20 ~0.80 0.40 ~1.00 0.90 ~1.50 0.15 ~0.60 0.15 ~0.60 ≤ 0.20
Bal. = balance.
WPNL2204
7.00 ~11.00 Bal. 5.00 ~9.00 5.00 ~9.00 Bal. Bal. Bal. ≤ 1.50
Mo
2.80 ~3.30 –
Cu
≤ 0.50
–
≤ 0.50
–
≤ 0.75
–
≤ 0.75
2.50 ~3.50 –
1.50 ~3.00 ≤ 0.20
Al
1.00 ~1.80
Ti
Nb + Ta
0.65 ~1.15 2.25 ~2.75 2.00 ~2.60 0.15 ~0.60 0.15 ~0.60 0.60 ~1.20 1.80 ~2.70
B
3.15 ~4.15 –
–
4.75 ~5.50 0.70 ~1.20 0.70 ~1.20 –
≤0.006
–
– – –
–
–
–
–
–
–
Creep-resistant steels
Designation
Specifications for creep-resistant steels: Japan
173
properties, including hardness, tensile, impact, creep and fatigue properties, are strongly influenced by microstructures resulting from heat treatment and chemical composition. In the case of turbine steels, the steel manufacturing process is very important to achieve the desired properties and performance. In the case of boiler/piping and pressure vessel applications, allowable stresses are provided in the METI Codes and JSME Codes, which feature criteria similar to ASME for determination of the allowable stress values for temperatures. In the future, specifications for creep-resistant steels will be updated in accordance with codes and standards developed in other areas worldwide.
4.7 1 2 3 4
References
http://www.jsa.or.jp/default_english.asp http://www.meti.go.jp/english/index.html http://www.jsme.or.jp/English/ Masuyama F. ‘History of power plants and progress in heat resistant steels’, ISIJ International, 2001 41, 612–625.
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5 Production of creep-resistant steels for turbines Y. T A N A K A, Japan Steel Works, Japan
5.1
Introduction
Increases in steam temperature and pressure have made a large contribution to improvement in the efficiency of fossil power plants. The importance of the increase in efficiency is a stringent problem for CO2 reduction and conservation of resources. Realisation of advanced fossil power plants that incorporate advanced steam conditions depends on the development of advanced heat-resistant steels and of production technology for large turbine rotor forgings. So far, many advanced type high pressure (HP) turbines, intermediate pressure (IP) turbines and high pressure low pressure combination turbines (HLP) have been developed successfully and contribute to an increase in plant efficiency. These advanced turbines are realised through the application of state-of-the-art production technology developed together with cultivated widely used and conventional technologies over the years. Points for the production of high-performance and reliable turbine rotor forging can be summarized as follows. Production of high purity steels with minimized residual elements and freedom from non-metallic inclusions is important in the steelmaking process. Homogeneous ingots with minimal segregation, delta ferrite, non-metallic inclusions and porosities should be made with a homogeneous distribution of chemistry throughout the casting process. For each forging process, a sufficient forging effect at the centre of large diameter in the ingots and forging blocks needs to be attained to consolidate the porosities in the ingots and sufficient forging strain should be given to the ingots to eliminate the solidification structure (e.g. dendrite) and promote formation of equiaxed grain through dynamic recrystallization. In the heat treatment process, heat treatment effects need to be exerted to develop the required properties at the centre of the forging. A fine grain microstructure needs to be obtained to assure sufficient detectability of defects. In this chapter, the development of production technologies for turbine rotor forgings is reviewed and several key processes are introduced. In 174 WPNL2204
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addition, the properties of forgings made from heat-resistant steels are introduced.
5.2
Overview of production technology of rotor shaft forgings for high temperature steam turbines
Figure 5.1 shows a schematic of the production process for large turbine rotor forgings. The detail of each process is essentially different depending on the designation of steels and the required properties of the rotor shaft. In this section, the facilities for the production and typical technologies are overviewed.
5.2.1
Steelmaking and casting process
Figure 5.2 shows the historical change in steelmaking methods and equipment in one company since 1950, as an example.1 Similar changes in facilities and production technologies have taken place in the majority of the forgemasters in the world. Turbine material used to be refined by open hearth furnaces and was cast in air. At that time, the absorption of hydrogen in steel was one of the most serious problems of the process since hydrogen causes defects such as flaking. Installation of vacuum degassing equipment like the Bochumer–Verein type was a solution for the degassing of hydrogen during casting. By use of vacuum degassing equipment, the basic open hearth furnace and the basic electric arc furnace (EAF), which tend to absorb hydrogen during refining but are superior in refining ability, were able to be used, leading to an improvement in the material properties. The efficiency of the vacuum degassing equipment has been improved to attain a higher vacuum. With the introduction of vacuum casting equipment, vacuum carbon deoxidization (VCD) technology has been successfully applied to steam turbine materials.2 With the increase in capacity of power plants, larger turbine rotor forgings were needed and new production technologies were developed. A pouring method using multiple furnaces was developed as the casting technology for such large forging. Ladle refining furnaces (LRF) were installed to keep steel molten after refining in an electric furnace. After that, several ladle refining furnaces were additionally installed and these make the production of ingots up to 600 tonnes possible using a fully ladle refined melt.3 On the other hand, an electroslag remelting (ESR) furnace was used for refining and casting of high quality steels. The capacity was enlarged to meet the demands of melting large high temperature turbines made from low alloy and high alloy steels. Vacuum induction melting (VIM) was also installed. The use of VIM for the production of rotor forging, however, is uncommon. Through
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Electric arc furnace
Press
Vacuum
Horizontal furnace
Vacuum
Steelmaking casting 1
2
Forging
3
Preliminary heat treatment
Vertical furnace Quenching
Rough machining
Vertical furnace Tempering
WQ Oil FC Quality heat treatment 5
4
FC Stress relieving Stress relieving
NDE, machining 6
7
5.1 Typical production process for large steam turbine forging.
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Ladle refining furnace
Vacuum casting
Year
1950
55
60
65
70
75
80
85
90
95
Acid and basic open hearth furnace Electric furnace Ladle refining furnace
Melting and refining
Holding furnace
150 ton 130 ton
30 ton x2
20 ton ESR
150 ton x2 100 ton ESR 5 ton VIM
Bochumer–Verein type mould stream degassing Ingot making
Air casting Mechanical pump
Steam ejector (high vacuum) Multipouring
Maximum ingot weight
140 ton
220 ton
250 ton
Low Si-VCD
Prediction of segregation
Double degassing
500 ton
570 ton
600 ton
400 ton
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5.2 The history of the steelmaking and casting process: an example.
Production of creep-resistant steels for turbines
Basic open hearth furnace Electric furnace
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Creep-resistant steels
the development of these facilities and production technologies, high performance and high reliability rotor forgings have been manufactured. Current steelmaking processes used in the production of high temperature steam turbine rotor forging are as follows. Basic electric arc furnace and ladle refining furnace Formerly, acid and/or basic open hearth furnaces were used for refining steels. Since the introduction of vacuum degassing facilities, basic electric arc furnaces have become the major equipment used in the melting and refining of turbine rotor steels. In this process, refining is performed by the double slag method. After melting the raw materials, oxidizing refining is effected by adding a basic oxidizing slag to reduce the C, Mn, Si and P content. Then the oxidizing slag is removed to avoid the oxidized elements returning to the molten steel. When removing oxidizing slag, it is especially important to remove the P, which is harmful to the material properties. Then reducing slag is added to decrease the S content. After adjustment of the chemistry, the molten steel is poured into a mould in the vacuum degassing chamber. In order to reduce the residual elements as low as possible to make high purity steels, the double slag process was further improved by using EAF and subsequently LRF. Figure 5.3 shows an example of a typical current steelmaking process using EAF and LRF.3 The raw materials are melted in an EAF where oxidizing refining is performed. Then the molten steel is poured into a ladle and the oxidizing slag in the EAF is completely removed. After reladling, the reduction refining and degassing processes take place in
Electric furnace Ladle furnace Vacuum Vacuum
Ar gas
Melting/ refining
Reladle
De-phosphorisation
Ar gas
Ladle refining
Ingot making
De-sulphurisation
5.3 Typical steelmaking process with EAF and LRF for high purity steel.
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a vacuum after agitation with Ar gas. Subsequently, the molten steel is cast in vacuum through mould stream degassing. Figure 5.4(a), (b) and (c) shows the appearance of tapping from the EAF to the ladle, ladle refining and casting from ladle to ingot mould in the vacuum chamber, respectively. The use of multiple ladles make it possible to cast a large ingot. Figure 5.5 shows the history of P and S content in materials for LP (low pressure steam turbine) rotor forging steel. The significant effect of advances
(a)
(b)
(c)
5.4 The steelmaking and casting process. (a) Tapping from an electric arc furnace, (b) ladle refining, (c) vacuum pouring. 0.024 0.022
: S% : P%
P and S contents (wt%)
0.020 0.018 0.016 0.014 0.012 0.010 0.008 0.006 0.004 0.002 0
1960
1970
1980 Year
1990
2000
5.5 Historical change of P and S content in steels for LP rotor forging.
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Creep-resistant steels Table 5.1 Elimination of tramp elements for clean steel Element
Technology and process to be applied
Mn Si P Si Sn As Sb Al H O N
Oxidizing refining in EAF → reladle Oxidizing refining in EAF and VCD Oxidizing refining in EAF → reladle Ladle refining Selection of raw material Selection of raw material Selection of raw material VCD Vacuum treatment Ladle refining and VCD Ladle refining
Consumable electrode Furnace Cooling water Crucible Slag pool
Consumable electrode
Vacuum
Cooling water
Molten steel pool
Molten steel pool Crucible Ingot
Ingot Cooling water
Cooling water (a)
(b)
5.6 Equipment for (a) ESR process and (b) VAR process.
in refining technology is demonstrated by the reduction of P and S. Table 5.1 summarizes the methods of reducing the impurity elements in current steelmaking processes and currently superclean steels has been developed by reducing not only residual elements such as P, S, As, Sn, Sb but also Si and Mn which are usually added to effect deoxidization.3–5 Electroslag remelting (ESR) The ESR equipment consists of a large capacity power supply and a water cooled crucible. Figure 5.6(a) shows a schematic of the ESR process and Fig. 5.7 shows an ESR facility. In the ESR process, an electrode is prepared by casting or forging after the conventional melting, refining and casting process. The melting of the electrodes occurs in the mould by heating caused by the electric resistance of the slag. As the droplets of electrode material
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5.7 An ESR facility.
fall, refining of the material proceeds. The material solidifies at the bottom of the molten material pool. By application of the ESR process, the purity, cleanliness of the material and homogeneity of the ingot can be improved. Formation of macro segregation is suppressed and the distribution of chemical elements becomes more uniform compared to ingots made by conventional casting. Slag composition is especially important in the ESR process in order to attain the expected properties. Presently, the ESR process is frequently used for the production of rotor forgings made from heat-resistant steels such as CrMoV steels, advanced CrMoV steels and various 12Cr steels. Hot topping process Several applications of ESR technology combined with a conventional casting method such as the Bohler electroslag topping process (BEST)6 and
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electroslag hot topping (ESHT-J)7 have been developed. In both methods, refining is performed by a conventional process in an EAF and the product is poured into an ingot mould. The slag is fed on to the surface of the molten steel in the mould and heating by the consumable electrode (together with a graphite electrode) takes place. By using these processes, the formation of porosity, shrinkages and segregation can be suppressed. Vacuum arc remelting (VAR) Figure 5.6(b) shows a schematic of VAR equipment. The electrode is arc melted in a water-cooled crucible inside a vacuum chamber. In contrast to the ESR process, molten steel refining does not proceed unless there is degassing through this process. Therefore, the electrode for the VAR process should be produced through a conventional refining and casting process or by vacuum induction heating process so as to attain the target chemistry and purity. On the other hand, cooling of the molten steel takes place at a faster rate than during the ESR process for a given mould dimension, resulting in a superior solidification structure with less segregation. The capacity of VAR is comparatively small owing to installation of a vacuum chamber; experience of application to the production of rotor steel forging is therefore limited. Vacuum induction melting (VIM) VIM uses a melting furnace for raw materials by induction heating in a vacuum. Since refining does not result from the VIM process, the raw material should be high purity ferroalloys and high purity metals, depending on the requirement of the products. Although application to large steel forging products is quite limited, the process coupled with VAR/ESR processes is indispensable to the production of super alloys coupled with VAR and/or ESR process.
5.2.2
Ingot making process
After the conventional melting and refining process, the molten steel is cast into an ingot mould from top or bottom of a mould made from cast iron. In the case of ingots for large steam turbine forging, top pouring into the ingot placed in vacuum tank is performed, as shown in Fig. 5.3. During pouring, a further reduction in the gas elements proceeds in the vacuum tank. For smaller rotor forgings, casting of ingots weighing several tens of tons of steel is done by bottom pouring. In this process, the refined molten steel is poured into the ingot mould from bottom of the mould set on a steel plate through refractory tube until the mould fills up. In the case of ESR and VAR, the re-melted steel is again solidified into ingots in a water-cooled crucible.
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In order to reduce segregation in the ingot and to improve the performance of the material, a low Si-VCD process is applied.2 In the VCD process, molten steel with a low Si content is cast in a high vacuum; C and O in the steel react and are exhausted as CO. During application of the VCD process, it is possible to achieve much lower levels of Si than in the Si deoxidation process. Decrease in Si content not only develops a fine solidification structure which suppresses formation of macro-segregation in large ingots but also decreases the susceptibility to temper embrittlement which is critical problem for materials used in the temperature range between 350 and 550 °C.8,9 A decrease in Si is also preferable for increasing creep and creep rupture strength. Figure 5.8 shows sulphur prints of the axial sections of rotor forging, although the material is NiCrMoV steel in an LP turbine, manufactured by a conventional Si-deoxidization and VCD method. The segregation streak in the forging made during the VCD process has almost disappeared while that in Si deoxidized forging is clearly observed. Low Si-VCD technology has been applied in the production of rotor forging not only for LP turbines but also for high temperature turbines. The design of the dimensions of the ingot is important, to reduce defects and inhomogeneities such as porosities and macro-segregation. The importance is increased for large ingots. The ratio of height to diameter (H/D), taper of the ingot and hot top design are major affecting factors. It is considered that an H/D of less than 1 is preferable to reduce the porosity size.10 A quality can be obtained equivalent to ingots prepared by ESR by reducing the H/D to 0.75 for the case of 12Cr steel. Nowadays, solidification behaviour can be predicted by the development of numerical simulation technology contributing to the manufacture of high quality ingots.
5.2.3
Forging process
The consolidation of porosities formed during solidification and homogenization of the material are the major aim of the initial stage of the forging. Then the material is forged to form the shape of products. Specific forging processes have been developed and applied to exert the optimum forging effects. These lead to sound and desirable properties from forging. Figure 5.9 shows an open die forging press. To consolidate porosities in the large ingot, a strong forging effect in the centre of a large ingot is required and many processes which optimize the forging temperature, shape and dimension of the dies, pressing sequence, and so on, have been developed and applied. For example, in the warm forging process, by cooling a uniformly heated material to a surface temperature of around 800°C, forging develops a temperature difference between the interior and surface of the material. As the result of the difference in flow stress between the interior and outer portions of the material, the process gives strong consolidation effect. Recently
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Bottom
500 mm
0
Dec. 1975
(a) C
Si
Mn
P
S
Ni
Cr
Mo
V
0.22
0.02
0.31
0.006
0.008
3.47
1.79
0.41
0.12
Bottom
Top
0
500 mm
Sep. 1952
(b) C
Si
Mn
P
S
Ni
Cr
Mo
V
0.34
0.38
0.51
0.021
0.008
3.60
0.19
0.43
0.15
5.8 Cross-sections of rotor forgings cast in air and cast through a VCD process. (a) NiCrMoV LP turbine rotor shaft from 140 ton VCD ingot; (b) acid open hearth furnace air cast 75 ton ingot. WPNL2204
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5.9 A free forging press.
the application of the finite element method (FEM) analysis to the problems of plastic deformation have become general and applied widely to simulate the effect of hot working and to determine the forging process. An example of a forging process for CrMoV rotor forging is shown in Fig. 5.10. Generally, the forging process consists of several steps. The optimum heating temperature for each hot working step is determined by considering the dynamic recrystallization behaviour of the material, its resistance to hot working, the grain growth behaviour, the diffusion effect of inhomogeneities such as segregation, and so on. Before the forging operation, generally, the hot top and bottom side of the ingot are discarded to remove the portion that contains heavy segregation and non-metallic inclusions. In the early stage of the forging process, upsetting is performed. The ingot height is reduced by upsetting and the diameter is increased to improve homogeneity and to increase the forging ratio. The forging ratio generally required to develop a homogeneous microstructure is around 3 for a conventional ingot. In case of an ESR ingot, a lower forging ratio is acceptable owing to its inherently
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Sketch of process Discard
Ingot
Making handling stem
Upsetting rounding
Finishing the body
Discard
Discard
Finishing the journal
5.10 An example of the forging process for a CrMoV rotor shaft.
(a)
(b)
5.11 (a) Upsetting and (b) finish forging.
good solidification structure. Dies and hot working steps are carefully designed to exert the largest forging effect. Figure 5.11(a) and (b) shows upsetting and finish forging.
5.2.4
Heat treatment
The role of heat treatment is not only in the development of target mechanical properties such as strength, toughness and creep strength in forging but also
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the formation of microstructure with good inspectability and heat stability. The level of performance of the forging that is achieved is defined by the fine and uniform microstructure. The overall process of heat treatment involves several heating steps and largely depends on the designation of the material. The difference in heating rate, cooling rate and holding time from the surface to the centre of the large diameter forging need to be considered in order to attain the target properties. Heat treatment of turbine rotor forgings consists of preliminary heat treatment which is first performed after forging and quality heat treatment which is performed subsequently. Stress relief heat treatment is also conducted after quality heat treatment. Preliminary heat treatment After the forging process, preliminary heat treatment is performed aimed at the relaxation of strain introduced by hot working and refining the coarse grain formed during the forging process. Since, generally, it is difficult to develop a small grain structure in large forgings through dynamic recrystallization by hot working, preliminary heat treatment is important in developing a fine grain microstructure that exhibits toughness and inspectability by an ultrasonic test (UT). A typical heat treatment that aims to refine coarse grain is a normalizing treatment and an alternative is heat treatment applying pearlite transformation. Figure 5.12 shows a schematic of normalizing heat treatment and pearlite transformation heat treatment. In the normalizing process, the material is T1 T2
AC3 AC1
(a) T1
T3
AC3 AC1
T3 : Temperature for pearlite transformation (b)
5.12 Two typical heat cycles for preliminary heat treatment after forging. (a) Normalizing and tempering heat treatment; (b) pearlite transformation heat treatment.
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Creep-resistant steels
cooled once after finish forging to develop the ferritic microstructure and then heated again to greater than the temperature at which austenite transformation (AC3) is completely finished. When the surface portion of the forging reaches the austenitizing temperature, the material is kept for enough time for the temperature at the centre of the forging to reach the austenitizing temperature. Setting the austenitizing temperature for normalizing is important in grain refining since certain types of material require a comparatively higher austenitizing temperature than AC3 to complete recrystallization as small austenite grains. Grain refining behaviour during austenite transformation may also be affected by the heating rate. Higher heating rates tend to develop finer austenite grain. However, in the large forging, the heating rate that is attainable at the centre is small. For forgings where there is difficulty in grain refining, depending on the designation of material and dimensions of the forging, normalizing heat treatment is repeated, reducing the grain in each treatment. Pearlite transformation is also a measure for grain refining. Pearlite transformation proceeds during cooling from the austenitizing temperature and the temperature and time for completion of pearlitic transformation depend on the chemistry of the material. The material needs to be kept at around the temperature of pearlite transformation nose of the time-temperature transformation (TTT) diagram. Since materials for rotor forging require good hardenability making it difficult to proceed with pearlite transformation, a considerable holding time may often be required to complete the pearlite transformation. After the completion of pearlite transformation, the material is again austenitized to develop a fine austenite grain. Quality heat treatment In the case of turbine forging materials, the target properties are developed by quenching heat treatment followed by tempering heat treatment. Quenching is heat treatment by cooling from the austenitizing temperature which is commonly selected at a temperature that dissolves the carbides in steels and develops the desired material properties such as creep strength. Attention should be given, however, to the avoidance of excessive grain coarsening by heating. In order to attain high strength and toughness for rotor forging, microstructures transformed at low temperatures like martensite and lower bainite are preferable. The formation of ferrite and pearlite, however, should be avoided in order to develop good toughness and creep strength. In large forgings, the cooling rate deep inside the forging is significantly slower than that of the surface region. Attention should be focused on developing the desired microstructures even at the centre of the forging. The microstructure of a large forging after quenching can be estimated by referring to the continuous
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cooling transformation (CCT) diagram of the material corresponding to the cooling rate during quenching in each portion of the forging. Figure 5.13 shows examples of CCT diagrams of quenching for a CrMoV steel and a 12CrMoV steel. Water quenching and water spray quenching are commonly applied in order to achieve sufficient cooling to develop a martensitic or bainitic microstructure throughout the whole volume of the large diameter forging. A 1000
Temperature (°C)
800 Ferrite 600
400
Bainite
200
HV 418 381 102
103
377 344
104 Time (s) (a)
340 335 105
1000
Temperature (°C)
800 Ferrite 600
400 Martensite 200 HV 102
488 489 433 103
450 370 363 362 269 228
104 Time (s) (b)
105
5.13 Examples of continuous cooling transformation (CCT) diagrams for (a) CrMoV steel at an austenizing temperature of 970°C and (b) 12CrMoVNbN steel at an austenizing temperature of 1050°C. HV is the Vickers hardness valve after cooling to room temperature.
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rather mild cooling effect can be obtained by immersion into an oil medium. Oil quenching reduces the risk of quenching crack but may also reduce the strength and toughness of the material. Forced air cooling and air cooling are other methods for effecting a mild cooling rate. In the case of high temperature turbine forging material like CrMoV steels, forced air cooling is sometimes applied in order to attain higher creep strength.11 After quenching heat treatment, tempering heat treatment is performed consecutively to develop the target mechanical properties below the initiation temperature of austenite transformation (AC1) of the material. The properties are controlled by the tempering temperature and holding time. Higher temperature and longer holding time generally reduces strength and increases the toughness. Since ultimate safety is required for turbines, the forging must have homogeneous microstructures and inhomogeneity leading to vibration must be avoided. In order to attain a uniform heat treatment effect, quality heat treatment of rotor forging is commonly performed in a vertical furnace under rotation of the forging. Figure 5.14 shows a rotor forging heated in the vertical furnace. Stress relief heat treatment is performed after machining to reduce the residual stresses generated in the forging by the heat treatment and machining. The heating temperature is set at more than 30°C lower than the tempering temperature so that no change in mechanical properties occurs.
5.2.5
Machining
Machining is performed in several stages of the production of rotor shaft forging as shown in Fig. 5.1. After the preliminary heat treatment, generally, machining using a lath is performed to make a smooth surface for ultrasonic examination. The smooth surface also contributes a homogeneous heat treating effect in quality heat treatment. Scale from oxides formed during the forging operation, surface defects and decarburized surface are removed by machining. After quality heat treatment, machining for gashing and finish machining are performed. Sometimes the central bore is machined following requests from turbine builders/users for non-destructive examination (NDE) of the centre of the forging. Turbine builders often conduct further machining for blade attachment, and so on.
5.2.6
Metallurgical and mechanical tests and nondestructive examination (NDE)
After heat treatment, specimens are removed from the forging and subjected to metallurgical and mechanical tests. These specimens are typically several portions taken from the surface and both ends of the forging. Metallurgical and mechanical tests on the centre core are also performed for forgings with
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5.14 A rotor forging heat treated in a vertical furnace.
a central bore. The followings are typical features of metallurgical and mechanical test, although all of these are not necessarily performed for all rotor forgings. Each test is performed in accordance with standards such as ISO, ASTM, DIN, JIS, and so on. • • • • • • • • • •
chemical analysis macrostructure and sulphurprint microstructure cleanliness hardness test tensile tests (at room temperature and elevated temperature) Charpy impact test (absorbed energy, shear fracture) creep test creep rupture test low cycle and high cycle fatigue tests
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• •
Creep-resistant steels
fracture toughness tests fatigue crack growth tests.
In order to assure the quality of the forging, NDE is also performed. Surface examination includes visual examination, liquid penetrant examination and magnetic particle examination. Ultrasonic examination is performed as a volumetric examination to detect defects in the forging. In order to have good detectability of the defects, as already mentioned, a fine and homogeneous grain microstructure in the forging should be formed in order to reduce the noise level and to be able detect small indications of effects in the forging. After stress relief heat treatment, a heat indication test (HIT) is performed to confirm the heat stability of the forging.12 Rotor forging placed in a test furnace is heated from room temperature up to approximately operating temperature under rotation and then is cooled down to room temperature. The movement of the central axis of the forging during the heat cycle is measured precisely. Inhomogenity, such as asymmetry of the microstructure and the existence of a large residual stress, may cause large abnormal movement of the central axis which may lead to vibration in the turbine. The amount of movement is limited to the specified value.
5.3
Production and properties of turbine rotor forgings for high temperature applications
5.3.1
Production and properties of CrMoV steel rotor forging
CrMoV steel designated ASTM A470 Class 8 and DIN 30 CrMoNiV 5 11 are typical low alloy steels for high temperature rotor forging for HP/IP turbines. The material bears V and a fine precipitation of vanadium carbide develops good creep strength. The CrMoV rotor is generally used at steam temperatures up to 566°C. Although the high temperature creep strength of the material is appreciable, CrMoV steel is poor in fracture toughness. Considerable discussions took place between turbine builders to determine the optimum balance of toughness and creep strength and heat treatments in the development of the target properties.11 In the early 1950s in the USA, a notch sensitivity problem caused by extremely high austenitizing temperatures over 1000°C was disclosed. This sensitivity disappeared when the austenitizing temperature was decreased to 954°C. Air cooling from the austenitizing temperature has been adopted by turbine builders preferring a higher creep strength. Oil quenching is also applied to improve toughness by turbine builders for whom toughness is important. In order to evaluate the effect of advanced steelmaking technology on the properties of CrMoV rotor forging, the Electric Power Research Institute (EPRI) carried out a project to evaluate the performance of three rotor forgings made by three different advanced
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steelmaking and casting processes.13,14 The processes applied are electroslag remelting (ESR), low S and vacuum carbon deoxidization (VCD). These evaluation tests on the forgings suggest that improvement in production technology significantly contributes to the performance of the rotor forgings. On the other hand, improvement in the toughness by alloy modification was successfully attained.15,16 2CrMoNiWV steel was proved to have superior toughness and equivalent creep strength to that of conventional CrMoV steel. Table 5.2 shows the chemistry of CrMoV steel and 2CrMoNiWV steel for rotor forging. Manufacturing processes for CrMoV rotor forgings are not necessarily the same for all the forgemasters. The process and detailed conditions are designated depending on their facilities and technologies. Figure 5.15 shows an example of the production process for CrMoV steel rotor forging. In this case, the steel is refined in EAF and LRF. Then the ingot is cast and forging takes place followed by preliminary heat treatment. After preliminary heat treatment, the surface of the forging is machined for a uniform heat treatment effect in quality heat treatment. Then quality heat treatment, quenching and tempering are performed followed by more machining. Test specimens for the mechanical tests are removed at this stage. After machining and NDE, stress relieved heat treatment is performed followed by the heat indication test. In general, EAF and subsequent LRF refining followed by vacuum casting into a mould is a typical process widely applied for steelmaking. The ESR process is also applied as an alternative refining and ingot-making process. An example of the forging process of a conventional CrMoV rotor shaft is shown in Fig. 5.10. In this case, four hot working steps are performed to complete the forging process. Figure 5.16 shows an example of heat treatment of CrMoV rotor forging. After finish forging, normalizing and subsequent tempering are performed as the preliminary heat treatment. Quenching is done by forced air cooling from the austenitizing temperature of 950°C. Then tempering at 670°C and stress relief heat treatment at 640°C are performed. The microstructure of the forgings quenched by forced air cooling and oil quenching is essentially an upper bainitic microstructure from the surface to the centre of the forging. Increase of cooling rate may enhance the transformation to lower bainite which gives better toughness. Table 5.3 shows an example of mechanical and impact properties at the surface and centre region of the CrMoV rotor forgings with a maximum body diameter of approximately 1.2 m investigated in an EPRI advanced CrMoV rotor project.13,14 The tensile properties are homogeneous from the surface to the centre portion. The fracture appearance transition temperature (FATT) of the central region is slightly higher compared to that of the surface region. The
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Steel
C
Si
Mn
P
S
Ni
Cr
Mo
ASTM A470 Cl.8
0.25–0.36
<0.10
<1.0
<0.015
<0.018
<0.75
0.9–1.5
1.0–1.5
0.2–0.3
DIN 30 Cr MoNiV 5 11*
0.28–0.34
<0.10
0.30–0.80
<0.007
<0.007
0.70–0.80
1.1–1.4
1.0–1.2
0.25–0.35
Alloy 88 (22CrMoNiWV8–8)*
0.21–0.23
<0.10
0.65–0.75
<0.007
<0.007
0.50–0.75
2.05–2.15
0.80–0.90
* VCD
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0.60–0.70
V
0.25–0.32
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Table 5.2 Typical chemistry of CrMoV rotor steel, in wt%
Production of creep-resistant steels for turbines
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Steelmaking (EAF, LRF), casting
Forging
Preliminary heat treatment
Rough machining NDE (Ultrasonic test) Quenching (oil)
Tempering
Machining NDE, mechanical and metallurgical tests Stress relief HIT Shipping
5.15 Production sequence for a CrMoV rotor shaft.
2CrMoNiWV steel is reported to develop better toughness from surface to centre compared to the conventional CrMoV steel forging. Addition of Ni and increase of Cr content in this steel is beneficial to improve the hardenability leading to better toughness. Figure 5.17 shows an example of creep rupture test of CrMoV steels from oil quenched and forced air cooled forgings showing similar rupture strength for both quenching methods. The creep rupture strength of the 2CrMoNiWV steel forging is equivalent to that of CrMoV steel.16
5.3.2
Production and properties of 12Cr steels
12Cr steel was introduced into service in 1960 for high temperature rotor forging operating at 566°C but its use was infrequent up until the early 1970s. The material contains V, Nb and N, in addition to Mo, V and Cr. Fine precipitates of vanadium carbide and niobium carbonitride in this type of steel suppress the recovery of microstructure during creep and develop high creep strength. The material also exhibits better toughness and higher creep and creep rupture strength than CrMoV steels. Table 5.4 summarizes the typical chemical composition of 12Cr steels for rotor forgings developed so far.17–23 Since the application of 12CrMoVNbN steel by the General Electric
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Creep-resistant steels 1020°C 720°C
FC
FC
Normalizing
Tempering (a)
950°C 670°C Forced air cool
FC
Quenching
Tempering (b) 640°C FC
(c)
5.16 An example of heat treatment for CrMoV rotor shaft forging: (a) Preliminary heat treatment; (b) quality heat treatment; (c) stress relieving heat treatment.
Table 5.3 Example of mechanical properties of EPRI advanced rotor forging project
Quenching Tempering
Low S forging
VCD forging
ESR forging
955°C × 25 h Forced air cool 680°C × 35 h
950°C × 23 h Forced air cool 670°C × 52 h
850°C × 28.5 h Forced air cool 670°C × 35 h
Surface
Centre
Surface
Centre
Surface
Centre
0.2% yield strength (MPa)
619
627
632
635
666
663
Tensile strength (MPa)
775
796
779
787
817
813
Elongation (%)
20.5
20.1
23
22.2
19.1
20.5
Reduction in area (%)
59.3
60.3
68.6
67.3
60.7
60.3
FATT (°C)
73
94
46
67
86
97
Uppershelf energy (J)
157
134
160
136
125
110
vE24°C (J)
14
11
49
18
22
16
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400 Forced air cool Oil quench
350 300
Stress (MPa)
250
200
150 CrMoV scatterband
100 17.0
17.5
18.0 18.5 19.0 19.5 20.0 20.5 Larson Miller parameter, T (20 + log t) 10–3 (K, h)
21.0
5.17 Example of creep rupture strength of CrMoV rotors.
Company, experience of the application of 12Cr rotor forging for high temperature, rotor has increased. Furthermore, with the development of advanced power plants such as ultra super critical (USC) plants, improved type 12Cr steels with higher creep and creep rupture strength such as TOS107, HR1100, TMK1, COST E and COST B rotors have been developed. Trial rotor forgings were evaluated in the EPDC project and COST501 in Europe. Improvement in the creep strength to withstand an increased steam temperature of around 593°C, and pressure condition was made by adding or increasing the contents of alloying elements such as W and Mo while reducing the C content. This enhances the solid solution strengthing and stabilization of carbides. The amount of W and Mo are generally controlled by referencing the Mo equivalent, Mo + W/2 (wt%), of 1.5. An increase in the Mo content has a beneficial effect on toughness whereas an increase in W is effective in developing better creep strength. A rotation test was performed on these advanced 12Cr steels and no problems were disclosed for operation at 593°C in the advanced type 12Cr steels that have been developed. However, a ferritic superalloy A286 forging tested at 650°C disclosed difficulty in application owing to progressive thermal fatigue damage caused by a large coefficient of thermal expansion and a small thermal conductivity.24 In order to realize a 12Cr steel that is serviceable up to 650°C, advanced 12Cr steels have been further improved as a substitute for superalloy A286. In these new 12CrMoV steels such as TOS110, MTR10A and HR1200, W content is increased to enhance solid solution strengthening; Co and B are also added. Addition of Co is effective in suppressing the
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Material
C
Si
Mn
12CrMoVNbN 12CrMoVTaN TOS107 HR1100 TMK1 TMK2 COST B COST E TOS110 MTR10A HR1200
0.19 0.18 0.14 0.13 0.14 0.12 0.17 0.12 0.11 0.12 0.10
0.30 0.27 0.03 0.28 0.07 0.06 0.07 0.10 0.08 0.05 0.06
0.65 0.62 0.52 0.58 0.51 0.49 0.06 0.45 0.1 0.05 0.46
P
S
0.016 0.009
0.017 0.0024
0.008
0.001
0.007 0.008
0.001 0.002
Ni
Cr
Mo
W
V
Ta
Nb
N
Co
B
Ref
0.6 0.30 0.73 0.58 0.60 0.52 0.12 0.74 0.2 <0.05 0.25
10.5 10.3 10.36 10.23 10.28 10.38 9.34 10.39 10.0 10.2 10.2
1.0 0.94 1.05 1.13 1.46 0.28 1.58 1.06 0.65 0.65 0.14
– – 1.06 0.23 – 1.98 – 0.81 1.8 1.75 2.51
0.20 0.25 0.21 0.22 0.17 0.19 0.27 0.18 0.2 0.2 0.21
– 0.089 – – – – – – – – –
0.085 – 0.07 0.06 0.056 0.047 0.059 0.045 0.05 0.06 0.07
0.060 0.0412 0.0414 0.045 0.046 0.051 0.015 0.052 0.02 0.02 0.017
– – – – – – – – 3.0 3.3 2.44
– – – – – – 0.0080 0.0002 0.01 0.002 0.013
17 18 19 20 21 21 17 22 23
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Table 5.4 Typical chemistry of developed 12Cr rotor steels in wt%
Production of creep-resistant steels for turbines
199
formation of delta ferrite during solidification of the ingot by reducing the Cr-equivalent without reducing the creep strength. Addition of B significantly contributes to stabilizing carbides and suppressing the progress of recovery.25 In Europe, COST522 was performed after COST50126 and the COST536 project is running with the aim of developing improved creep strength 12Cr steels for turbine rotor forging operating up to 650°C.27 Figure 5.18 shows an example of the production sequence for a 12Cr rotor forging. The production process for 12Cr steel rotor forging is almost the same as that for CrMoV steel with regard to machining after quality heat treatment, but differs for the journal overlay welding which is followed by stress relieving heat treatment. In order to develop desirable properties for 12Cr steels, production of a sound ingot is particularly important. In casting large 12Cr steel ingots, possible problems are the formation of delta ferrite and the macro-segregation in the ingot. Formation of delta ferrite basically depends on the shape of the
Steelmaking and ingot making
Forging
Preliminary heat treatment
Rough machining UT Quenching
Tempering
Machining NDE, Mechanical and metallurgical test Journal overlay
Stress relieving
Machining of journal NDE for Journal, HIT Shipping
5.18 Production sequence for 12Cr steel rotor shaft.
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Creep-resistant steels
gamma-loop in the phase diagram of the material. Alloy modification by adding or increasing alloying elements such as Cr, Mo, W for better creep strength tends to increase the potential for formation of delta ferrite. The existence of delta ferrite reduces the strength, ductility and toughness of the material. Since delta ferrite is hardly eliminated by heating and hot working during the forging process, formation of delta ferrite during solidification of an ingot must be avoided. The Cr equivalent value is mostly used as a measure of the susceptibility to form delta ferrite. The value depends on the chemistry of the material and is given as follows: Cr-eq = Cr + 6Si + 4Mo + 1.5W + 11V + 5Nb – 40C – 30N – 4Ni – 2Mn – 2Co (%) It is generally said that delta ferrite will not present at a Cr-eq of less than 10. The potential for delta ferrite formation, however, strongly depends not only on the chemistry but also solidification rate which is significantly affected by the size of the ingot. Therefore, the Cr-eq value should be controlled based on experience of particular compositions and casting processes. The major aim of Co addition in new 12Cr steels is to decrease the Cr equivalent to avoid the formation of delta ferrite. The formation of segregation and precipitation in eutectic carbonitrides like NbCN is another problem. Enrichment of alloying elements in the segregation zone promotes formation of eutectic carbonitrides and causes significant deterioration in the toughness and ductility of the material. In order to avoid the formation of NbCN, the content of C and Nb need to be controlled. As shown in Fig. 5.18, the raw materials are melted and subjected to oxidizing refining in an EAF. Then reduction refining and degassing is performed in LRF after which the material is poured into an ingot mould during a VCD process. An ESR process can be also applied as an alternative method of refining and ingot making and several turbine builders have asked for 12CrMoV steel ingots to be made by ESR in order to eliminate problems related to segregation. In addition, elctroslag hot topping is also applied in the manufacture of 12Cr steel ingots. New 12Cr rotors with alloying elements like Mo, W, Co and B are being produced by the ESR process to avoid segregation and to develop homogeneous material properties. The forging process for 12Cr steels is basically the same as that for CrMoV steels, as shown in Fig. 5.10. In case of the 12Cr steels, the resistance to deformation in hot working is larger than that of low alloy steels and forging of 12Cr steel requires a larger load and sometimes needs additional steps to finishing compared with CrMoV steels. Forging at high temperature is preferable from the standpoint of hot workability. However, increasing the heating temperature encourages formation of delta-ferrite and grain growth. In addition, careful attention is needed to avoid loss of ductility caused by the formation of low melting boride for boron bearing 12Cr steels.23
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An example of a heat treatment diagram for conventional 12Cr steel is shown in Fig. 5.19. After the forging process, normalizing heat treatment followed by tempering is performed before quality heat treatment. The pearlite transformation process can be applied as an alternative preliminary heat treatment. The major purpose of these preliminary heat treatments is grain refining to provide good toughness and enhanced detectability indicated by NDE. It is essential to attain a rapid cooling rate in the quenching process of 12Cr steels, to avoid precipitation of carbides in the austenite phase and to complete martensitic transformation. Since the material has good hardenability, cooling is performed by oil quenching or water spray quenching. The austenitizing temperature for quenching is set at the resolving temperature of Nb carbides. In 12Cr steels, some of the retained austenite may exist in the material after quenching. In order to eliminate retained austenite, double tempering heat treatment is commonly performed after quenching. The first tempering is performed at around 600°C for complete transformation of the retained austenite. The second tempering is to develop the target mechanical properties. After quality heat treatment, stress relieving heat treatment is usually conducted by heating at more than 30°C below the second tempering temperature. Overlay welding on the journal of rotor forging is a peculiar process in the manufacture of 12Cr steel forging. Since the thermal conductivity of 12Cr steel is small, seizure is likely to occur in the bearing area.28 In order to 1100°C
700°C FC
FC
Normalizing
Tempering (a)
1050°C Oil quench Quenching
660°C
570°C FC 1st Tempering (b)
FC 2nd Tempering
630°C FC (c)
5.19 An example of heat treatment for 12Cr steel rotor shaft forging. (a) Preliminary, (b) quality and (c) stress relieving heat treatments.
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Creep-resistant steels
avoid seizure, a CrMo steel sleeve, platings and overlay weldings have been applied to the journal of the 12Cr rotor forging. The reliability of the overlay is superior to other steel sleeve and platings and is currently applied to the 12Cr rotor steel forging. Overlay welding is often applied to CrMo type materials. Tensile residual stress occurs owing to the difference in thermal expansion coefficients between the forging material and the overlay material. On the standpoint of fatigue strength, the overlay surface is sometimes rolled to attain compressive residual stress in the bearing area. Figure 5.20 shows a journal overlay just after welding. The applied welding process and overlay material are selected according to the properties of the forging material. Since 12Cr steels develops good hardenability, the microstructure of the 12Cr rotor forging shows a tempered martensite structure from the centre to the surface of the forging. Table 5.5 shows the typical mechanical properties of developed 12Cr steels. The balance of strength versus toughness is much better than that of CrMoV steels. Generally, alloying to increase the creep strength increases the martensitic transformation temperature and reduces the toughness. Figure 5.21 shows the creep strength of conventional, advanced and new 12Cr steels. The 105 h rupture temperature under a stress of 100 MPa is around 580°C for conventional 12CrMoVNbN steels and around 600°C for advanced type steels. Although the new 12Cr steels are aimed at application at 650°C, the long term properties tend to deteriorate, so that presently 630°C is assumed to be the optimum. Efforts have been made to characterize premature fracture and to develop a material serviceable at 650°C.
Overlay Overlay
5.20 A 12CrMoV rotor forging with overlay welding on the journal.
WPNL2204
Material
12CrMoVNbN 12CrMoVTaN TOS107 HR1100 TMK1 TMK2 COST B COST E TOS110 MTR10A
Body diameter (mmϕ)
1262 1200 1200 1220 840 1150 1296 1200
Austenitizing (°C)
Tempering (°C)
1050 1050 1050 1050 1090 1050 1100 1070 1070
570 570 570 665 550 550 590 570 570
+ 620 + 640 + 660 + + + + +
665 680 700 690 690
0.2%YS (MPa)
TS (MPa)
El (%)
RA (%)
Centre FATT (°C)
Centre vE (J)
(0.02%YS (0.02%YS (0.02%YS 765 (0.02%YS 642 801 (0.02%YS 730
935 920 918 900 883 801 914
16.7 20.9 20.8 20 19.0
43.1 55.7 52.6 60 59.0
61 47 43 20
24 35 35 70
60 5
33 86
830
20
= 709) = 700) = 695) = 759)
= 615)
YS, yield strength; TS, tensile strength; EL, elongation; RA, reduction in area.
65
62
Production of creep-resistant steels for turbines
Table 5.5 Typical mechanical properties of the 12Cr rotor steels
203
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Creep-resistant steels
100 000 h rupture strength (MPa)
400
950
Temperature (F) 1050 1100
1000
1150
1200
300 Advanced 12Cr steel 200
New 12Cr steel
12Cr steel
CrMoV steel 100 90 80 70 60 50 500
550
600 Temperature (°C)
650
5.21 Representative creep rupture strengths of CrMoV and 12Cr steels.
5.3.3
High pressure–low pressure combination rotor
A combined cycle power plant, which acomplishes high efficiency, consists of gas turbines and steam turbines. A high pressure–low pressure combination rotor (HLP) has been often used in combined cycle power plants since the combination type turbine has advantages, such as smaller building space, lower cost, easy maintenance, and so on, compared with the separate type steam turbines. Figure 5.22 shows a schematic of HLP turbine rotor forging and its required properties. Generally, both the high and low temperature sections of the HLP turbine are made from the same material. Optimization of chemistry of the material and manufacturing process is important in developing target creep strength in the HP section and good centre toughness in the LP section simultaneously. With the increase in the capacity of the combined cycle power plant, a larger HLP turbine has been required. Extensive research on the development of material has been carried out.29–34 The material needs to have good hardenability to develop a low temperature transformation microstructure, which leads to good toughness, even at the centre of a large diameter LP section. A creep strength equivalent to that of CrMoV steel is also required in the HP section. Table 5.6 summarizes the typical chemistry of the low alloy materials developed for large HLP turbine rotor forging. The chemistry of these materials is almost midway between the 3.5NiCrMoV steel and CrMoV steel with a small addition of the other elements. In order to attain the good toughness at the centre of the LP section, the Cr and Ni content are optimized taking account of the balance with creep strength. Addition or increase in the content of elements such as
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Production of creep-resistant steels for turbines
205
LP turbine HIP turbine High strength High toughness
Creep strength
LP and HIP turbine (separate type)
LP portion HP portion High strength High toughness
Creep strength
HLP turbine (single cylinder type)
5.22 Requirements for materials in a HLP turbine.
Nb, W is effective in increasing the creep strength. The chemistry of the materials was determined taking account of the optimum balance of toughness and creep strength which is achieved by the application of the special heat treatment process. The steelmaking process for materials for HLP turbine rotor forging is almost the same as those applied to CrMoV steels. Conventional refining process by EAF and LRF and subsequent casting into moulds can be applied. The ESR process is also used to avoid the segregation and increase the homogeneity of the ingots. The forging process for an HLP turbine is similar to those for HP/IP turbine rotor shaft forging. After the upsetting, cogging and finish forging are performed. Normalizing and tempering or pearlite transformation process can be applied as the preliminary heat treatment. Differential quality heat treatment is a peculiar process in developing the creep strength in the HP portion and high centre toughness in the LP portion simultaneously. Figure 5.23 shows a schematic of special equipment for differential heat treatment. Figure 5.24 shows an example of a differential heat treatment diagram. The HP section is heated to a higher temperature and subjected to forced air cooling. On the other hand, the LP section is heated at a comparatively low austenitizing temperature to suppress grain growth and develop the fine grain microstructure. Then the LP portion is water spray cooled to form a low temperature transformation microstructure such as lower bainite, which develops good toughness in all the LP portion. Differential heat treatment is performed in a vertical furnace which consists of two separate temperature sections under rotation of the forging. In another differential quenching technology, control of the cooling rate has been developed to simulate that of oil quenching by adjusting the intensity of air and water spray.29 Subsequently, each section is tempered, usually differentially, to reach the target strength level.
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EPRI-Europe
C
Si
Mn
P
S
Ni
Cr
Mo
W
V
Nb
Ref.
0.22
0.07
0.67
0.006
0.002
0.74
2.10
0.84
0.66
0.30
–
29
0.22
0.03
0.02
0.003
0.001
2.49
1.58
1.19
–
0.23
–
30
0.24
0.02
0.45
0.004
0.0009
1.69
2.22
1.08
0.19
0.19
0.015
31
2CrMoNiWV EPRI-Japan 2.5NiCrMoV 2.25CrNiMoVWNb 2Cr1.8NiMoV
0.23
0.01
0.20
0.004
0.002
1.74
2.03
1.17
–
0.26
–
32
2CrMo1.5NiV
0.24
0.02
0.20
0.004
0.004
1.48
1.99
1.69
–
0.23
–
33
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Creep-resistant steels
Table 5.6 Chemical composition of the HLP rotor steels, in wt%
Production of creep-resistant steels for turbines
207
Electric furnace Austenitizing temp. for HP portion
Fan
HP
HP
LP
LP Water spray
Austenitizing temp. for LP portion Differential heating
Differential cooling
5.23 Schematic of differential heat treatment.
Owing to the differential heat treatment, a transition area of mechanical and impact properties unavoidably exists in the forging although the forging has a bainitic microstructure. Several trail rotors of various chemistries, shown in Table 5.6, were investigated in detail and it was confirmed that a good balance between strength and toughness is attained at the centre of the low pressure portion, sufficient to cope with a larger body diameter and a creep rupture strength exceeding that of CrMoV steel was attained in the high pressure portion.29–34 Table 5.7 summarizes the mechanical and impact properties of the HLP rotor forgings developed. Figure 5.25 shows the relationship between strength and 50% fracture appearance transition temperature (FATT) at the centre of the low pressure portion in the developed low alloy steel rotor forging. The balance between strength versus toughness is remarkably improved in these advanced steels compared with that of CrMoV steel. Figure 5.26 also shows an example of creep rupture strength.30 Including these results, the HP section of all the trial rotor forging listed in Table 5.6 is confirmed to have the equivalent creep rupture strength to that of CrMoV steel. Figure 5.27 shows an HLP rotor forging.
5.4
Future trends
The production technology for rotor shaft forgings made of low alloy steels and 12Cr steels has been established and high quality forgings are being
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Creep-resistant steels 1100°C 680°C FC Pearlite transformation (a) 900°C
(∗) Quenching
625°C
580°C
Water spray
LP
FC
FC (∗) 2nd tempering
1st tempering
970°C Forced air cool
HP
640°C
580°C FC
(∗) Quenching
FC (∗)
1st tempering
2nd tempering
(b) 630°C FC (c)
5.24 An example of a heat treatment process for HLP turbine rotor forging. (a) Preliminary, (b) quality and (c) stress, relieving heat treatments. (*) represents differential heat treatment.
manufactured and operated in power plants. Many forgemasters have been equipped with advanced and large capacity facilities for the production of heat-resistant steel forgings. With regard to ferritic heat-resisitant steels, as already mentioned, new 12Cr steel developed for application up to 650°C still presents the problem of premature fracture at around 650°C. Therefore, efforts to eliminate the premature fracture and manufacture ferritic steel at 650°C have been continuing.27 The importance of higher efficiency fossil power plant is still increasing as a solution for the problem of global warming. Several projects are attempting to realise advanced ultra super critical (A-USC) power plants using steam temperatures over 700°C, which is around 100°C higher than the current typical USC power plants. The THERMIE project was first commenced in
WPNL2204
Table 5.7 Mechanical properties of the HLP rotor forgings Body diameter (mmϕ)
Quenching (°C)
Tempering (°C)
0.2%YS (MPa)
TS EL (MPa) (%)
RA (%)
Centre FATT (°C)
Centre vE at RT (J)
EPRI/Japan 2.5MoCrMoV
LP HP
1750 1720
935WSQ 950FAN
650 650
651 653
805 972
22 22
68 67
22 3
165 157
EPRI/Europe 2CrMoNiWV
LP HP
1750 1250
954WSQ 954WSQ/Air
655 655
716 726
925 840
20 19
70 66
57 55
52 41
2.25CrNiMoVWNb
LP HP
1750 1000
900WAQ 970FAN
625 640
682* 655*
839 781
22 23
61 63
23 40
71 46
2Cr1.8NiMoV
LP HP
1720 1320
950WSQ 970FAN
666*
864
21
59
46
48
2CrMo1.5NiV
LP HP
1356 968
910WQ 940FAN
620
620
18
60
–5
72
650 658
*0.02%YS; Data is not available for the blank rows.
Production of creep-resistant steels for turbines
Material
209
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CrMoV 120
LP centre FATT (°C)
Advanced CrMoV
80 2Cr1.8NiMoV 2.5NiCrMoV
40
2.25CrNiMoVWNb 2CrMoNiWV 2CrMoNiWV 2.5NiCrMoV 0 2CrMoNiWV
–40
700
800 900 1000 Tensile strength (MPa)
1100
5.25 Balance between toughness of LP centre and tensile strength. 700 550°C
600 500
600°C 630°C
Tangential Centre material of HP portion
Longitudinal
400 300 Stress (MPa)
210
200
Mean creep curve of conventional CrMoV steel
100
538°C, 105 h 50 16
17 18 19 20 21 22 Larson Miller parameter, T (20 + log t ) × 10–3
23
5.26 An example of creep rupture strength of material for HLP rotor.
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211
5.27 An HLP rotor forging.
1998 to develop such a A-USC plant and was superceded by the AD700 and COMTES projects.35,36 In Japan and USA, activities to realize A-USC power plant are also continuing. One of the most critical points for enabling the provision of A-USC plants will be the material used for high temperature application over 700°C. For high temperature turbines, several candidate materials, Alloy 617, Alloy 625, Alloy 263 and Alloy 718 were selected for forgings in the THERMIE project. Large super alloy ingots of Alloy 718 and Alloy 706 have been made by ESR and/or VAR and much experience of production exists for landbased gas turbine disks. However the weight of ingots produced so far is at most 20 ton. Formation of segregation depends significantly on the chemistry of the alloys and the ingot diameter. Production of a sound ingot from super alloys become more difficult with an increase in the ingot diameter. Production was demonstrated to be feasible for the application of superalloys to turbine rotor forging, and a 1-m diameter ESR ingot was produced.36 Weld type construction type is proposed in the THERMIE project for the design of a high temperature rotor shaft in a A-USC plant.35,36 The turbine is constructed by welding a superalloy and 12Cr steel. This makes it possible to construct a large turbine using smaller pieces of superalloy forging(s) Many high efficiency plants such as A-USC, combined cycle, IGCC, IGFC, and so on, have been designed. Current production technology of heat-resistant steel turbine forgings supports the realization of these power
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plants and the development of advanced production technologies undoubtedly contributes to a further increase of plant efficiency.
5.5
References
1 Tanaka Y and Ishiguro T, ‘Development of high-purity large-scale forgings for energy service’, Phisica Stati Solidi(a), 1996, 160 305–320. 2 Sawada S and Kawaguchi S, ‘The beneficial effect of vacuum carbon deoxidation on rotor forging properties’, Workshop on Rotor Forgings for Turbines and Generators, Palo Alto, EPRI, 1980, 5–1. 3 Ikeda Y, Yoshida H, Tanaka Y and Fukuda T, ‘Production and properties of superclean monoblock LP turbine rotor forging’, in Clean Steel: Superclean Steel, Nutting J and Viswanathan R (eds), The Institute of Metals, 1996, 71–87. 4 Mayer W, Bauer R and Zeiler G, Development of production technology and manufacturing experiences with superclean 3.5NiCrMoV steels’, in Clean Steel: Superclean Steel, Nutting J and Viswanathan R (eds), The Institute of Metals, 1996, 89–100. 5 Jaffee R I, ‘Development of superclean rotor steels’, in Superclean Steels, Jaffee R I (ed.), Pergamon Press, 1991, 3–27. 6 Fielder H, Richter G and Scharf G, ‘Application of special metallurgical processes for the production of highly stressed forgings’, The 8th International Forgemasters Meeting, Paper No. 15, Kyoto, Japan, 1977. 7 Morinaka K, Futamura Y, Kitagawa I and Watanabe S, ‘The manufacture of the large ESHT-J ingot’, I&SM, 1989, April, 9–15. 8 Gould G C, Long Time Isothermal Embrittlement in 3.5Ni, 1.75Cr, 0.50Mo, 0.20C Steel, ASTM STP407, ASTM, 1968, 90–105. 9 Tanaka Y, Azuma T, Yaegashi N and Ikeda Y, ‘10000H isothermal ageing test results of NiCrMoV Steels for low pressure steam turbines’, in Clean Steel: Superclean Steel, Nutting J and Viswanathan R (eds), The Institute of Metals, 1996, 71–87. 10 Takenouchi T, Ikeda Y and Tanaka T, ‘Production of 12CR rotor forgings for steam turbines using advaced VCD process’, Recent Developments in Rotor Forging Steels, Iron & Steel Society, Warrendale, PA, 1990. 11 Kolar M, ‘Discussion on Heat Treatment Practices’, Workshop on Rotor Forgings for Turbines and Generators, Palo Alto, EPRI, 1980 5–105. 12 E-472/472M : Standard Test Method for Heat Stability of Steam Turbine Shaft Forgings, in ASTM Standard, ASTM 01.05, 2006. 13 Swaminathan V R, Steiner J E and Mitchel A, Advanced RotorForgings for HighTemperature Steam Turbine-Vol 1, EPRI Report CS-4516, Palo Alto, EPRI, 1986. 14 Swaminathan V R, Steiner J E and Mitchel A, Advanced Rotor Forgings for HighTemperature Steam Turbine – Vol 2, EPRI Report, CS-4516, Palo Alto, EPRI, 1986. 15 Finkler H and Potthast E, New 2Cr-Mo-Ni-V Steel for High-Pressure Rotors, ASTM STP 903, ASTM, 1986, 107–123. 16 Potthast R, Viswanathan R and Wieman W, ‘Advanced 2%CrMoNiWV Steel for Combination Rotors, Proceedings of 12th International Forgemasters Meeting, Chicago, Forging Industry Education and Research Foundation, Cleveland, OH, 1994. 17 Tsuda Y, Yamada M, Ishii R, Tanaka Y, Azuma T and Ikeda Y, ‘Development of high strength 12% Cr Ferritic steel for turbine rotor operated above 600°C’, Proceedings of 13th International Forgemasters Meeting, Pusan, Korea, October 1997, 417–428.
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18 Iijima K, Siga M, Yoshioka T, Fukui Y, Kaneko R, Simomura K and Sasaki R, ‘Steam turbine materials for improved coal fired plants’, The First International Conference on Improved Coal-Fired Power Plants, Palo Alto, CA, USA, EPRI, 1986. 19 Tsuchiyama T, Suzuki K, Kohno M, Arihara H, Okamura M and Ohizumi H, ‘Manufacturing and quality of large ESR 12%Cr rotor forging’ , Proceedings of 11th International Forgemasters Meeting, Terni, Italy, 1991 IX-4, 1–10. 20 Hizume A, Takeda Y, Yokota H, Takano Y, Suzuki A, Kinoshita S, Kohno M and Tushichiyama T, ‘The probability of a new 12% Cr rotor steel applicable for steam temperatures above 593°C’, Proceedings of International Conference on Advances in Material Technology for Fossil Power Plants, Chicago, ASM International, 1987, 143. 21 Berger C, Beech S M, Mayer K H, Scarlin R B and Thorton D V, ‘High temperature rotor forgings of high strength 10%CrMoV steel’, Proceedings of 12th International Forgemasters Meeting, Chicago, 1994, Section 8 – 1, 1–17. 22 Kagawa Y, Tamura F, Ishiyama O, Matsumoto O, Honjo T, Tsuchiyama T, Manabe Y, Kadoya Y, Magoshi R and Kawai H, ‘Development and manufacturing of the next generation of advanced 12Cr steel rotor for 630C steam temperatures’ , Proceedings of 14th International Forgemasters Meeting, Wiesbaden, Germany, 2000, 301– 308. 23 Arai M, Doi H, Fukui Y, Kaneko R, Azuma T and Fujita T, ‘Improvement of long time creep rupture properties of High WcoB containing 12Cr rotor steels for use of 650°C in USC power plant’, Proceedings of the 3rd Conference on Advances in Material Technology for Fossil Power Plants, Wales, UK, 2001, 415–423. 24 Muramatsu K, ‘Development of ultra-super critical plant in Japan’, in Advanced Heat Resistant Steel, Viswanathan R and Nutting J (eds), IOM Communications, London, UK, 349–364. 25 Azuma T, Kazuhiro M and Tanaka Y, ‘Effect of boron on creep strengthening in 12% Cr heat resistant steel’, 14th International Fogemasters Meeting, Wiesbaden, Germany, September 2000, 283–289. 26 Thornton D V and Mayer K H, ‘European high temperature materials development for advanced steam turbines’, in Advanced Heat Resistant Steel, Viswanathan R and Nutting J (eds), IOM Communications, London, UK, 349–364. 27 Scarlin B, Kern T-U and Staubli M, ‘The European efforts in material development for 650C USC Power Plants-COST 522’, in Advances in Material Technologies for Fossil Power Plant, Viswanathan R, Gandy D and Coleman K (eds), ASM International, 2004, 80–99. 28 Haas H, Simmermann A and Termuehlen H, ‘Turbines for advanced steam conditions – operational experience and development’ The First International Conference on Improved Coal-Fired Power Plants, Palo Alto, USA, EPRI, 1986. 29 Potthast E, Poppenhager J, Wiemann W and Mayer K H, ‘Advanced 2%CrMoWV steels for combination rotors’, Proceedings of 11th International Forgemasters Meeting, Terni, Italy, Societa Delle Fucine, 1991, IX–8. 30 Tanaka Y, Ikeda Y, Ohnishi K, Kawaguchi S, Watanabe O, Kaplan A, Schwant R C, Jaffee R I and Poe G, ‘Development of a superclean 2.5NiCrMoV rotor steel for HP and LP application’, Proceedings of 11th International Forgemasters Meeting, Terni, Italy, Societa Delle Fucine, 1991, IX–7. 31 Yamada M, Tsuda Y, Watanabe O, Miyazaki M, Tabaja Y, Takenouchi T and Ikeda Y, ‘HLP single cylinder steam turbine rotor forgings for combined cycle power
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32
33
34
35
36
Creep-resistant steels plants’, Proceedings of the Robert I. Jaffee Memorial Symposium on Clean Materials Technology, Chicago, USA, ASM International, 1992, 161. Fukui Y, Shiga M, Hidaka K, Kaneko R and Tan T, ‘Development of superclean 0.2Mn-1.8Ni-Cr-Mo-V steel rotor for Hp-Lp turbine’, Proceedings of the Robert I. Jaffee Memorial Symposium on Clean Materials Technology, ASM Materials Week, Chicago, USA, ASM International, 1992, 249. Tsuchiyama T, Miyakawa M, Okamura M, Matsumura K, Morita M, Yamamoto T and Nishida S, ‘Development and production of a new HP-LP combined turbine rotor of 2CrMoV steel’, Proceedings of the Robert I. Jaffee Memorial Symposium on Clean Materials Technology, ASM Materials Week, Chicago, USA, ASM International, 1992, 181. Kitagawa, K. Soeda, Tsuji I and Kadoya Y, ‘Manufacturing of 2 1/4CrMoV steel HP-LP merged type steam turbine rotor forgings’, Proceedings of 11th International Forgemasters Meeting, Terni, Italy, Societa Delle Fucine, 1991, IX–10. Kern T-U, Wieghardt K and Kirchner H, ‘Material and design solutions for advanced steam power plants’, in Advances in Material Technologies for Fossil Power Plant, Viswanathan R, Gandy D and Coleman K (eds), ASM International, 2004, 20–34. Scarlin B, ‘Material developments for ultrasupercritical steam turbines’, in Advances in Material Technologies for Fossil Power Plant, Viswanathan R, Gandy D and Coleman K (eds), ASM International, 2004, 51–67.
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6 Physical and elastic behaviour of creep-resistant steels Y. F. Y I N and R. G. F A U L K N E R, Loughborough University, UK
6.1
Introduction
The general physical and elastic properties of creep-resistant steels are sometimes the first to be considered in design calculations. For example, ferritic/martensitic steels are sometimes more favoured than austenitic steels in power plant owing to their merit of lower thermal expansion and higher thermal conductivity. The former provides more structural stability and the latter reduces temperature gradient within a component and therefore yields lower thermal stress levels and greater heat transfer rates when the temperature of a component changes during heat up and cool down. The historical problem with superheaters, reheaters and associated components has been that of thermal fatigue, in which problems are exacerbated by weak, heavy section components, complex geometries and bending stresses (Starr, 2002). A low coefficient of expansion and low elastic modulus is obviously advantageous, because the thermal stress, σthermal caused by a temperature change of ∆T is directly proportional to the product of the coefficient of linear thermal expansion, α and Young’s modulus, E, i.e. σthermal = α ∆TE
[6.1]
Fatigue stresses can result from pipework movement in the plant, during heat up and cool down when load changes occur (secondary stresses). Here, the advantage is with strong thin wall members having innate flexibility and whose deadweight does not overwhelm pipe support systems. However, during start up, rapid changes in temperature in the plant can lead to significant through wall temperature differences. This situation occurs more frequently nowadays as two-shifting operations of power plants are more common. In this case, as well as good high strength properties, a high thermal conductivity is also of advantage. An estimate of likely susceptibility to thermal fatigue of steel is given by the thermal stress parameter (TSP), which is defined as: TSP = αE k
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where α is the coefficient of linear thermal expansion with units of 10–6K–1, E is Young’s modulus in GPa and k is the thermal conductivity in Wm–1K–1. The lower the thermal stress parameter, the more resistant is the alloy to thermal stress. Thus, low thermal expansion, low elastic modulus and high thermal conductivity are desirable. More recent studies show the susceptibility to thermal fatigue is also related to the yield strength of the material at a particular temperature and a modified parameter, R-value, has been proposed to describe the resistance-to-crack of a material (Skelton and Beckett, 1987):
R=
SY 0.2 k αE
[6.3]
where SY0.2 is the 0.2% yield strength. Opposite to TSP, higher R-values indicate lower probability for thermal induced crack initiation. R-values for some common steels and a Ni-base alloy at 650°C (except those specified) are shown in Fig. 6.1. It is clear from Fig. 6.1 that ferritic/martensitic steels and Ni-base alloys are much more resistant to thermal induced crack initiation. In this chapter, the general physical behaviour, mainly the thermal properties of creep resistant steels will be discussed. Where available, data concerning the measured properties will be given. The implications of these properties for industrial applications of creep-resistant steels are also discussed.
7000 Austenites 6000
Ni-base alloy Martensites New 1Cr NiCr20 10Cr (550C) TiA1 (600C)
R-value
5000
4000 A286
3000 Alloy 2000 800 Esshete Type 1250 347 Type Type 1000 304 316 Type 321
X8CrNi MoBNb 16–16 NF 709 X8CrNi X8CrNi MoNb MoVNb 16–16 16–13
0
6.1 R-values (resistance-to-crack) of some ferritic/martensitic, austenitic and Ni-base alloys at 650°C (except those individually labelled in the figure).
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6.2
219
Elastic behaviour
The response of any material to externally applied forces is deformation, that is a change of their size, shape and/or volume. These changes in dimension and volume will be reversed, remain permanent or be a combination of the two when the externally applied load is withdrawn. When the changes are reversed, this is called elastic deformation and when the changes are permanent, this is called plastic deformation. This chapter only deals with elastic behaviour of creep-resistant steels. Plastic deformation will be discussed in later chapters. To understand the behaviour of a material under load, it is necessary to define the terms stress and strain.
6.2.1
Stress and strain
When a pair of balanced forces, F, acting on the opposite side of a bulk material with cross-sectional area, A, the force can be resolved to two components, one perpendicular to the surface, Fn, and the other within the plane of the surface, Fτ, as illustrated in Fig. 6.2. The action of such a pair of applied forces can be represented by two stresses, the tensile stress, σ, and the shear stress, τ. They are defined as follows: σ=
Fn A
[6.4]
Fn
F
Fτ
A
Fτ
F
Fn
6.2 Tensile and shear stresses caused by a pair of externally applied forces.
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Creep-resistant steels
τ=
Fτ A
[6.5]
The metric unit of stress is N m–2. In engineering applications, this unit is too small and is often replaced by MPa (1 MPa = 106 N m–2). The basic response of a block of material to applied load is its dimensional, and/or volumetric and/or shape change depending on the nature of the applied load. In the simplest case, the applied stress is a pure tensile stress (Fig. 6.3(a)) and the block would elongate from the original length, l0, to a new length, l0 + u. This dimensional change is described by the nominal tensile strain, εn, which is defined as:
εn = u l0
[6.6]
In general, the block would shrink sidewise when it is stretched. This can be represented by a nominal lateral strain:
εl = – w w0
[6.7]
These two strains are related through Poisson’s ratio: ν=–
εl εn
[6.8]
A shear stress will cause a shear strain. The engineering shear strain is defined as (see Fig. 6.3(b)): σ
u /2
w
τ
w p
w0
l0
u /2
l
τ
p
V0
V0 – ∆V
p
σ
w /2
w /2 (a)
(b)
(c)
6.3 Illustration of various forms of deformation and the corresponding strains. (a) Tensile, (b) shear, (c) dilatation.
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p
Physical and elastic behaviour of creep-resistant steels
γ= w l
221
[6.9]
The strain caused by hydrostatic pressure is called dilatation and is defined as (see Fig. 6.3(c)): ∆ = ∆V V0
6.2.2
[6.10]
Modulus of elasticity
At small strains, the relation between strain and the applied stress obeys Hooke’s Law, that is the deformation of the material is linear–elastic. Therefore, in the case of simple tension, the nominal tensile strain εn is proportional to the tensile stress: σ = Eεn
[6.11]
where E is the Young’s modulus. Similarly, the following relations hold τ = Gγ
[6.12]
p = – K∆
[6.13]
where G and K are called the shear modulus and the bulk modulus, respectively. The four elastic constants, namely E, G, K and ν, are related to each other by the following equations (Cottrell, 1995): K=
E 2(1 – 2 ν )
[6.14]
G=
E 2(1 + ν )
[6.15]
E=
9 KG 3K + G
[6.16]
It is sometimes useful to know that the above equations can be approximately rewritten as (Ashby and Jones, 1996): K ≈ E, G ≈ 3/8E and ν ≈ 0.33
[6.17]
The values of Young’s modulus of some common creep-resistant steels (chemical composition as shown in Table 6.1) are listed in Table 6.2, together with values for pure iron for comparison. From Table 6.2, it is clear that the values of Young’s modulus do not change very much from one alloy to another within a specific type of alloy. For example, austenitic steels listed in Table 6.2 have Young’s modulus varying from 193–200 GPa at 500°C, although their compositions differ considerably.
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Steel
C
Si
Mn
Cr
Ferritic/ martensitic
T22 P91 P92 E911 HCM12A 410
0.12 0.1 0.07 0.1 0.11 0.14
0.3 0.4 0.03 0.17 0.1 1
0.45 0.45 0.45 0.47 0.6 1
2.25 9 9 9 12 12.5
Austenitic
A286 304L 304 304H 309 310 316 321 347
0.08 0.03 0.08 0.08 0.2 0.08 0.06 0.08 0.08
1 0.6 0.6 0.6 0.75 0.6 1 0.75 0.75
2 2 2 1.6 2 1.6 2 2 2
13.5–16 18 18 18 22–24 25 17 17–19 17–19
Ni 0 0 0 0 0 0.5 24–27 8 8 8 12–15 20 12 9–13 9–13
Mo
W
V
Nb
B
N
Cu
Ta
Ti
P
Ref.
1 1 0.5 1 0.4 0
0 0 1.8 1 2 0
0 0.2 0.2 0.2 0.2 0
0 0.08 0.05 0.07 0.05 0
0 0 0.004 0 0.003 0
0 0.05 0.06 0.07 0.03 0
0 0 0 0 1 0
0 0 0 0 0 0
0 0 0 0 0 0
0 0 0 0.007 0 0.04
1 1 1 1 1 2
1–1.75 0 0 0 0 0 2.5 0 0
0 0 0 0 0 0 0 0 0
0.1–0.5 0 0 0 0 0 0 0 0
0 0 0 0 0 0 0 0 10 × C–1
0.003–0.01 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0.1 0 0
0 0 0 0 0 0 0 0 0
0 0 0 0 0 0 0 0 (b)
1.9–2.3 0 0 0 0 0 0 (a) 0
0 0.04 0.04 0.04 0.045 0 0.045 0.045 0.045
3 4 5 6 7 5 5 8 9
(a) 5 × (C + N) – 0.7 (b) 10 × C-1 minus Nb 1, Klueh (2005); 2, Tsai et al, (2002); 3, De Cicco et al, (2005); 4, Ravi Kumar et al. (2006); 5, Alyousifa and Nishimura (2006); 6, ZeladaLambri et al. (1999); 7, Ye and Wang (2006); 8, Chênea et al. (2007); 9. Laha et al. (2007).
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Table 6.1 Chemical compositions of some creep-resistant steels (weight percent, Fe balance)
Temperature T (°C)
Thermal expansion coefficient α (10–6 K–1)
Thermal conductivity k (Wm–1K–1)
Young’s modulus E (GPa)
Heat capacity C (J kg–1 K–1)
699
Fe
600
14.5
38.9
196
T22 P91 P92 E911 HCM12A 410 9Cr–0.12C–1Mo
600 600 600 600 550 600 20
14.6 12.6 13.1 ~12 12 11.6 11.15
33 30 29.8 ~27 29.5 24.9 26
167 168 170 ~180 179
18.7 18.7 19.8 17.6 17.1 17.5 17.6 18.5 18.4
21.5 21.5 21.4 18.7 18.7 19.5 23.8 21.4 21.4
193 193 200 200 200 193 163 193 193
Austenitic
Iron
Ferritic/ martensitic
Steel
304L 304 304H 309 310 316 A286 321 347
500 500 500 500 500 Various 538 500 500
770 630
Electrical resistivity ρ(10–6Ωm) at 20°C 10.1
0.992
744 402 500 500 500 502 502 500 461 500 500
57 49.9 72 72 72 78 78 74 72 72
References
Brandes and Brook (1992) Starr (2002 Haarmann et al. (2002) Richardot et al. (2000) Starr (2002) Yoo (2004) Sandmeyersteel (2007) Brandes and Brook (1992) Matweb (2007) Matweb (2007) Hightempmetals Hightempmetals Hightempmetals Matweb (2007) Hightempmetals Matweb (2007) Matweb (2007)
(2007) (2007) (2007) (2007)
Physical and elastic behaviour of creep-resistant steels
Table 6.2 Some physical properties of the steels listed in Table 6.1, together with those for pure iron
223
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Creep-resistant steels
The values of Young’s moduli of austenitic steels are also very close to that of pure iron at 600°C. However, the difference between these values for austenitic steels and the Young’s modulus of ferritic steels is marked; about 10–15%. This indicates that it might be difficult to achieve a desired Young’s modulus value by simple addition of alloying elements. This property is mainly determined by the type of the steel, that is ferritic/martensitic or austenitic. Indeed, from an atomic point of view, Young’s modulus is determined by (1) the bonding between atoms; (2) the number of atoms contributing to holding an atom in place in the structure; and (3) the distance between the atoms. The bonding of atoms in steels is mainly metallic, whether it is ferritic/martensitic (BCC) or austenitic (FCC). However, the FCC austenitic steel is generally more densely packed when compared with the BCC martensitic steel. In another words, the average distance between atoms in austenitic steels is less than that in martensitic steels. Therefore, the attracting forces between atoms which hold the atoms together are stronger in austenitic steels than those in martensitic steels. This means that higher stresses must be applied to cause the same deformation, that is, a higher Young’s modulus. At higher temperatures, the atoms are more active because they have more kinetic energy. Therefore, it is easier to pull them apart and this results in a decrease in Young’s modulus. An example of decreasing Young’s modulus as temperature increases is shown in Fig. 6.4 for T/P91 and T/P92 steels. Clearly, the dependence of Young’s modulus on temperature is very significant. The slope of the curves increases with increasing temperature. It is also
Young’s modulus, E (GPa)
210 200 190 180 170 160
P91 P92 200
300
400 500 Temperature, T (°C)
600
700
6.4 Dependence of Young’s modulus of T/P91 (Haarmann et al., 2002) and T/P92 (Richardot et al., 2000) on testing temperature.
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225
worth noting that the Young’s modulus of T/P92 in Fig. 6.4 is it is slightly higher than that of T/P91 at the same testing temperature. At normal service temperatures, Young’s modulus of creep-resistant steels can be as low as half that at room temperature. Therefore, it is necessary to give the testing temperature when reporting Young’s modulus values, otherwise confusion may occur.
6.3
Thermal properties of creep-resistant steels
6.3.1
Thermal expansion
Thermal expansion describes the response of a material to heat input. The extent of expansion or contraction depends on the material and temperature change and is commonly described by the linear coefficient of thermal expansion or CTE in short. Assume a steel bar of initial length l0 undergoes a temperature increase from T to T + ∆T and its length changes to l = l0 + ∆l. The CTE, α, is defined as: α = ∆l l 0 ∆T
[6.18]
Therefore, CTE is the relative increase in length of material when temperature increases by one degree Kelvin. It is clear that the metric unit of CTE is m · m–1 · K–1 or K–1. In practical applications, the expansion or contraction of a material is very small, this unit is too big and 10–6K–1 is commonly used instead. The values of some common power plant creep resistant ferritic/ martensitic and austenitic steels are listed in Table 6.2. Design of the structures or components must take into consideration the change of dimension of the material as temperature changes during normal service. This means that in the design process, the structure must be allowed to accommodate the dimensional change of the material from room temperature to normal operational temperature. This produces the requirement to put expansion joints or expansion loops in the structure. If the component is restrained, thermal expansion may cause buckling or bending of the component. Buckling and bending can also occur when two steels with markedly different thermal expansion coefficients are fabricated together and subsequently heated or cooled. Another effect is thermal stress. When a component expands, tensile forces are created. On the other hand, compressive forces are created when a component contracts. If we recall that strain in a pure tensile situation is defined as the relative increase of the length of the material, (Equation [6.6], we can define a thermal strain here as well, ε thermal = ∆l l0
[6.19]
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Creep-resistant steels
From Equation [6.18], we obtain εthermal = α∆T
[6.20]
32
212
Temperature (°F) 392 572 752
932
20
1112 11.2
8.4
15
10 0
100
200 300 400 Temperature (°C)
500
5.6 600
Coefficient of linear expansion (10–6 /°F)
Coefficient of linear expansion (10–6 K)
Therefore, thermal strain is directly proportional to the change in temperature. However, it should be noted that the CTE varies with temperature, thus Equation [6.20] only can be applied in a small temperature range or in a sense of average behaviour. In fact, CTE as defined in Equation [6.18] is the mean CTE of the material over the temperature range T to T + ∆T. Outside this range, the CTE might have a different value. To obtain a CTE at temperature T, the range ∆T must be acceptably small. Therefore, in reporting thermal expansion coefficients of steels, it is common for the mean values of CTE from room temperature to a test temperature to be cited. Of course, these mean values are also dependent on the test temperature. Figure 6.5 shows the mean thermal expansion of three creep-resistant steels, namely T/P91, T/P22 and TP304H, as a function of test temperature. In the temperature range shown, mean CTE for all materials increases approximately linearly with increasing temperature and the increase in CTE from room temperature to 600°C is about 25%. Of course, in design of power plant components, the mean CTE in the range from room temperature to the operational temperature is more relevant. The difference in thermal expansion coefficient between different creepresistant steels is significant. The CTEs for austenitic steels are about 50% higher than those for ferritic and martensitic steels, as shown in Fig. 6.5 and Table 6.2. CTE generally increases with increasing bond energy (Incropera et al., 2006). Bond energy depends on the nature of the interaction between atoms forming the solid and the bond length. The stronger the interaction, the higher is the bond energy. On the other hand, the shorter the bond length,
6.5 The temperature dependence of mean linear thermal expansion coefficients of T/P91, T/P22 and TP304H (Haarmann et al., 2002).
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the higher is the bond energy. In the case of steels, the main bonding is between iron atoms and the nature of interaction is the same in all the cases of ferritic, martensitic and austenitic steels. Therefore the bond energy is mainly determined by the bond length. The bond length can be approximated using the equilibrium distance between the centres of atoms. Thus, the bond energy increases with decreasing interatomic distance. As mentioned earlier, the FCC austenitic steel is denser than the BCC ferritic/martensitic steel, hence it has a higher thermal expansion coefficient. However, the bond strength may be altered by adding different alloying atoms to the material. This provides the possibility for developing low thermal expansion steels via alloying. The significant variation of CTE within the ferritic/martensitic group is good evidence of such an alteration. The variations of CTE as a function of the main alloying elements, namely carbon and chromium for ferritic/martensitic steels and carbon, chromium and nickel for austenitic steels are shown in Figs. 6.6 and 6.7. Figure 6.6 shows clear trends of the decrease of CTE with increasing concentration of both carbon and chromium in the ferritic/martensitic steels. The only exception in Fig. 6.6(a) is T/P22 which has exceptionally low chromium content (2.25 wt%). This exception is believed to be due to the effect of chromium which outweighs the effect of carbon. However, the effects of both carbon and chromium on CTE of austenitic steels are not so clear, as shown in Fig. 6.7(a) and (b). On the other hand, Fig. 6.7(c) shows a clear decrease in CTE with increasing nickel content of the steels. On the whole, it is reasonable to conclude that CTE decreases with increasing concentration of C, Cr and Ni in both ferritic/ martensitic and austenitic steels. Yamamoto et al. (2003) have carried out some regression analysis of CTE data for Ni-base superalloys and found the contribution of different alloying elements to CTE from room temperature to 700°C could be described by the formula below: α700°C = 13.8732 + 7.2764 × 10–2 × [Cr] + 3.751 × 10–2 × [Ta + 1.95Nb} + 1.9774 × 10–2[Co] + 7.3 × 10–5 × [Co] × [Co] – 1.835 × 10–2 × [Al] – 7.9532 × 10–2 × [W] – 8.2385 × 10–2[Mo] – 1.63381 × 10–1 × [Ti] [6.21] where α700°C is the mean thermal expansion coefficient from room temperature to 700°C in 10–6 K–1 and [element name] is the concentration of the element in weight percent. This kind of equation is very useful in developing low CTE alloys. Thermal stress caused by thermal expansion is described by Equation [6.1], that is, σthermal = Eεthermal = α∆TE. Localized stresses from thermal
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Creep-resistant steels 15 T/P22
CTE, α (10–6 K–1)
14
T/P92
13
T/P91 E911
12
HCM12A AISI410
11 0.06
0.08
0.10 0.12 0.14 Carbon content, C (wt%)
0.16
(a) 15 T/P22
CTE, α (10–6 K–1)
14
T/P92
13
T/P91 E911
12
HCM12A AISI410
11 0
2
4
6 8 10 Cr content, C (wt%)
12
14
16
(b)
6.6 Linear thermal expansion coefficient of some ferritic/martensitic creep-resistant steels as a function of the content of (a) carbon and (b) chromium at 600°C.
expansion during heating and cooling can contribute to the problem of stress corrosion cracking in an environment which would not normally attack the steel. These applications require design to minimize the adverse effects of temperature variations such as the use of expansion joints to permit movement without distortion of the component and the avoidance of notches and abrupt changes of section. Thermal stress is also a contributor to fatigue as discussed in the introduction.
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Physical and elastic behaviour of creep-resistant steels
CTE, α (10–6 K–1)
20
229
304H
19 304L 321
304 347
316
A286
18 309
310
17
16 0.00
0.05
0.10 0.15 0.20 Carbon content, C (wt%) (a)
0.25
21
CTE, α (10–6 K–1)
20
304H
19 321 347
18 A286
304/304L
309
316
310
17
16 15
CTE, α (10–6 K–1)
20
19
18 21 24 Cr content, C (wt%) (b)
27
304H
304/304L 321 347
18 309
316
A286 310
17
16 8
12
16 20 Ni content, C (wt%) (c)
24
28
6.7 CTE as a function of the content of (a) carbon, (b) chromium and (c) nickel in some austenitic creep-resistant steels.
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230
6.3.2
Creep-resistant steels
Thermal conductivity
In physics, thermal conductivity, k, is the property of a material that measures its capability to conduct heat. Heat transfer by conduction involves transfer of energy within a material without any motion of the material as a whole. The rate of heat transfer depends on the temperature gradient and thermal conductivity of the material. Higher thermal conductivity indicates higher ability for transferring heat. This implies that if the same volume of two different materials is heated, the one with higher thermal conductivity will have lower thermal gradient throughout the bulk. Therefore, higher thermal conductivity is desirable in at least two ways. First, higher thermal conductivity reduces the temperature gradient within a component and therefore reduces thermal stress caused by temperature change. This has a direct influence on the thermal stress parameter, as indicated in Equation [6.2]. Second, higher thermal conductivity allows thicker sections of material to be used, hence decreasing the demand on the strength of the material. The definition of thermal conductivity is straightforward. Assume a thin slab of a material with cross-sectional area A and thickness ∆x; one side of the slab is kept at temperature T + ∆T and the other kept at T, as shown in Fig. 6.8, then the temperature gradient across the slab is ∆T . The heat Q ∆x transferred through area A from the high temperature to low temperature side is directly proportional to the temperature gradient ∆T , area A and the time ∆x ∆t during which the heat transfer takes place, that is: Q = – κ ∆T A∆t ∆x
Q
[6.22]
A
T + ∆T
T ∆x
6.8 Definition of thermal conductivity.
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The minus sign here signifies that the heat transfer is from the higher temperature to lower temperature side, that is, opposite to the direction of temperature gradient. The proportionality constant is the thermal conductivity of the material. Rearranging Equation [6.22], κ=–
Q ∆T A∆t ∆x
[6.23]
Therefore, thermal conductivity is the quantity of heat, Q, transmitted in unit time through a unit normal area caused by a unit temperature gradient under steady state conditions. Different materials transfer heat in different ways and therefore have vastly different thermal conductivities. Gases transfer heat by direct collisions between molecules and their thermal conductivities are low compared to most solids since they are dilute media. Non-metallic solids transfer heat by lattice vibrations. In metals, heat transfer is accomplished by mobile electrons which also participate in electrical conduction. Therefore, metals are much better thermal conductors than non-metals. Because the carriers of heat in metals are the mobile electrons, the thermal conductivity of metals is determined by the number density and mobility of the mobile electrons in the metal. Higher mobile electron density and higher electron mobility result in higher thermal and electrical conductivity. Thermal conductivities of some creepresistant steels are listed in Table 6.2. Thermal conductivity of ferritic/martensitic steels is about two-thirds of that of pure iron, but is about 50% higher than that of austenitic steels (Table 6.2). As mentioned before, austenitic steels have denser structures than ferritic/ martensitic steels; therefore the number density of mobile electrons is higher than that in ferritic/martensitic steels. This should lead a higher thermal conductivity for austenitic steels. However, owing to the smaller inter-atomic distance (this is also a measure of the atom size), the attraction between the nuclei and the electrons is stronger and the mobility of mobile electrons in austenitic steels is much lower than that in ferritic/martensitic steels. This effect outweighs the effect of the number density of mobile electrons and therefore the thermal conductivity of austenitic steels is lower than that of ferritic/martensitic steels. As thermal conductivity is a kind of measure of the effectiveness of energy transfer by the collision of mobile electrons, anything that affects the mobility of electrons may have an influence on thermal conductivity. For example, grain boundaries, second phase particles and non-metallic inclusions in steels all affect the thermal conductivity of the material. Figure 6.9 shows thermal conductivity as a function of steel composition for (a) carbon steels as a function of carbon content; (b) chromium steels as a function of chromium content; (c) nickel steels as a function of nickel concentration; and (d) tungsten
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95 90 85 80 75 70 65 60 0.4
0.6 0.8 1.0 1.2 1.4 Carbon content, C (wt%) (a)
1.6
Thermal conductivity, k (W m–1 K–1)
Creep-resistant steels
Thermal conductivity, k (W m–1 K–1)
Thermal conductivity, k (W m–1 K–1)
Thermal conductivity, k (W m–1 K–1)
232
140 120 100 80 60 40 20 0 0
20 40 60 Ni content, C (wt%) (c)
140 120 100 80 60 40 0
5 10 15 Cr content, C (wt%) (b)
20
130 120 110 100 90 80
80
0
2
4 6 8 W content, C (wt%) (d)
10
6.9 Thermal conductivity of some steels as a function of content (a) carbon, (b) chromium, (c) nickel and (d) tungsten. Data from www.EngineersEdge.com.
steels as a function of tungsten concentration. Figure 6.9 is plotted according to data from www.EngineersEdge.com. As can be seen from Fig. 6.9, there is a general trend of decreasing thermal conductivity with increasing concentration of alloying element. The effect is significant. For example, the thermal conductivity of chromium steel decreases to less than one-third of its original value when the concentration of chromium increases from 0 to about 20 wt%. Interestingly, there is an upturn in the relation between thermal conductivity of nickel steel and the concentration of nickel at 40 wt% (Fig. 6.9(c)). Therefore, the minimum of thermal conductivity appears to occur at a nickel concentration of ~50 wt%. Further increase of nickel content results in the material becoming a nickelbase alloy rather than a steel. Nickel-base alloys have tighter composition specifications than steel, which means that they contain fewer impurities. Therefore, thermal conductivity increases.
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233
Creep-resistant steels are more complicated, often containing up to around 15 or more alloying elements. Therefore, no straightforward relationships between conductivity and composition are expected. However, it still holds that there is a general trend of decreasing thermal conductivity with increasing concentration of alloying element, as shown in Fig. 6.10 for some common ferritic/martensitic creep-resistant steels. Thermal conductivity of ferritic/ martensitic steels is much lower than that of pure iron and decreases with increasing concentration of carbon or chromium. The trend in Fig. 6.10(a) is not so clear owing to the exceptionally high thermal conductivity of T/P22.
Thermal conductivity, k (W m–1 K–1)
40 Pure Fe
35 T/P22
T/P91
30
T/P92
HCM12A E911
25
AISI410 0.00
0.04 0.08 0.12 Carbon content, C (wt%) (a)
0.16
Thermal conductivity, k (W m–1 K–1)
40 Pure Fe
35 T/P22
30
T/P92 E911
HCM12A
T/P91 25
AISI410 0
2
4 6 8 10 Chromium content, C (wt%) (b)
12
14
6.10 Thermal conductivity of some common ferritic/martensitic creep-resistant steels as a function of (a) carbon and (b) chromium content.
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Creep-resistant steels
This is because T/P22 has very low chromium content and the effect of lower chromium concentration outweighs that of higher carbon concentration. Similar trends are found for austenitic steels as shown in Fig. 6.11. However, thermal conductivity values for austenitic steels are much lower than those for ferritic/martensitic steels and its variation as a function of composition is smaller than that in the case of ferritic/martensitic steels. Thermal conductivity of creep-resistant steels is also a function of temperature. As mentioned earlier in this section, metallic solids conduct heat via the collision of mobile electrons and thermal conductivity is a measure of the effectiveness of the energy transfer during the collision. Generally speaking, electrons move faster at higher temperatures and the probability of collision with others is much higher. Therefore, energy can be more effectively transferred from locations with higher temperature to locations where temperature is lower and thermal conductivity generally increases with increasing temperature. Figure 6.12 shows that this is true for T/P91, TP304H and T/P22 at low temperatures.
6.4
Electrical resistivity and conductivity of creepresistant steels
Metals, including steels, are good electrical conductors. However, they do resist the flow of electrical current. The ability of a metal to resist electrical flow is represented by its electrical resistivity. If a sample of metal with uniform cross-sectional area A and length l has an electrical resistance R, then its resistivity, ρ is defined by R=ρ l A
[6.24]
ρ= R l/ A
[6.25]
or
Therefore, the resistance of a sample of steel is inversely proportional to its conducting area. This has implications for the monitoring of crack initiation and propagation in steels. If a crack grows or propagates in a steel component, the conducting area will be reduced and the resistance to electrical current will be increased. The simplest way of measuring resistance is to use Ohm’s law, that is: I = V or R = V [6.26] R I where I is the electrical current and V the applied electrical potential or voltage. In SI units, the unit of electrical resistivity is ohm-meter or Ωm. For steels, the resistivity is in the range 1~10 × 10–7 Ωm and some values are listed in Table 6.2.
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Thermal conductivity, k (W m–1 K–1)
Physical and elastic behaviour of creep-resistant steels 40
235
Pure Fe
35
30
25
A286 304/304H 321/347 310
304L 20
316
309 15
Thermal conductivity, k (W m–1 K–1)
0.00
0.05 0.10 0.15 Carbon content, C (wt%) (a)
0.20
40 Pure Fe 35
30
25
A286 304L/304/304H 321/347
20
316
310 309
15
Thermal conductivity, k (W m–1 K–1)
0
5 10 15 20 Chromium content, C (wt%) (b)
25
40 Pure Fe 35
30
25 304L/304 304H
20
321/347
A286
316 309
310
15 0
5
10 15 20 Nickel content, C (wt%) (c)
25
30
6.11 Thermal conductivity of some austenitic creep-resistant steels as a function of (a) carbon, (b) chromium and (c) nickel content.
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Creep-resistant steels
32
212
Temperature (°F) 392 572 752
932
Thermal conductivity (Wm K–1)
40
1112 22.4
T/P22 16.8
30 T/P91
11.2
20 TP304H
10 0
100
200 300 400 Temperature (°C)
500
5.6 600
Thermal conductivity ((Btufth–1 °F–1)
236
6.12 Examples of thermal conductivity as a function of temperature for some creep-resistant steels.
Electrical conductivity, often represented using the Greek letter σ, is a measure of the ability of a conductor to conduct electrical current. It is defined as the reciprocal of electrical resistivity, i.e. σ= 1 ρ
[6.27]
Metals conduct electricity by the movement of free electrons, which also conduct heat as discussed earlier. Therefore, electrical conductivity of steels is related to their thermal conductivity. Generally speaking, a good thermal conductor is also good at conducting electricity. This relationship is described by the Wiedemann–Franz Law: κ = LT σ
[6.28]
where κ is thermal conductivity, σ is electrical conductivity and T is the absolute temperature. The proportional constant L is called Lorenz number and it has the value: 2 2 L = π k2 = 2.45 × 10 –8 W Ω K –2 3e
where k is the Boltzmann constant and e is the charge of electrons. Electrical resistivity of creep-resistant steels is dependent on various factors, such as the purity of the material and the temperature. From Table 6.2, electrical resistivity of both ferritic/martensitic and austenitic steels is much higher than that of pure iron, indicating that electrical resistivity increases with the degree of alloying. For austenitic steels, electrical conductivity does
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237
not differ very much from alloy to alloy, but higher than the usual values, as for ferritic/martensitic steels, indicates that the structure of the material plays a vital role in conducting electrical current. Electrical resistivity is also temperature dependent. Examples of the dependence of electrical resistivity on temperature for ferritic/martensitic and austenitic steels are shown in Fig. 6.13. Clearly, electrical resistivity of both ferritic/martensitic and austenitic steels increases with increasing temperature. Austenitic steels have much higher electrical resistivity than ferritic/ martensitic steels, that is, ferritic/martensitic steels are much better electrical conductors than austenitic steels, which is as expected because ferritic/ martensitic steels have much higher thermal conductivity, as discussed in Section 6.3. It is interesting to note that the temperature dependence of electrical conductivity of steels is opposite to that of thermal conductivity. In Section 6.3 we have shown that thermal conductivity generally increases with increasing temperature. Figure 6.13 shows that electrical resistivity of steels rises as temperature increases. Therefore, electrical conductivity decreases with increasing temperature. This is due to the different mechanisms of conducting heat and electricity. Although metals conduct both heat and electricity by the movement of mobile electrons, the nature of the movement is different. Heat is transferred from a location of higher temperature to that of lower temperature by collision of electrons which depends on the random movement of mobile electrons, no directional motion of the electrons as a whole is involved. Temperature in fact is a measure of the intensity of such random movement and the higher the temperature, the higher is the frequency of collision. As a result, mobile electrons conduct heat more effectively at higher temperatures. It can be shown that thermal conductivity is directly
Electrical resistivity, ρ (Ωm)
120
100
80
60 9Cr–0.12C–1Mo AISI 321/347 40 0
200
400 600 Temperature, T (°C)
800
1000
6.13 Examples of temperature dependence of electrical resistivity of ferritic/martensitic and austenitic creep-resistant steels.
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Creep-resistant steels
proportional to the average speed of electrons, which in turn is directly proportional to the square root of absolute temperature. Therefore, thermal conductivity increases with increasing temperature. In the case of conducting electricity, the mobile electrons move in the direction opposite to the electrical current as a whole, in addition to the random movement. Electrical conductivity is dependent on the mass directional movement (or sometimes called drift) of the electrons rather than the random motion. In fact, the collision between the electrons owing to random movement (and lattice vibration) scatters the electrons and prevents them from drift. Therefore, electrical conductivity is inversely proportional to the average speed of Brownian motion of electrons. At higher temperatures, the random movement is more intense and it is more difficult for electrons to drift in a certain direction. Therefore electrical conductivity decreases with increasing temperature. Combining κ ∝ Franz Law, Equation [6.28].
6.5
T and σ ∝ 1 , we obtain the Wiedemann– T
Implications for industries using creep-resistant steels
Although general physical properties of creep-resistant steels are not discussed so much as creep strength in the literature, they have important implications for the design and service of the component. Generally speaking, lower modulus of elasticity, lower thermal expansion and higher thermal conductivity are desirable if secondary stresses caused by heat input and temperature changes are considered. As two-shifting and weekend close down become common, thermal fatigue caused by such operations at power plants must be taken into consideration. Thermal stress caused during cool down and heat up of power plant components is directly linked to all three parameters discussed in this chapter. Higher thermal expansion coefficient leads to higher thermal strain. Higher Young’s modulus results in higher thermal stress at the same thermal strain. Material with lower thermal conductivity cannot conduct heat effectively and gives rise to a higher temperature gradient and therefore a higher thermal strain. If the strength of the material is acceptable under operational conditions, the material with lower thermal expansion, higher thermal conductivity and lower Young’s modulus should be used. If this material cannot meet the strength requirements and the use of materials with higher thermal expansion, lower thermal conductivity and higher Young’s modulus are inevitable and acceptable ways of reducing or avoiding related problems must be found. For example, based on calculations for thick section austenitic pipework, Fleming et al. (1997) pointed out that to avoid unacceptable thermal stresses in a fully austenitic plant it would take several days to heat up and cool down.
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239
The combination of high thermal expansion and low thermal conductivity means that precautions must be taken to avoid adverse effects. For example, during the welding of austenitic steels, measures such as low heat input and dissipation of heat using other media may be required. When heated from room temperature (20°C) to 600°C, an austenitic pipe of length 1 m may expand by 1 cm. Expansion loops/joints may need to be used to accommodate such a dimensional change. Precautions must also be taken to reduce stress corrosion cracking caused by or assisted by thermal stress caused by low thermal conductivity and high thermal expansion. On the other hand, higher thermal conductivity allows thicker section components to be used and this reduces the demand on the strength of the material.
6.6
Future trends
Owing to their low thermal expansion and high thermal conductivity, tremendous efforts have been made to develop better ferritic/martensitic steels. In recent years, material scientists and engineers seem to have concluded that with less than 11–12 wt% chromium content, ferritic/martensitic steels cannot overcome the problem of steam oxidation. Therefore, much attention has been paid to developing higher chromium ferritic/martensitic steels. However, the attempt has not been successful. These higher chromium ferritic/ martensitic steels seem to promote the formation of the so-called Z phase which has detrimental effects on the long term creep strength of the alloy. For both environmental and economic reasons, power plant designers are keen to operate plant at higher temperatures and pressures to improve the efficiency. Under such conditions, the problem of oxidation of ferritic/ martensitic steels becomes more serious. Thus, more recently, materials scientists and engineers have suggested abandonment of ferritic/martensitic steels and development of new austenitic steels for higher temperature operation. Although thermal expansion may be reduced via the route of alloying, it is difficult to obtain low thermal stress during heat up or cool down periods via alloying because alloying inevitably decreases thermal conductivity of the alloy. Under such considerations, means of avoiding unacceptable thermal stress must be found in the future at the design stage to accommodate higher thermal expansion and lower thermal conductivity of austenitic steels.
6.7
References
Alyousifa O M and Nishimura R (2006), ‘The effect of test temperature on SCC behaviour of austenitic stainless steels in boiling saturated magnesium chloride solution’, Corrosion Science, 48, 4283–4293. Ashby M F and Jones R H (1996), Engineering Materials 1 – An Introduction to their Properties and Applications, Butterworth Heinemann, Oxford.
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Brandes E A and Brook G B (eds) (1992), The Smithells Metal Reference Book, Butterworth Heinemann, Oxford. Chênea J, Brassa A-M, Trabucb P and Gastaldi O (2007), ‘Role of microstructure and heat treatments on the desorption kinetics of tritium from austenitic stainless steels’, Journal of Nuclear Materials, 360, 177–185. Cottrell A (1995), An Introduction to Metallurgy, 2nd edition, The Institute of Materials, London. De Cicco H, Luppo M I, Raffaeli H, Di Gaetano J, Gribaudo L M and Ovejero-García (2005), ‘Creep behavior of an AISI286 type stainless steel’, Materials Characterization, 55 97–105. Fleming A, Buchanan L W and Maskell R V (1997), ‘Overview of material development requirements for boiler plant’, Materials Issues in Heat Exchangers in Boilers Conference, Starr F and Meadowcroft B (eds), IOM London, 109–117. Haarmann K, Vaillant J C, Vandenberghe B, Bendick W and Arbab A (2002), The T91/ P91 Book Vallourec and Mannesmann tubes, Boulogne. Hightempmetals (2007), www.hightempmetals.com Incropera F P, DeWitt D P, Bergman T L and Lavine A S (2006), Fundamentals of Heat and Mass Transfer, 6th edition, Wiley, New York. Klueh R L (2005), ‘Elevated temperature ferritic and martensitic steels and their applications to future nuclear reactors’, International Materials Review, 50 (5), 287–310. Laha K, Kyono J and Shinya N (2007), ‘An advanced creep cavitation resistance Cucontaining 18Cr–12Ni–Nb austenitic stainless steel’, Scripta Materialia, 56, 915– 918. Matweb (2007), www.matweb.com Ravi Kumar B, Das S K, Mahato B, Arpan Das and Ghosh Chowdhury S (2006), ‘Effect of large strains on grain boundary character distribution in AISI 304L austenitic stainless steel’, Materials Science and Engineering A, 454–455, 239–244. Richardot D, Vaillant J C, Arbab A and Bendick W (2000), The T92/P92 book, Vallourec & Mannesmann Tubes, Boulogne. Sandmeyersteel (2007), www.sandmeyersteel.com Skelton P and Beckett B E (1987), ‘Thermal Fatigue Properties of Candidate Materials for Advanced Steam Plant’, Conference on Advances in Material Technology for Fossil Power Plants, ASM, Ohio, 359–366. Starr F (2002), ‘Potential issues in the cycling of advanced power plants’, OMMI, 1 (1), 1–19. Tsai M C, Chiou C S, Du J S and Yang J R (2002), ‘Phase transformation in AISI 410 stainless steel’, Materials Science and Engineering A, 332, 1–10. Yamamoto R, Kadoya Y, Kawai H, Magoshi R, Noda T, Hamano S, Ueta S and Isobe S (2003), ‘New wrought Ni-based superalloys with low thermal expansion for 700C steam turbines’, Energy Technology, 21, 1351–1360. Ye W, Li Y and Wang F (2006), ‘Effects of nanocrystallization on the corrosion behaviour of 309 stainless steel’, Electrochimica Acta, 51, 4426–4432. Yoo, Y-S (2004), ‘Study of LBB assessment methodology applied to a 12Cr series ferrite steel piping structure for FBRs’, 24, 27–36. Zelada-Lambri G I, Lambria O A and Rubiolob G H (1999), ‘Amplitude dependent damping study in austenitic stainless steels 316H and 304H. Its relation with the microstructure’, Journal of Nuclear Materials, 273, 248–256.
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7 Diffusion behaviour of creep-resistant steels H. O I K A WA and Y. I I J I M A, Tohoku University, Japan
7.1
Introduction
In the early 1950s, Sherby and coworkers (e.g., Sherby et al., 1954) analysed creep data and diffusion data of several metals and claimed that the temperature dependence of both phenomena is similar to each other. This finding first gave us a sound physical base for discussing creep mechanisms. In those days, however, creep data were very limited and diffusion data were also a few and their values were scattered over a wide range. Therefore, they had to base their discussion on average values of these limited data. Today, the view is well established (Sherby and Burke, 1967; Mukherjee, et al., 1969) that the temperature dependence of creep at high temperature is close to that of diffusion in pure metals (see Fig. 7.1). A similar correlation has also been found in solid solution alloys (Monma et al., 1964). Diffusion is one of most fundamental processes governing creep deformation. In this chapter diffusion behaviour in metals and alloys will be outlined and then the role of diffusion in creep deformation will be discussed. Finally, some fundamental diffusion data which are deemed to be useful in discussing creep of steels will be cited.
7.2
Diffusion and creep
7.2.1
Activation energies
It is worthwhile noting that there is an essential difference in the physical meaning of temperature dependence of creep and of diffusion, although the temperature dependences of these two phenomena are close to each other under some conditions. Diffusion in simple metals can be recognized as a thermally activated process, where the activation energy is the sum of the formation energy and the migration energy of vacancies. The pre-exponential (or frequency) term 241 WPNL2204
242
Creep-resistant steels 1000
Mo
500
Qd / (kJ mol–1)
Ta Ni Pt Cu 200 β Ti Mg Pb 100
Ag
α Fe
Nb γ Fe β Co
Au α Ti
Al
Zn Cd
50
100
200 Qc /(kJ mol–1)
500
1000
7.1 Correlation between the activation energies of high-temperature creep, Qc, and of lattice self-diffusion, QD, in pure metals.
is essentially the number of lattice sites in the system and not a function of temperature. Creep is a time-dependent deformation process which is a result of complex dislocation behaviour. Its temperature dependence is influenced by many factors which may have their own temperature dependence. The so-called activation energy of creep is a temperature dependence of the steady-state (or minimum) creep rate under a given stress, derived on the assumption that one thermally activated process occurs in the phenomenon. Therefore, the activation energy of creep is an apparent one in the sense of rigorous thermal activation rate theory. It must be kept in mind during discussion of creep mechanisms that the pre-exponential term in the form of a thermally activated rate equation contains many factors which might depend on temperature even under constant stress.
7.2.2
Time-dependent deformation and diffusion
Creep deformation is time-dependent straining under a constant applied stress, or under a given load in many cases. At very high temperature under very low stress, ‘diffusional creep’ occurs in (pure) metals. Under these conditions, creep strain arises directly from the movement of atoms. The temperature dependence (the activation energy) of creep is the same as that of vacancy diffusion. At temperatures higher than about a half of the melting temperature, 0.5Tm, and under ordinal creep conditions, the temperature dependence of
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243
creep (strain rate) is also similar to that of diffusion. This similarity, however, does not mean that creep strain arises directly from the movement of atoms. Rather, the similarity indicates that the rate-controlling step of ‘so-called’ high-temperature creep is a kind of restoration process relating intimately to diffusion. When steady-state creep-rates and/or minimum creep rates are taken as the parameter of a creep process, a similarity in the temperature dependence between creep and diffusion can be observed, at least in (pure) metals and many simple alloys. Correlation between diffusion and creep behaviour of practical alloys under practical creep conditions is not simple, as in the case of pure metals and solid solution alloys at high temperatures. In these conditions temperature is usually less than 0.5Tm, significant structure changes occur in the matrix and the influence of surrounding atmosphere becomes obvious with time. These factors strongly affect the creep behaviour of material and any simple relation between creep and diffusion is difficult to observe.
7.2.3
Influence of short-circuit diffusion
In creep of metals and alloys, the essential rate-determining stage is a kind of restoration process which is governed by diffusion of vacancies and/or alloying elements. At temperature higher than about 0.5Tm, diffusion of vacancies/atoms through the nearly perfect crystal lattice governs creep process and the temperature dependence of creep rate is similar to that of the lattice diffusion of vacancies/atoms. At temperatures lower than about 0.5Tm, however, the effects of shortcircuit diffusion (see Section 7.3.4) become obvious in many cases. The absolute values of short-circuit diffusion coefficients are always larger than the lattice diffusion coefficients, but the net effect of these types of diffusion is not always large, because their cross-sections are small. In some cases, a change in the activation energy of creep is observed at temperatures below 0.5Tm. This change in the activation energy can be explained not by a change in creep mechanism, but simply by the effect of short-circuit diffusion (Lüthy et al., 1980).
7.3
Diffusion characteristics
7.3.1
Moving species: vacancies
In most metallic materials, direct exchange of an atom with an adjacent atom hardly ever occurs without vacant lattice sites. In pure metals, therefore, (long-range) atom movement has the same meaning as vacancy travel but in the opposite direction. The (self-)diffusion coefficient of constituent atoms
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Creep-resistant steels
using radio isotopes, D*, is essentially the same as the diffusion coefficient of vacancies. In (binary) solid solutions consisting of the solute atoms B with the solvent atoms A, the measured diffusion coefficients using isotopes, D*B or D*A, at a fixed concentration of B, are again an indication of the movement of vacancies replacing the lattice site of atom B or atom A. These coefficients do not relate to the concentration gradient in a material. This type of diffusion coefficient in alloys have been measured in limited cases only. When the site exchange between atom A and vacancies occurs independently of the presence of other kinds of atoms (B), diffusion of atom A and atom B occurs as parallel reactions independent with each other. The resultant vacancy diffusion is simply the sum of vacancies exchanging with atoms A and with atoms B. Diffusion of vacancies under this condition, D , can be expressed by:
D = NA D*A + NB D*B
[7.1]
Here, NA and NB are the mole fraction of element A and B, and D*A and D*B are self-(tracer) diffusion coefficients in the alloy, respectively. Diffusion of this type is affected preferentially by the faster atoms. This type of diffusion results in a segregation of atoms, because of the difference between D*A and D*B . When the condition does not allow segregation (long-range concentration fluctuation), then vacancy diffusion must proceed as a series of reactions of the movement of atoms A and atoms B, D ′ :
D ′ = D*AD*B/(NAD*B + NBD*A)
[7.2]
This equation can be more easily understood from another form of expression: 1/ D ′ = NA/D*A + NB/D*B
[7.3]
This formulation is analogous to series-combined resistance which is a simple sum of resistance of each component. Diffusion of this type is affected preferentially by the diffusion of slower atoms. The climb motion of an edge dislocation is controlled by vacancy formation (or annihilation) at the elementary jog and the moving away of this vacancy from the jog. The diffusion coefficients suitable for use in the analyses are not D (parallel vacancy diffusion coefficients), but D ′ (combined vacancy diffusion coefficients).
7.3.2
Moving species: constituent atoms
In diffusion experiments with alloys, usually two pieces of alloys with different levels of the solute (B) concentration are welded and the concentration change of B with time (homogenization process) is measured. The available diffusion
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245
coefficients are called interdiffusion coefficients, or chemical diffusion coefficients, and expressed customarily as D˜ . In interdiffusion reactions, the concentration profile changes with time; usually a kind of homogenization process proceeds. When diffusion of atom A occurs independently of that of atom B, then the resultant diffusion is a simple sum of movement of A and B (in the inverse direction from each other under ordinal conditions) and the so-called Kirkendall effect can be seen. Interdiffusion coefficients, D˜ , can be estimated from the equation (Darken equation): D˜ = NAD*B + NBD*A
[7.4]
When a change of concentration profile is not allowed, but the change in the position only is allowed, then the movement of atoms (solute and solvent) is expressed by: D˜ ′ = D*AD*B/(NAD*A + NBD*B)
[7.5]
This equation can be more easily understood from another form of expression: 1/ D˜ ′ = NA/D*B + NB/D*A
[7.6]
This formulation is analogous to series-combined resistance which is a sum of resistance of each component. In analyses of the glide motion of an edge dislocation, which is controlled by diffusion of the solute atmosphere, the diffusion coefficient suitable for use in the analyses is not D˜ (interdiffusion, a parallel sum of independent diffusion of constituent components), but D˜ ′ (combined diffusion coefficients of constituent components). When a solute atmosphere is being formed around a fresh dislocation, the diffusion coefficients that should be used in the analyses are ordinary interdiffusion coefficients. An example of concentration dependence of several types of diffusion coefficient is shown schematically in Fig. 7.2 for an A–B binary solid solution alloy system. Diffusion coefficients suitable for employment in analyses of creep phenomena are those under the condition of a no-concentration gradient, that is, D ′ and/or D˜ ′ , not D and/or D˜ . Unfortunately most diffusion data reported are D and/or D˜ . Values of D ′ and/or D˜ ′ are available in very limited cases.
7.3.3
Diffusion paths: lattice diffusion
In common metals and alloys, atom movement occurs through interchange of an atom in a crystal lattice site with a vacant site. Diffusion in nearly perfect crystals is called lattice diffusion or volume diffusion. Lattice diffusion in metals depends on Tm and the crystal lattice system. Among metals with
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DB* (self)
Diffusion coefficient (log. scale)
246
DA* (imp) DB*
~ D′
~ D
D
D′ D* A
DB* (imp)
DA* (self) 0 (A)
0.5
1.0 (B)
Mole fraction of B
7.2 Schematic diagram of concentration dependence of diffusion coefficients in an A–B binary solid solution alloy.
the same crystal lattice system, (self-) diffusion coefficients depend closely on their Tm value. In metals with a higher Tm, diffusion coefficients are smaller than those in metals with lower Tm values. In metals with closepacked lattice systems (e.g. face centred cubic (fcc) and hexagonal close packed (hcp)), diffusion is slower than that in those with less close-packed lattice systems (e.g. body centred cubic (bcc)). In pure iron (see Fig. 7.3) self-diffusion coefficients in the γ-phase are more than two orders of magnitude smaller than those in the δ- and α-phases.
7.3.4
Diffusion paths: short-circuit diffusion
There are some diffusion paths that go through imperfect crystal lattice sites and these are termed short-circuit diffusion. Typical examples are those through grain boundaries, Dgb, and over the surface, Ds. Diffusion along dislocation cores, Dd, is also a typical example of short-circuit diffusion. Short-circuit diffusion coefficients are larger than lattice diffusion coefficients even at the melting temperature of the crystal. The difference between the lattice diffusion coefficient and other short-circuit diffusion coefficients increases with decreasing temperature (see Fig. 7.4). The difference can be many orders of magnitude. The practical effect of short-circuit diffusion, however, is not very large, at least above 0.5Tm, because the effective crosssections of short-circuit diffusion are usually very small. At lower temperatures the effect of short-circuit diffusion becomes obvious in some cases. In these
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T (K) 1800
1400
1000
800
10–10 δ 10–12
* (m2s–1) DFe
10–14
αp γ
10–16 10–18 αf –20
10
10–22
0.6
0.8
1.0
1.2
T –1 (10–3 K–1)
7.3 Lattice self-diffusion coefficients in pure iron as a function of temperature (Oikawa,1982).
10
1100
–6
900
T (K) 800 700
600
10–7 10–8 10–9 10–10
10–12 10–13 10–14
Dgb
10–15 10–16
Dd
10–17
Tα–γ = 1184 K
10–18 10–19 10–20 10–21
Tc = 1043 K
Dv, Dd and Dgb (m2 s–1)
10–11
Dv
10–22 10–23
0.8
1.0
1.2 1.4 T –1 (10–3 K–1)
1.6
1.8
7.4 Short-circuit diffusion coefficients in high purity α-iron (Iijima et al., 2005).
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cases, effective (total) diffusion coefficients, Deff, is estimated by the following equation: Deff = flDl + fgbDgb + fdDd
[7.7]
Here, fl , fgb and fd are the cross-sections of the lattice (volume), the grain boundary and the dislocation core diffusion, respectively. Practical creep conditions of common heat-resistant steels are located in the temperature range where the effects of short-circuit diffusion cannot be neglected. It is unfortunate that diffusion data in this temperature range are very limited in number and also the reliability of most existing data is not high. Short-circuit diffusion coefficients are greatly affected by the segregation of minor elements and also the character of grain boundaries. Practical materials usually contain many minor alloying elements and also some impurities. The character of grain boundaries in practical materials depends greatly on the history of their heat treatments. Hence, actual estimation of the short-circuit diffusion coefficients is not easy in heat-resistant steels.
7.4
Roles of atom/vacancy movement in creep
7.4.1
Deformation of matrix: dislocation climb controlled
In pure metals belonging to fcc, bcc and hcp lattice systems, dislocations can move rather easily on their glide planes, but find it difficult to get out of these planes. To continue straining under a given stress, dislocations must be eliminated from the glide plane by cross glide in the case of screw dislocations or climbing out of the glide plane in the case of edge dislocations. Cross gliding can occur with increasing stress level, but in climbing in edge dislocations (motion of jogs) a net change in number of atoms around the dislocation is necessary. To move jogs while keeping mass balance, vacancies must be created (atoms must go away) or disappear (atoms must come in) at the jogs. Therefore, formation and migration of vacancies (atom flow in the inverted direction) are an essential stage in order to continue creep deformation. In creep of most pure metals, the thermally activated rate-controlling step is believed to be diffusion of vacancies (Sherby and Weertman, 1979). In solid solution alloys, usually easiness of dislocation glide becomes less than that in pure metals. However, in some alloys, dislocation climbing becomes significantly more difficult and diffusion of vacancies (jog movement) remains as the rate-controlling step. This type of behaviour can be seen in some alloys of low stacking-fault energy. A typical example has been reported in creep of α-Cu–Al solid solution alloys (Hasegawa et al., 1972). A similar situation may be seen in some austenitic steels, but not in ferritic steels in which the stacking-fault energy is reasonably high.
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7.4.2
249
Deformation of matrix: dislocation glide controlled
In some solid solution alloys, solute atoms gather around edge dislocations and make a solute atmosphere. A typical cause of such segregation (clustering) is the elastic strain around edge dislocations and this type of segregation is called a Cottrell(-type) atmosphere. Once this type of atmosphere is generated, edge dislocations receive extra resistance to move even on the original glide plane. Sometimes the glide velocity is retarded to the level of the climbing velocity. In this case the motion of edge dislocations is controlled by the movement of solute atmosphere being formed around the dislocations. Therefore, in many alloys the creep rate is controlled by the diffusion of solute atoms rather than vacancies. This situation can easily be observed in alloys with solutes that have a large size misfit. An example can be seen in α-Fe–Mo alloys (Oikawa et al., 1980). In practical heat-resistant steels, especially in ferritic steels, most alloying elements have this type of strengthening effect. Chromium is an exceptional solute, because the size difference of Cr (versus α-Fe) is very small.
7.4.3
Influence through microstructure change
During creep deformation of complex alloys such as steels, significant changes in microstructure occur. At the very first stage of creep, a significant change in dislocation configuration occurs, followed by slow changes, not only in dislocation structures, but also in the kind, shape, size and amount of precipitates. Diffusion of the solute plays an important role in the formation of a new phase (precipitates). In this case the solute atoms gather to make precipitates. This process is a de-homogenization process and D˜ (parallel process) is suitable for analysing the process. If a strengthening solute element is consumed from the matrix to make a different phase as precipitates, the concentration of the solute causing drag resistance decreases and the material is weakened by the precipitate formation. When the third (minor) element affects the diffusion of main strengthening alloying elements, this minor element may act as a decelerator (or accelerator) of the weakening of the material. A typical example can be seen in the addition of Re in α-Fe–15Cr–5W alloy (Kunieda et al., 2006). In the early stage of precipitation, the size of the newly precipitated phase is very small but the number of precipitates is very large. Precipitate strengthening may occur at this stage, and some of the weakening by solute consumption may be cancelled. In a later stage of precipitation, the size of the precipitates increases and the number of precipitates decrease. In this precipitate coarsening stage, the
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diffusion of precipitate constituent atoms, D˜ , must be taken into account. In this case, the average concentration of the solute is unchanged, but still the strength of the material decreases because of weakening of the pinning force of precipitates on dislocations.
7.4.4
Failure
Correlation between creep rupture (time to failure) and diffusion is not well established. Failure is a more large-scale and complicated phenomenon than the straining of matrix and many factors other than simple diffusion must be taken into account in the analyses. At the very beginning stage of failure in creep of a specimen, the movement of vacancies may have a role in clustering vacancies and formation of embryos or nuclei of microcracks. The role of diffusion, however, may not be significant in the overall weakening of the matrix itself, because creep failure is a complicated phenomenon. Oxidation/ cracking, resulting stress concentration and so on, come to play more essential roles. The diffusion of oxygen into the matrix, especially through grain boundaries, diffusion of oxygen and solute atoms into/within oxide layers and diffusion behaviour in the coating layers are important processes to be considered.
7.5
Influence of some factors on creep through their effects on diffusion
7.5.1
Magnetic transformation
The ferromagnetic state retards diffusion, because of its spin ordering. The Curie temperature, TC, of ferritic iron is 1043 K so that most practical creep conditions of ferritic steels lie in the ferromagnetic state and diffusion is affected by this ferromagnetism (see Fig. 7.5). In pure iron, the effect of the ferromagnetic state on (self-)diffusion has been well studied. The effect can be quantitatively expressed as Dmag = Dnon-mag × exp( –αQ s 2/RT)
[7.8]
Here, s is the spontaneous magnetization relative to the value at 0 K and α a factor for the energy increase in the ferromagnetic state. The value s was measured by Potter (1934) and Crangle and Goodman (1971) for iron and s2 is shown as a function of temperature in Fig. 7.6. The value of α in α-iron has been reported for several elements (see Section 7.6). The magnitude of the effect depends greatly on the solute element and α is well correlated with the change in magnetization of the first and second nearest neighbours of the diffusing atom in iron (Nitta and Iijima, 2005). Cobalt diffusion in α-iron is an exceptional case, in which the ferromagnetic effect can be recognized up to about 100 K over TC.
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T (K) –14
10
1173 1073
973
873
773 Fe Co Cr Nb Mo
10–16
Tα–γ
10–15
10–18
Tc
D (m2 s–1)
10–17
10–19
Mo
10–20
Nb
10–21 10–22 10–23 0.8
0.9
1.0
1.1 1.2 T –1/(10–3 K)
1.3
1.4
7.5 Self-diffusion and impurity diffusion coefficients in α-iron around TC (Iijima et al., 1988; Nitta and Iijima, 2005).
1.0
0.8
s2
0.6
0.4
0.2
0.0 500
600
700
800 T (K)
900
1000
1100
7.6 Square of the spontaneous magnetization (relative to the value at 0K) of iron as a function of temperature.
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Creep rate is also affected by the ferromagnetic transformation through its effect on diffusion. An obvious change in the temperature dependence of creep rate is observed near TC in α-iron (e.g. Karashima et al., 1972) and many α-iron base alloys (e.g. Karashima et al., 1968b). The temperature dependence of creep rate becomes larger in the lower temperature range than that in the higher temperature range, that is, the (apparent) activation energy of creep is larger in the temperature range below TC, similar to the case of diffusion. In creep of α-Fe–Co alloys, a discrepancy between the effect of ferromagnetism on creep and TC of α-iron has been observed (Karashima et al., 1968b) as in the case of Co diffusion in α-iron. In some alloys such as the α-Fe–Mo system, the effect of ferromagnetism on creep can be observed up to 50–100 K higher than TC, although no such a discrepancy is reported in diffusion. A similar discrepancy between diffusion and creep in the effect of ferromagnetic transformation is also reported in γ-Fe–Ni high alloys (Karashima et al., 1968a).
7.5.2
Grain boundaries
It is well known that grain boundaries act as a strengthening factor in low temperature deformation, because of their resistance (barriers) to dislocation motion. In high temperature deformation, in contrast, grain boundaries act as sources and sinks of vacancies and also render a path of rapid diffusion for atoms. As a result, grain boundaries act as a weakening factor in high temperature deformation. Grain boundary diffusion coefficients depend greatly on the character of grain boundaries. They are largest in random boundaries, decreasing with a decrease in the randomness of boundaries, to being smallest in twin boundaries (∑3 coincidence boundary). The effects of grain boundaries on creep depend greatly on the character of grain boundaries as in the case of grain boundary diffusion. A significant effect can be recognized in random boundaries, but the effect becomes smaller as the degree of the randomness decreases. The smallest (almost nil) effect is expected in twin boundaries. In most experiments on grain boundary diffusion, Dgb was derived from penetration a small distance from the surface. Under this condition, diffusion through grain boundaries of relatively less random types can be determined. Recently, careful experiments were done on diffusion through grain boundaries of more random types. Values of Dgb determined in these experiments are significantly higher and the temperature dependence is lower than those reported hitherto. An example is shown in Fig. 7.7 for * Dgb in high purity α-iron. Similar results have been reported for grain boundary diffusion of Cr in austenitic stainless steels (Mizouchi et al., 2004). Therefore, a great care must be exercised in experiments on short-circuit
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10
1100
–16
900 800
T (k) 700
600
253
500
Iijima (2005) Bokstein (1959) Leymonie (1960) Borisov (1964) James (1965) Bernardini (1981) Hänsel (1985)
10–17 10–18 10–19
δDgb (m3 s–1)
10–20 10–21 10–22 10–23
10–26
10–27 0.8
Tc = 1043 K
10–25
Tα–γ = 1184 K
10–24
1.0
1.2
1.4
1.6
1.8
2.0
T –1 (10–3 K–1)
7.7 Grain boundary diffusion coefficients in α-iron showing the effect of penetration depth (Iijima, 2005).
diffusion to realize the types of short-circuit diffusion actually observed in the experiments. Speaking generally, the order of relative values in three kinds of diffusion coefficients is Dl << Dd < Dgb (see Fig. 7.4 for α-iron). The order of these diffusion coefficients is not likely to change with temperature, but their ratios depend greatly on temperature, increasing with decreasing temperature. This effect of temperature is significant near TC.
7.5.3
Segregation of minor elements
Care must be exercised because short-circuit diffusion coefficients are significantly sensitive to the segregation of impurity atoms and/or minor elements at the dislocation core and/or grain boundaries. Grain boundary diffusion coefficients depend greatly on the segregation of other elements at grain boundaries. The segregation usually retards the grain boundary diffusion, although the absolute value of Dgb is still quite high when comparing diffusion in the matrix.
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A similar situation is expected in the diffusion along dislocation cores, Dd. In Fig. 7.8 the effect of carbon content in impurity levels on diffusion along dislocation cores in α-iron is shown as a typical example. The presence of 16 ppm C causes a decrease in Dd of an order of magnitude or more. Practical creep conditions of heat-resistant steels may be located in an intermediate temperature range, where grain boundaries can act as a strengthening factor and also as a weakening factor. Heat-resistant steels contain many impurity (minor) elements other than the (intended) alloying elements. Therefore, actual grain boundaries may have quite complicated effects on diffusion and so also on creep strength.
7.5.4
Stacking faults
In metals and alloys with a low stacking-fault energy, γsf, a perfect dislocation may separate into two partial (imperfect) dislocations, and the separation distance between these two partial dislocations depends inversely on γsf. 1073 973 873 10–8
T (K) 773 673
573
Shima et al. (0.5 mass ppm C) Mehrer and Lübbehuesen (16 mass ppm C)
10–10
10–14
10–16
Tc = 1043 K
10–18
Tα–γ = 1183 K
Dd / (m2 s–1)
10–12
10–20
Dv
10–22 9
10
11
12 13 14 15 T –1 (10–4 K–1)
16 17
18
7.8 Dislocation (core) diffusion in α-iron showing the effect of carbon content (Shima et al., 2002; Mehrer and Lübbehusen, 1989).
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Between these two partial dislocations, a stacking-fault layer is usually formed. In solid solution alloys, the activity coefficient of the solute element in the stacking fault may be different from that of the (perfect) matrix and the solute element segregates in the stacking fault. When the partial dislocations move (on the glide plane), segregated atoms must move with the partial dislocations. Therefore, stacking faults may cause extra resistance to gliding of the (perfect) dislocation. The glide velocity of dissociated dislocations can thus be retarded. In some alloy systems, γsf decreases significantly with increasing solute concentration. In this type of alloy, the solute segregates easily in the stacking faults. The segregation of the solute may accelerate the decrease of γsf, which may result in more increase in creep strength. Even in a perfect dislocation of a pure screw type, at least one partial dislocation has some degree of edge component, so that the segregation of the solute (Cottrell atmosphere type) occurs and this solute atmosphere may raise resistance to the glide of the dislocation. The effect, however, of the solute segregation in the stacking fault and around a partial dislocation (having some edge component) on the glide velocity is removed by a pronounced retardation of the climb process in these low γsf materials. Therefore, in low γsf alloys dissociated dislocations glide viscously with the solute atmosphere, but the rate-determining step is still the climb motion, similar to that in pure metals. Some austenitic steels have intermediate/low γsf values and the glide of dislocations in these steels may be of a viscous type, but the rate-controlling step is likely to be the climbing (jog movement). Hence, the diffusion that needs to be considered in creep of austenitic steels is not that of solute atoms but of the vacancies.
7.6
Diffusion data in iron and in some iron-base alloys
7.6.1
Diffusion in α-iron and in bcc alloys
Diffusion parameters for self-diffusion and impurity diffusion in α-iron are listed in Table 7.1 and the temperature dependence of diffusion coefficients is shown in Fig. 7.9. References for data in the table are presented in a concise format for simplicity, that is, the second and subsequent authors are omitted and only the initial page number is cited. Most data were obtained using radioactive tracers. In cases where radioisotopes are difficult to use, diffusion data at very dilute concentrations of the solute are listed, which are indicated by (ID). No influence of the magnetization can be seen in the diffusion of interstitial elements, such as carbon, nitrogen and oxygen. The effect of the magnetization
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Table 7.1 Self-diffusion and impurity diffusion in α-iron Diffusant
Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
Reference
C N O Al (ID) 7 Be 57 Co 51 Cr 64 Cu 59 Fe 54 Mn (ferro) 54 Mn (para) 99 Mo 95 Nb 63 Ni 32 P (ferro) 32 P (para) Si (ID) 44 Ti 48 V 181 W
238–1168 244–1146 1023–1123 1173–1373 1073–1773 859–1173 885–1174 1063–1175 766–1148 973–1033 1073–1173 833–1163 823–1163 788–1160 932–1017 1078–1153 1100–1173 948–1174 1058–1172 833–1173
2.0 ×10–6 1.4 ×10–6 3.8 ×10–7 5.2 ×10–4 1.7 ×10–3 2.8 ×10–4 3.7 ×10–3 4.2 ×10–4 2.8 ×10–4 1.5 ×10–4 3.5 ×10–5 1.5 ×10–2 1.4 ×10–1 4.2 ×10–3 1.4 ×10 2.9 ×10–2 1.7 ×10–4 2.1 ×10–1 1.2 ×10–2 1.5 ×10–2
83.9 79.1 92.1 246 228 251 (α 267 (α 244 251 (α 234 220 283 (α 300 (α 268 (α 332 271 229 293 (α 274 287 (α
1 1 2 3 4 5 6 7 8 9 9 10 11 12 13 13 14 15 16 17
= 0.23) = 0.133) = 0.156)
= 0.074) = 0.061) = 0.060)
= 0.079) = 0.086)
1. Weller M (1996), Nichtmetalle in Metallen ’96, DGM Informations-gesellschaft mbH, Verlag, Oberursel, Germany 2. Takagi J (1986), Z Metallkd, 77, 6. 3. Nishida K (1971), Trans Jpn Inst Metals, 12, 310. 4. Grigorev GV (1968), Fiz Metallov Metalloved, 26, 946. 5. Iijima Y (1993), Mater Trans JIM, 34, 20. 6. Lee, CG (1990), Mater Trans JIM, 31, 255. 7. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 27. 8. Iijima Y (1988), Acta Metall, 36, 2811. 9. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267. 10. Nitta H (2002), Acta Mater, 50, 4117. 11. Oono N (2003), Mater Trans, 44, 2078. 12. Cermák J (1989), Z Metallkd, 80, 213. 13. Matsuyama T (1983), Trans Jpn Inst Metals, 24, 587. 14. Bergner D (1989), Defect Diffusion Forum, 66, 1407. 15. Klugist P (1995), Phys Stat Sol (a), 148, 413. 16. Geise J (1987), Z Metallkd, 78, 291. 17. Takemoto S (2006), Phil Mag, 87, 1619.
is reported for many substitutional elements. For some elements the value of * can be estimated for a wide α in Equation [7.8] is determined and Dimp range of temperature. In cases of manganese and phosphorus, diffusion parameters, D0 and Q, are reported as two sets for the paramagnetic (higher temperature) and the ferromagnetic (lower temperature) ranges, respectively. Tracer diffusion of constituent elements in iron-base bcc alloys are listed in Table 7.2. Most data are obtained in the paramagnetic α-phase. In some
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10–12 10–13 Si Be
10–14 Al Cu 10–15 Mn
P
D (m2 s–1)
10–16
V Ti Mn
10–17
P
–18
10
Cr
10–19
Mo Nb
10–20 Co 10–21
Ni W Fe
10–22 10–23 8
9
10 11 T –1 (10–4 K–1)
12
13
7.9 Self-diffusion and impurity(tracer) dffusion coefficients in α-iron.
alloys diffusion parameters have been determined over a wide range of temperature. Some are represented by the factor α in Equation [7.8]. Others are represented by two sets of parameters for the para- and ferro-magnetic ranges, respectively. Interdiffusion data in iron-base bcc alloys are listed in Table 7.3. Only a few systematic studies have been reported on interdiffusion and all data are those in the paramagnetic region.
7.6.2 Diffusion in γ-iron and fcc alloys Diffusion parameters for self-diffusion and impurity diffusion data in γ-iron are listed in Table 7.4 and the temperature dependence of diffusion coefficient is shown in Fig. 7.10. The mark (ID) has the same meaning as in Table 7.1. In the case of oxygen, diffusion data were obtained from the observation of microstructure changes and are indicated by (RD). Reliable diffusion data in the γ-phase alloys have been reported for limited systems only. Tracer diffusion values for constituent elements in binary and ternary alloys are listed in Table 7.5. There are some data for alloys of a Fe– Cr–Ni ternary system, which is the basic system of austenitic heat-resistant steels.
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Table 7.2 Tracer diffusion in bcc alloys Composition (atom%)
Diffusant Temperature (K)
D0 (m2 s–1) Q (kJ mol–1)
Reference
Fe–6 Al Fe–10Al (para) Fe–10Al (ferro) Fe–6.8Co(para) Fe–6.8Co(ferro) Fe–0.87Cr Fe–1.43Cr Fe–3.09Cr Fe–5.05Cr Fe–13Cr Fe–16Cr Fe–19Cr Fe–9.13Cr(ferro) Fe–15.2Cr(ferro) Fe–15.2Cr(para) Fe–19.7Cr(ferro) Fe–19.7Cr(para) Fe–0.4Mo Fe–0.4Mo Fe–1.5Mo Fe–1.5Mo Fe–1.48Si Fe–1.87Si Fe–2Ti Fe–2Ti Fe–4Ti Fe–6Ti Fe–1.8V Fe–2V Fe–5V Fe–5.3V
59
4.2 ×10–5 2 ×10–6 6 ×10–6 5.7 ×10–9 4.7 ×10–5 1.2 ×10–4 2.8 ×10–4 6.7 ×10–4 8.5 ×10–5 6.4 ×10–5 1.9 ×10–5 1.8 ×10–5 9.3 ×10–4 1.3 ×10–4 2.7 ×10–5 6.5 ×10–5 1.8 ×10–5 4.6 ×10–4 2.0 ×10–2 5.7 ×10–4 2.7 ×10–2 1.0 ×10–4 7.7 ×10–3 2.8 ×10–4 5.6 ×10–5 2.7 ×10–5 4.0 ×10–5 1.4 ×10–4 3.9 ×10–4 3.0 ×10–4 1.9 ×10–4
1 1 1 2 2 3 3 3 3 4 4 4 5 5 5 5 5 6 6 6 6 7 7 8 4 4 4 8 9 9 8
1. 2. 3. 4. 5. 6. 7. 8. 7.
Fe Fe 59 Fe 60 Co 60 Co 59 Fe 59 Fe 59 Fe 59 Fe 51 Cr 51 Cr 51 Cr 59 Fe 59 Fe 59 Fe 59 Fe 59 Fe 59 Fe 99 Mo 59 Fe 99 Mo 59 Fe 59 Fe 59 Fe 59 Fe 59 Fe 59 Fe 59 Fe 48 V 48 V 59 Fe 59
1088–1478 1156–1450 765–942 1153–1193 903–1073 1040–1173 1040–1173 1040–1173 1073–1173 1073–1673 1073–1673 1073–1673 848–999 868–950 999–1050 848–919 963–1098 823–1173 834–1143 823–1173 838–1120 1063–1373 1063–1373 1173–1473 1273–1673 1273–1673 1273–1673 1173–1773 1273–1723 1273–1723 1173–1466
198 184 196 146 187 241 249 256 237 232 218 217 231 227 216 217 208 253(α 285(α 253(α 290(α 276 276 276 216 205 209 237 241 239 240
= = = =
0.153) 0.090) 0.151) 0.076)
Raghunathan VS (1981), Phil Mag A, 43, 427. Hirano K (1970), Trans Jpn Inst Metals, 13, 96. ∨ Ku cera J (1974), Acta Metall, 22, 135. Bowen AW (1970), Met Trans, 1, 2705. Ray SP (1968), Acta Metall, 16, 981. Nitta H (2006), Acta Mater, 54, 2833. Treheux D (1981), Acta Metall, 29, 931. Lai DYF (1967), US Atomic Energy Commission Report–50314. Bowen AW (1970), Met Trans, 1, 2767.
Interdiffusion values in fcc binary alloys are listed in Table 7.6. Very limited data have been reported on interdiffusion in the γ-phase. Concentration dependences of interdiffusion parameters are reported in Fe–Co, Fe–Mn and Fe–Ni systems.
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10–9 O
10–10
C
10–11 N
D (m2 s–1)
10–12
P Be
10–13
W Ti
10–14 10
Mo Ni Fe
–15
Co
Cu Nb V Mn Cr
10–16 10–17 6
7 8 T –1 /(10–4 K–1)
9
7.10 Self-diffusion and impurity(tracer) diffusion coefficients in γ-iron.
Table 7.3 Interdiffusion in bcc alloys Composition (at %)
Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
Reference
Fe–9Al Fe–17Al Fe–3Co Fe–5Co Fe–2.5Mo Fe–4.1Mo Fe–5.0Mo Fe–0.47Si Fe–0.94Si Fe–1.87Si Fe–2.81Si Fe–3.74Si
1193–1483 1193–1483 1123–1168 1123–1168 1073–1573 1073–1573 1073–1573 1175–1708 1175–1708 1175–1708 1175–1708 1175–1708
2.7 ×10–5 3.6 ×10–7 3.0 ×10–1 4.1 3.6 ×10–4 4.2 ×10–4 4.0 ×10–4 7.8 ×10–5 7.7 ×10–5 9.6 ×10–5 1.1 ×10–4 1.3 ×10–4
188 142 318 343 257 259 262 221 220 220 219 220
1 1 2 2 3 3 3 4 4 4 4 4
1. 2. 3. 4.
Hishinuma A (1968), J Jpn Inst Metals, 32, 516. Hirano K (1977), Trans Iron Steel Inst Jpn, 17, 194. Nohara K (1977), Tetsu-to-Hagane, 63, 926. Borg RJ (1970), J Appl Phys, 41, 5193.
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Table 7.4 Self-diffusion and impurity diffusion in γ-iron Diffusant
Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
C (ID) N O (RD) 7 Be 60 Co 51 Cr 64 Cu 59 Fe 54 Mn Mo (ID) 95 Nb 63 Ni 32 P Ti (ID) 48 V W (ID)
1123–1578 1173–1623 1223–1373 1373–1623 1411–1613 1173–1618 1378–1641 1444–1634 1193–1553 1323–1633 1210–1604 1426–1560 1553–1623 1348–1498 1210–1607 1258–1578
2.3 ×10–5 9.1 ×10–5 1.3 ×10–4 3.3 ×10–5 1.3 ×10–4 1.7 ×10–5 4.3 ×10–5 4.9 ×10–5 1.6 ×10–5 3.6 ×10–6 8.3 ×10–5 1.1 ×10–4 2.8 ×10–3 1.5 ×10–5 6.2 ×10–5 5.1 ×10–5
148 169 166 256 305 264 280 284 262 240 267 297 292 251 274 272
1. 2. 3. 4. 5. 6. 7. 8. 7. 10. 11. 12. 13. 14. 15. 16.
Reference 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Agren J (1986), Scr Metall, 20, 1507. Grieveson P (1964), Trans Met Soc AIME, 230, 407. Takagi J (1986), Met Trans, 17A, 221. Grigorev GV (1986), Fiz Met Metalloved, 25, 836. Suzuoka T (1961), Trans Jpn Inst Metals, 2, 176. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 22. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 27. Heumann Th (1968), J Phys Chem Solids, 29, 1613. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267. Oikawa H (1983), Tech Reports Tohoku Univ, 48, 36. Geise J (1985), Z Metallkd, 76, 622. Hanatake Y (1978), Trans Jpn Inst Metals, 19, 667. Seibel G (1963), Compt Rend Acad Sci, C256, 4661. Moll SH (1959), Trans TMS-AIME, 215, 613. Geise J (1987), Z Metallkd, 78, 291. Ruzickova J (1981), Kov Mater, 19, 3.
7.7
Concluding remarks
7.7.1
Practical application of diffusion data
When the thermally activated step controlling creep phenomena is the climb process as in the case of pure metals, vacancy diffusion coefficients should be used in (theoretical) analyses. This will also be the case in some solid solution alloys of low/intermediate γsf , where dislocations glide viscously but the climb velocity is affected more significantly with the addition of solute elements, because of a wide stacking fault which is not easy to climb. In heat-resistant steels with an austenite base, the diffusion of vacancies should be noted.
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261
In alloys of high γsf, viscous gliding of (edge) dislocations controls the creep straining. In this case, the diffusion coefficients that should be used in (theoretical) analyses of creep are those of the solute elements. In heatresistant steels with a ferrite base, diffusion of the solute elements should be noted. It is unfortunate that not all suitable data are available and we have to use our second choice, estimated values in practice. When self-(tracer) diffusion coefficients of the constituent elements are already available as a function of solute concentration D ′ of vacancies, D˜ ′ of the solutes can be derived for a given composition. This situation, however, occurs in very limited alloy Table 7.5 Tracer diffusion in fcc alloys Composition (at%)
Diffusant Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
Reference
Fe–6.8Co Fe–2Cr Fe–6Cr Fe–9.13Cr Fe–1.04Mn Fe–1.04Mn Fe–2.03Mn Fe–2.03Mn Fe–2.97Mn Fe–2.97Mn Fe–4.90Mn Fe–4.90Mn Fe–14.9Ni Fe–14.9Ni Fe–29.7Ni Fe–29.7Ni Fe–15Cr–20Ni Fe–15Cr–20Ni Fe–15Cr–20Ni Fe–15Cr–20Ni–1.4Si Fe–15Cr–20Ni–1.4Si Fe–15Cr–20Ni–1.4Si Fe–15Cr–45Ni Fe–15Cr–45Ni Fe–15Cr–45Ni Fe–17Cr–12Ni Fe–17Cr–12Ni Fe–18Cr–8Ni Fe–18Cr–8Ni Fe–20Cr–25Ni/Nb Fe–20Cr–25Ni/Nb Fe–22Cr–45Ni Fe–22Cr–45Ni Fe–22Cr–45Ni
60
1.1 ×10–3 3.2 ×10–4 1.2 ×10–4 1.2 ×10–5 9.0 ×10–6 5.5 ×10–6 1.1 ×10–5 2.0 ×10–6 5.8 ×10–6 9.6 ×10–7 6.6 ×10–6 1.7 ×10–6 2.1 ×10–4 1.9 ×10–4 1.0 ×10–3 2.4 ×10–4 5.3 ×10–4 8.3 ×10–4 1.5 ×10–4 5.1 ×10–4 7.1 ×10–4 4.8 ×10–4 2.1 ×10–4 4.0 ×10–4 1.8 ×10–4 3.6 ×10–5 1.3 ×10–5 5.8 ×10–5 8.0 ×10–6 1.7 ×10–4 2.1 ×10–5 1.5 ×10–4 4.1 ×10–4 1.1 ×10–4
326 245 237 237 265 250 263 235 256 222 255 229 286 289 306 292 308 309 300 303 303 310 288 293 293 279 264 281 245 284 248 286 295 291
1 2 2 3 4 5 4 5 4 5 4 5 6 6 6 6 7 7 7 7 7 7 7 7 7 8 8 9 10 11 12 7 7 7
Co Cr 51 Cr 59 Fe 59 Fe 54 Mn 59 Fe 54 Mn 59 Fe 54 Mn 59 Fe 54 Mn 59 Fe 63 Ni 59 Fe 63 Ni 59 Fe 51 Cr 57 Ni 59 Fe 51 Cr 57 Ni 59 Fe 51 Cr 57 Ni 59 Fe 51 Cr 59 Fe 51 Cr 59 Fe 54 Mn 59 Fe 51 Cr 57 Ni 51
1283–1583 1073–1673 1073–1673 1173–1313 1263–1513 983–1573 1263–1513 983–1573 1263–1513 983–1573 1263–1513 983–1573 1258–1578 1258–1578 1258–1578 1258–1578 1236–1673 1236–1673 1236 –1673 1236–1673 1236–1673 1236–1673 1236–1673 1236–1673 1236–1673 873–1570 849–1568 1173–1473 923–1123 1029–1560 823–1523 1236–1673 1236–1673 1236–1673
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Table 7.5 Continued Composition (at%)
Diffusant Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
Reference
SUS316 SUS316 SUS316 SUS316 SUS316 SUS316
59
1.2 ×10–6 6.3 ×10–6 4.1 ×10–5 1.8 ×10–3 1.0 ×10–3 1.7 ×10–4
229 243 260 296 184 143
13 14 15 16 16 16
1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16.
stainless stainless stainless stainless stainless stainless
steel steel steel steel steel steel
Fe Cr 54 Mn 60 Co 95 Zr 99 Mo
1178–1483 1023–1473 1023–1473 1242–1423 1200–1490 1178–1497
51
Hirano K (1972), Trans Jpn Inst Metals, 13, 96. Bowen AW (1970), Met Trans, 1, 2705. Ray SP (1970), Trans Ind Inst Metals, 23, 77. Nohara K (1973), J Jpn Inst Metals, 37, 51. Nohara K (1971), Trans Iron Steel Inst Jpn, 11, 1267. Milliom B (1981), Mater Sci Eng, 50, 43. Rothman SJ (1980), J Phys F, 10, 383. Perkins RA (1973), Met Trans, 4, 2535. Linnenbom V (1955), J Appl Phys, 26, 932. Stawstorm C (1969), J Iron Steel Inst, 207, 77. Smith AF (1969), Met Sci, 3, 93. Smith AF (1975), Z Metallkd, 66, 692. Patil RV (1982), Met Sci, 16, 387. Smith AF (1975), Met Sci, 9, 375. Smith AF (1975), Met Sci, 9, 181. Patil RV (1980), Met Sci, 14, 525.
Table 7.6 Interdiffusion in fcc alloys Composition (atom%)
Temperature (K)
D0 (m2 s–1)
Q (kJ mol–1)
Reference
Fe–5Co Fe–10Co Fe–15Co Fe–3Cr Fe–5Mn Fe–10Mn Fe–15Mn Fe–10Ni Fe–20Ni
1323–1573 1323–1573 1323–1573 1173–1473 1123–1573 1123–1573 1123–1573 1273–1561 1273–1561
4.5 ×10–3 2.1 ×10–3 1.1 ×10–3 1.2 ×10–7 7.2 ×10–6 1.8 ×10–6 3.0 ×10–5 5.3 ×10–4 8.9 ×10–4
329 319 314 219 393 264 269 318 317
1 1 1 2 3 3 3 4 4
1. 2. 3. 4.
Hirano K (1977), Trans Iron Steel Inst Jpn, 17, 194. Davin A (1963), Rev Metall, 60, 275. Nohara K (1973), J Jpn Inst Metals, 37, 51. Goldstein JI (1965), Trans Met Soc AIME, 233, 691.
systems. When interdiffusion coefficients D˜ at/around the composition in question have already been determined, then this D˜ can be used as a rough estimation of D˜ ′ for the solute. In dilute solid solution alloys, this alternative usage may give a reasonable value.
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Without having the self-(tracer) diffusion coefficients of the solute and solvent at the composition in question, it is not easy to estimate diffusion coefficients of vacancies. In many cases only D of the impurity, Dimp, and self-(tracer) diffusion coefficients, D*, in the pure (base) metal are known experimentally, but no other diffusion data. In these cases we inevitably use D* and Dimp as rough alternatives for D ′ and D˜ ′ in the alloy.
7.7.2
Diffusion data searching
Diffusion data up to the end of the 1980s were well compiled in the chapters of a Landort–Börnstein book (Mehrer H (ed), 1990). It contains data for impurity diffusion (LeClaire A D and Neumann G 1990), self-diffusion in binary alloys (Bakker, 1990), chemical diffusion in binary alloys (Murch and Bruff 1990), diffusion in ternary alloys (Dayananda, 1990) and boundary diffusion (Kaur and Gust, 1990). Diffusion data reported thereafter can be found in the review journal Defect and Diffusion Forum published in ZürichUetikon by Trans Tech Publications.
7.8
References
Bakker H (1990), ‘Self-diffusion in homogeneous binary alloys and intermetallic phases’, in Diffusion in Solid Metals and Alloys Mehrer H (ed,), Landort-Börnstein, New Series, Group 3, Springer Verlag, Berlin, Volume 26, 213–278. Crangle J and Goodman G M (1971), ‘The magnetization of pure iron and nickel’, Proc Roy Soc London, Ser A, 321, 477–491. Dayananda M A (1990), ‘Diffusion in ternary alloys’, in Diffusion in Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3, Springer Verlag, Berlin, Volume 26, 372–436. Hasegawa T, Ikeuchi Y and Karashima S (1972), ‘Internal stress and dislocation structure during sigmoidal transient creep of a copper-16at.%aluminium alloy’, Met Sci J, 6, 78–82. Iijima Y (2005), ‘Diffusion in high purity iron: influence of magnetic transformation on diffusion’, J Phase Equil Diff, 26, 466–471. Iijima Y, Kimura K and Hirano K (1988), ‘Self-diffusion and isotope effect in α-iron’, Acta Metall, 36, 2811–2820. Iijima Y, Nitta H, Nakamura R, Takasawa K, Inoue A, Takemoto S and Yamazaki Y (2005), ‘Precise measurement of low diffusion coefficients using radioactive tracers’, J Jpn Inst Met, 69, 321–331. Karashima S, Motomiya T and Oikawa H (1968a), ‘High temperature creep of nickel– iron alloys”, Technol Rept Tohoku Univ, 33, 193–206. Karashima S, Oikawa H and Watanabe T (1968b), ‘The effect of ferromagnetism upon creep deformation of alpha-iron and its solid solution alloys’, Trans Metall Soc AIME, 242, 1703–1708. Karashima S, Iikubo T and Oikawa H (1972), ‘On the high-temperature creep behaviour and substructures in alpha-iron single crystals’, Trans Jpn Inst Met, 13, 176–181. Kaur I and Gust W (1990), ‘Grain and interphase boundary diffusion’, in Diffusion in
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Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3, Springer Verlag, Berlin, Volume 26, 630–716. Kunieda T, Yamashita K, Murata Y, Koyama T and Morinaga M (2006), ‘Effect of Re addition on W diffusivity in Fe–Cr alloys’, Mater Trans, 47, 2106–2108; erratum 47, 2888. LeClaire A D and Neumann G (1990), ‘Diffusion of impurities in solid metallic elements’, in Diffusion in Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3, Springer Verlag, Berlin, Volume 26, 85–212. Lüthy H, Miller A K and Sherby O D (1980), ‘The stress and temperature dependences of steady-state flow at intermediate temperature for pure polycrystalline aluminum’, Acta Metall, 28, 169–178. Mehrer H (ed) (1990), ‘Diffusion in Solid Metals and Alloys’, Landort-Börnstein, New Series, Group 3, Springer Verlag. Berlin, Vol. 26, p. 747. Mehrer H and Lübbehusen M (1989), ‘Self-diffusion along dislocations and in the lattice of alpha-iron’, Defect Diffus Forum, 66–69, 591–604. Mizouchi M, Yamazaki Y, Iijima Y and Arioka K (2004), ‘Low temperature grain boundary diffusion of chromium in SUS316 and 316L stainless steels’, Mater Trans, 45, 2945– 2950. Monma K, Suto H and Oikawa H (1964), ‘Relation between high-temperature creep and diffusion in alloys’, J Jpn Inst Met, 28, 387–400. Mukherjee A K, Bird J E and Dorn J E (1969), ‘Experimental correlations for hightemperature creep’, Trans Amer Soc Met, 62, 155–177. Murch G E and Bruff C M (1990), ‘Chemical diffusion in inhomogeneous binary alloys’, in Diffusion in Solid Metals and Alloys, Mehrer H (ed.), Landort-Börnstein, New Series, Group 3, Springer Verlag, Berlin, Volume 26, 279–371. Nitta H and Iijima Y (2005), ‘Influence of magnetization change on solute diffusion in iron’, Philos Mag Lett, 85, 543–548. Oikawa H (1982), ‘Lattice self-diffusion in solid iron: a critical review’, Technol Rept Tohoku Univ, 47, 67–77. Oikawa H, Saeki M and Karashima S (1980), ‘Steady-state creep of Fe-4.1 at%Mo alloy at high temperatures’, Trans Jpn Inst Met, 21, 309–318. Potter H (1934), ‘The magneto-caloric effect and other magnetic phenomena in iron’, Proc Roy Soc London, Ser A, 146, 362–387. Sherby O D and Burke P M (1967), ‘Mechanical behaviour of crystalline solids at elevated temperature’, Progress in Materials Science, 13, 325. Sherby O D and Weertman J (1979), ‘Diffusion-controlled dislocation creep: A defense’, Acta Metall, 27, 387–400. Sherby O D, Orr R L and Dorn J E (1954), ‘Creep correlation of metals at elevated temperatures’, Trans Amer Inst Min Metall Engr, 200, 71–80. Shima Y, Ishikawa Y, Nitta H, Yamazaki Y, Miura K, Isshiki M and Iijima Y (2002), ‘Selfdiffusion along dislocations in ultra high purity iron’, Mater Trans JIM, 43, 173–177.
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8 Fundamental aspects of creep deformation and deformation mechanism map K. M A R U YA M A , Tohoku University, Japan
8.1
Introduction
This chapter deals with the phenomenology of creep deformation. Mechanical properties such as yield stress and ultimate tensile strength are determined by tensile tests under constant strain rate. Deformation modes are different below (elastic) and above (plastic) the yield stress. On the other hand, creep tests under constant stress (or load) provide the creep properties necessary for structural materials for high temperature use. Creep deformation rate varies during a test from a high speed just after loading to a low speed around the minimum creep rate. It should be understood what is going on in the course of creep deformation. This chapter provides suggestions on how to understand creep behavior of engineering materials based on the stress– strain response of materials. Special attention is also paid to stress-accelerated creep tests for evaluating long term creep behavior. The athermal yield stress is a key point for appropriate evaluation of creep properties under service conditions. The athermal yield stress is the critical stress above which instantaneous plastic deformation takes place during loading. The difference in creep deformation behavior above and below the athermal yield stress will be examined. Creep tests are carried out over wide ranges of stress and temperature, and creep deformation rate varies widely from 10–8 to 10 h–1. Several creep deformation modes appear depending on the creep test conditions. The deformation mechanism map is another subject of this chapter. It can provide useful information, for example which creep mode is operative under a given creep condition.
8.2
Stress–strain response of materials
Suppose a specimen is subjected to a tensile test at a strain rate, ε˙ , at room temperature. The specimen first deforms elastically and stress, σ, increases linearly with increasing strain, ε, as shown in Fig. 8.1. Then plastic deformation 265 WPNL2204
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True stress
Stress, σ
σu Engineering stress σy
Plastic Elastic
Strain, ε
8.1 Difference between a true stress–strain curve and an engineering stress–strain curve at room temperature.
starts at the yield stress, σy, and the stress–strain response deviates from the elastic line. Strain hardening proceeds in the course of plastic deformation and the flow stress, σ, increases with increasing strain usually due to increase in dislocation density and formation of dislocation substructures. In scientific considerations the true stress σΤ should be used. The true stress is defined by: σT = L/A
[8.1]
where L is the applied load and A is the current cross-section area at strain of ε. Engineering stress, σE, defined by the following equation is often more practical and we often use it in engineering expressions: σE = L/A0 = σT/(1 + ε)
[8.2]
where A0 is the initial cross-section area and ε in this equation is the engineering strain (elongation/initial length of specimen). As Fig. 8.1 shows, the engineering stress is lower than the true stress in tensile tests and their difference increases with increasing strain. The engineering stress–strain curve has a maximum caused by the reduction of cross-section area and the subsequent necking, despite the absence of such a peak on the true stress–strain curve. The maximum stress is called ultimate tensile strength (UTS). In the ASME code, allowable stress is given by two-thirds of yield stress when the twothirds of yield stress is lower than one-quarter of UTS, but given by onequarter of UTS when the one-quarter of UTS is lower. UTS together with yield stress play an important role when determining the allowable stress. In the case of brittle materials, UTS corresponds to fracture stress. However,
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the UTS of a ductile material gives the maximum load attainable during a tensile test and may not have sufficient scientific importance. Stress–strain curves tested at elevated temperatures show a slightly different shape from those tested at room temperature, at which diffusion of atoms can be neglected. At elevated temperature dislocations can climb and annihilate themselves with other dislocations. Therefore, after a sufficient amount of deformation, dislocation density reaches a stationary value determined by the dynamic balance between multiplication and annihilation of dislocations. The steady state dislocation substructure results in a steady state flow under a constant true stress as depicted in Fig. 8.2. In the case of pure metals and solid solution alloys with a low density of initial dislocations, dislocation density increases to a steady state value. The increase results in the true stress–strain curve of strain hardening type drawn in Fig. 8.2. High Cr ferritic steels often contain a high density of dislocations introduced during martensitic transformation. Plastic deformation assists in annihilation of dislocations in such steels. The recovery of dislocation substructures results in a true stress– strain curve of a strain softening type. We sometimes use the steady state flow stress–strain rate relation as a substitute for the steady state creep rate– stress relation in creep.
8.3
Temperature and strain rate dependence of yield stress
The yield stress of steel increases with decreasing temperature at room temperature and below (see Fig. 8.3). In plastic deformation at low temperature
σy
Steady state
Elastic
True stress, σ
line
Strain softening
Strain hardening σy
Strain, ε
8.2 Strain hardening type and strain softening type stress–strain curves at elevated temperature.
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Normalized yield stress, σy /E
268
. Low ε
Athermal yield stress
σa
. High ε 0K Temperature, T
8.3 Temperature dependence of yield stress σy normalized with Young’s modulus E, and effects of strain rate on the curve. σa is the athermal yield stress independent of temperature and strain rate.
the Peierls barrier (a short range obstacle) is the main obstacle to dislocation motion. Dislocations overcome the obstacle with the assistance of thermal vibration of atoms and applied stress. Since the thermal energy decreases with decreasing temperature, the yield stress increases at low temperature. At intermediate temperatures ranging from 100–450°C for steel, the thermal vibration energy becomes large enough for dislocations to overcome the Peierls barrier and the main obstacle to dislocation motion changes to another long range obstacle, such as other dislocations and particles. Since these obstacles are too large for dislocations to overcome by thermal energy, dislocations can overcome the obstacles with the aid of applied stress only. Therefore, the yield stress at intermediate temperature is essentially independent of temperature and strain rate. Such a yield stress is called the athermal yield stress σa. The yield stress is related to dislocation density ρ and interparticle spacing λ by the following equations: σa = α M G b ρ
[8.3]
σa = β M G b/λ
[8.4]
where α is a constant of about 0.4, M is the Taylor factor (= 3), G is the shear modulus, b is the length of Burgers vector and β is a constant of about 0.8. The yield stresses show weak dependence on testing temperature owing to the temperature dependence of G. They are truly independent of temperature when normalized by an elastic constant such as G or Young’s modulus E. Dynamic strain aging caused by carbon and nitrogen atoms may introduce a peak on the yield stress–temperature curve at intermediate temperature, but the peak is insignificant when 0.02 or 0.2% proof stress is used for the yield stress.
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At elevated temperature, diffusion of atoms and vacancies assists dislocations to pass through the athermal obstacles, resulting in another temperature and strain rate dependent plastic deformation, namely creep. This temperature range is called the creep regime. Yield stress at high temperature is expressed as: σy = E ( ε˙ / ε˙ 0 )1/ n exp(QD/nRT)
[8.5]
where ε˙ is the strain rate, ε˙ 0 is a material constant, n is the stress exponent which is usually greater than 3, QD is the activation energy for lattice diffusion, R is the universal gas constant and T is the absolute temperature. Therefore, the yield stress of high temperature deformation decreases with increasing temperature or decreasing strain rate. In low temperature deformation, thermal energy assists dislocations to overcome the Peierls barrier. The assistance of thermal energy increases with decreasing strain rate ε˙ , resulting in a decrease in yield stress. However, σa is the lower limit of yield stress at low temperature, since applied stress below σa does not allow dislocations to pass through the long range obstacles. At high temperature, yield stress increases with increasing strain rate, as expected from Equation [8.5]. However the yield stress cannot exceed the athermal yield stress σa, since above σa dislocations can pass through the athermal obstacles without the aid of diffusion. The athermal yield stress is an important value even in high temperature deformation, since it is the upper limit of yield stress.
8.4
Deformation upon loading of creep test
During loading of creep tests at elevated temperature specimens deform at a high strain rate, suggesting athermal deformation during the loading. To examine this expectation the amount of deformation εi upon loading of 2.25Cr– 1Mo steel is plotted in Fig. 8.41 as a function of creep stress σ normalized by Young’s modulus E. The steel was normalized at 930°C for 20 min and then tempered at 720°C for 2 h. The εi versus σ/E relationships tested at various temperatures overlap with each other, confirming the athermal nature of the deformation during the loading. The εi versus σ/E curve resembles a typical stress–strain curve of tensile test. Linear response, namely elastic deformation occurs at low stress. Above a critical stress (σ/E = 6–7 × 10–4) the εi versus σ/E curve deviates upwards from the elastic line. One can easily obtain 0.02% proof stress (micro yield stress) and 0.2% proof stress (macro yield stress) from the εi versus σ/E curve. The proof stresses normalized by E are independent of testing temperature, confirming the prediction of Fig. 8.3. The micro yield stress is two-thirds of the macro yield stress in this example. The micro yield stress is used as the athermal yield stress σa in this chapter.
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Elongation upon loading, εi
10–1
–2
10
450°C 475°C 500°C 525°C 550°C 600°C 650°C
10–3
σa
0.2%
% 0.02
2/3
10–4 0.1
0.2
0.4 0.6 1 Normalized stress, σ/10–3 E
Elastic
2
4
8.4 Elongation upon loading in creep test for 2.25Cr–1Mo steel as a function of stress normalized with Young’s modulus E.
8.5
Creep behavior below and above athermal yield stress
Suppose a creep stress of σ1 (σ1 < σa) is applied to a specimen, then the specimen first deforms elastically to the strain corresponding to point A as shown in Fig. 8.5 (a). The deformation stops at this point at room temperature. Atoms and vacancies can move a sufficient distance above 40% of the melting point of the material; 450°C in the case of iron. At these elevated temperatures, dislocations can climb over particles on their slip planes with the aid of diffusion. Recovery of dislocation substructure proceeds during high temperature exposure, resulting in a decrease in σa given by Equation [8.3]. Diffusional flow of atoms themselves can bring about creep deformation (diffusion creep). Because of these reasons, time-dependent plastic deformation, namely creep, occurs below the athermal yield stress as depicted in Fig. 8.5 (b). The strain–time curve is called the creep curve. Slope of the creep curve (dε/dt) is called creep rate. On the creep curve, creep rate first decreases owing to strain hardening (primary creep) and reaches a stationary value (secondary creep). Degradation of microstructures and accumulation of creep damage, such as cavities and micro cracks, proceed during creep, resulting in creep acceleration to rupture. This stage is called tertiary creep. Since the stationary creep rate in the secondary creep stage is minimum during the whole creep process, the creep rate is often called the minimum creep rate and used as a measure of creep deformation resistance. The minimum (or steady state) creep rate ε˙ m is expressed as: ε˙ m = ε˙ 0 (σ/E)n exp(–Qc/RT)
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Fundamental aspects of creep deformation (a)
271
(b)
Strain, ε
σ2
σ1 . εm
C
Tertiary
Secondary A σ1
B σ2 σa Stress, σ
Primary
tr Time, t
8.5 (a) Athermal stress–strain curve and (b) creep curves at stresses σ1 (below athermal yield stress σa) and σ2 (above σa).
where ε˙ 0 is a material constant, σ is the creep stress and Qc is the activation energy for creep. The values of n and Qc are closely related to creep mechanism as will be explained in Section 8.7. Practical heat-resistant steels are strengthened by various obstacles, such as dislocation substructures, precipitates and so on. Degradation of these obstacles starts at a very early stage of creep. In such materials the minimum creep rate is attained by dynamic balance between strain hardening and microstructural degradation.2,3 At applied stress σ2 higher than σa, the specimen deforms elastically to point B and then plastically to point C within a short period of time. The specimen is strain hardened during the plastic deformation by introducing dislocations. The athermal plastic deformation stops at this point and no further deformation continues at room temperature. Since recovery of dislocation substructure proceeds at elevated temperature, another timedependent plastic deformation (creep) continues at elevated temperature. Creep behavior above σa is similar to that below σa. However, it should be noted that above σa, dislocations are introduced during fast plastic deformation upon loading and the dislocation substructure at the beginning of creep tests above σa is different from that below σa.
8.6
Change in creep behavior at athermal yield stress σa
Let us take a look at how creep characteristics change at σa in a particlestrengthened material. Orowan stress, σOr, of a material with average particle spacing λ is given by Equation [8.4]. As depicted in Fig. 8.6, the particles
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(c) Region H: σ > σmy Orowan loop
(b) Region M: σmy > σ > σOr Particle
(a) Region L: σOr > σ
8.6 Changes in dislocation substructures in particle-strengthened material with decreasing creep stress. (a) Above macro yield stress σmy, (b) between σmy and Orowan stress σOr and (c) below σOr.
alone can sustain the applied stress σ below σOr without introducing additional dislocations. Above σOr dislocations pass through the particles leaving an Orowan loop at each particle. The Orowan loops bring about strain hardening without appreciable increase in dislocation density. This type of strain hardening occurs below the macro-yield stress σmy. Above σmy the particles with dislocation loops cannot sustain the applied stress and dislocations travel long distance upon loading, introducing additional dislocations into the material. These drawings explain how the dislocation substructure just after loading above σmy is different from that below σOr.4 Experimental results for an aluminum alloy containing particles (2.2 vol% Al6Mn) are given in Fig. 8.7;4 minimum creep rates ε˙ m are plotted against creep stress σ normalized by shear modulus G. The positions of σΟr (0.02% proof stress) and σmy (at 0.08% plastic strain) measured experimentally by tensile test are indicated in the figure. There are three regions with a different stress exponent n: the value of n is low (n = 9) in the high stress region H (σ > σmy), takes a medium value (n = 14) in the intermediate stress region M (σOr < σ < σmy), and is high (n = 26) in the low stress region L below σOr.
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10–1
Minimum creep rate (s–1)
Al–2.2vol% Al6 Mn M
10–3
H 10–5 L σmy
–7
10
350°C 300°C 250°C
σOr 10–9
0.6
0.8
1
2
3
Normalized stress, σ/10–3 G
8.7 Minimum creep rate of a particle-strengthened alloy as a function of creep stress σ normalized with shear modulus G. σOr and σmy are Orowan (micro-yield) stress and macro-yield stress, respectively.
8
σ/10–3 G
Al–2.2vol% Al6Mn
1.0 (M) 0.98 (M) 0.95 (L) 0.81 (L)
Dislocation density (1013 m–2)
6
4
2
1 0.8 (a) 250°C 0.6 0
0.1
σ/10–3 G 2.0 (H) 1.3 (M/H) 1.0 (L) 0.2
(b) 350°C 0
0.1
0.2
Strain
8.8 Changes in dislocation density during creep tests. The solid symbol is the density before testing. Stress levels H, M and L correspond to the three regions indicated in Fig. 8.7.
The ε˙ m versus σ curves exhibit threshold like behavior typical of particle strength materials. Figure 8.84 shows the evolution processes of dislocation substructures in the alloy during loading and the subsequent creep deformation. In Fig. 8.8
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the stress levels L, M and H correspond to the three regions indicated in Fig. 8.7. The dislocation density versus strain relationships accord very well with the prediction from Fig. 8.6. In region L (σ < σOr) the dislocation density is kept close to the original value during the whole creep test, since no additional dislocation is needed to sustain the applied stress. Bowing out of dislocations between particles can explain the slight increase in dislocation density from the value before loading. At σ > σmy (Region H) the presence of particles is not sufficient to sustain the applied stress and a high density of dislocations is introduced into the alloy during loading. The recovery of dislocations proceeds towards a steady state value during the primary creep. A similar recovery process of dislocation substructures has often been reported in creep of pure metals. In region M (σOr < σ < σmy) an obvious increase in dislocation density is absent upon loading since the alloy can sustain applied stress only by introducing Orowan loops around particles. The dislocation density increases up to a steady state level during the primary creep, suggesting that the hardening by Orowan loops is replaced by hardening by dislocation substructures. The stationary densities of dislocations in the Al–Mn alloy are plotted against applied stress in Fig. 8.9.4 The boundaries between regions L and M and between M and H are indicated in the figure. The boundaries L/M shift slightly among the three testing temperatures. The following well known relationship holds between the dislocation density ρ and creep stress σ normalized by shear modulus G:
8
[8.7]
Al – 2.2vol% Al6Mn
6 4
250
L/M 350 300
Stationary dislocation density (1013 m–2)
ρ = (σ/α MGb)2
2 250°C 300°C 350°C
M/H 1 0.8 0.6
0.6
0.8
1
2
3
Normalized stress, σ/10–3 G
8.9 Stationary dislocation density as a function of creep stress σ normalized with shear modulus G. L/M and M/H indicate the boundaries between regions L and M and between regions M and H, respectively.
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The equation is essentially the same as Equation [8.3]. Below the boundary L/M, on the other hand, the dislocation density is independent of creep stress and kept at a value close to the original dislocation density before creep loading. As demonstrated in this section creep behavior above macro-yield stress should be different from that below micro-yield stress. The minimum creep rates for the 2.25Cr–1Mo steel used in Fig. 8.4 are plotted in Fig. 8.10 as a function of creep stress normalized by Young’s modulus E. 0.02% (σ0.02) and 0.2% (σ0.2) proof stresses are indicated in the figure. The stress exponent decreases with decreasing stress and then discontinuously jumps up to 15 below σ0.02. This threshold like-behavior of the 2.25Cr–1Mo steel resembles Fig. 8.7 which was plotted for a particle-strengthened aluminum alloy. Engineering materials are used below the micro-yield stress and creep behavior changes at the micro-yield stress. Their creep properties under service conditions should be evaluated from short-term creep data obtained below the microyield stress.
8.7
Deformation mechanism maps
A material deforms by several mechanisms at elevated temperatures depending on its creep conditions, namely stress and temperature. Representatives of the mechanisms are diffusion creep controlled by volume diffusion (Nabarro– Herring creep) or grain boundary diffusion (Coble creep) and dislocation creep controlled by volume diffusion (high temperature power law creep) or 10–2
Minimum creep rate (h–1)
2.25Cr–1Mo steel
10–4
650°C 600°C 550°C 525°C
10–6
σ0.02
500°C 475°C 450°C
σ0.2
–8
10
0.1
0.2
0.4
0.6
1
2
4
Normalized stress, σ/10–3 E
8.10 Stress dependence of minimum creep rate of normalized and tempered 2.25Cr–1Mo steel.
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by pipe diffusion (low temperature power law creep). The creep rates of all the mechanisms are represented by the following general equation:
ε˙ = ε˙ 0 (σ/G)nd pD
[8.8]
where ε˙ 0 is a material constant characteristic of the mechanism and material, d is the grain size, p is the grain size exponent and D is the diffusion coefficient relevant to the mechanism. The values of n, p and D are typical of each creep mechanism and are listed in Table 8.1. Dl, Dp and Dgb are the diffusion coefficient of lattice, dislocation pipe and grain boundary diffusion, respectively. The four representative creep mechanisms are independent of each other and the creep strain produced by each mechanism contributes additively to the total creep strain. Therefore, the mechanism giving the highest value of ε˙ dominates creep deformation at a given stress and temperature. This assumption predicts that a mechanism with a high stress exponent n operates at high stress. The diffusion creep mechanisms take the stress exponent of n = 1 and are operative at low stress, while a low temperature power law creep of n = nd + 2 dominates at high stress, where nd is the stress exponent of high temperature power law creep and nd = 3 – 5. A high temperature power low creep of n = nd is operative at intermediate stress. The Coble creep (controlled by grain boundary diffusion) and low temperature power law creep (by dislocation pipe diffusion) are rate controlled by short circuit diffusion. Activation energies for such short circuit diffusion paths are lower and almost half that of lattice diffusion. Therefore, the Coble creep and the low temperatures power law creep with low activation energy appear at lower temperatures than the Nabarro–Herring creep and the high temperature power law creep. The aforementioned consideration provides the deformation mechanism map drawn in Fig. 8.11.5 The diffusion creep mechanisms (Coble creep and Nabarro–Herring creep) appear in the lowest stress range and the dislocation creep mechanisms (high temperature and low temperature power law creep) in the intermediate stress range. A dislocation glide mechanism without the aid of diffusion takes over the role of plastic deformation above the athermal yield stress. The critical stress for the onset of the dislocation glide mechanism Table 8.1 Stress exponent n, grain size exponent p and diffusion coefficient D in Equation [8.8] for each creep mechanism Deformation mechanism
n
p
D
Dislocation creep Low temperature power law creep High temperature power law creep
5–7 3–5
0 0
Dp Dl
Diffusion creep Coble creep Nabarro–Herring creep
1 1
3 2
Dgb Dl
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10–1
277
Ideal strength Dislocation glide
10–3
σa Low temperature power law creep
High temperature power law creep
10–5 Nabarro–Herring creep
Coble creep 10–7
0
0.2
0.4 0.6 0.8 Normalized temperature, T/Tm
1
8.11 Schematic drawing of deformation mechanism map. Tm is the melting temperature of the material under consideration.
corresponds to the athermal yield stress σa, and is independent of temperature and strain rate as mentioned in Sections 8.3 and 8.4. Nabarro–Herring creep and high temperatures power law creep operate at higher temperatures owing to their higher activation energy for creep. Creep rates of both dislocation creep mechanisms are independent of grain size, but those of the diffusion creep mechanisms increase with decreasing grain size. Therefore, the diffusion creep fields, especially the Coble creep field, expand, while the dislocation creep fields shrink with decreasing grain size. Deformation mechanism maps of various materials are available in a book by Frost and Ashby.5 By using this kind of map we can predict a relevant deformation mechanism of a material under any creep conditions. Correct evaluation of creep properties requires accelerated creep tests conducted in the same deformation mechanism field as the service conditions. Below the athermal yield stress, creep deformation by one of the four mechanisms starts after elastic deformation upon loading, whereas plastic deformation by the dislocation glide mechanism proceeds during loading above the athermal yield stress. The deformation mechanism map is correct in this sense. However, above the athermal yield stress dislocation creep begins after the athermal plastic deformations upon loading (see Fig. 8.10).1 The main obstacles for dislocation creep below the athermal yield stress are the inherent obstacles present in the material, such as solute atoms in solutiontreated materials, precipitates in usual engineering steels and dislocation substructure in tempered martensite steels. On the other hand, the main obstacles above the athermal yield stress are dislocation substructures introduced into specimens during loading, as demonstrated in Section 8.6. One should keep this fact in mind when using deformation mechanism maps. WPNL2204
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Creep-resistant steels
Concluding remarks
Plastic deformation at room temperature is essentially time independent and finishes within a short period of time. At elevated temperatures, thermally activated migration of atoms and vacancies, namely diffusion, occurs extensively and time-dependent plastic deformation (creep) continues for a long period of time until fracture. Several deformation mechanisms, such as diffusion creep and dislocation creep appear, depending on testing stress and temperature. Deformation mechanism maps are useful in predicting the deformation mechanism operative under the creep condition of interest. In the case of heat-resistant steel, the stress and temperature ranges of interest are 1 × 10–4E – 3 × 10–3E (E: Young’s modulus) and 0.4Tm – 0.55Tm (Tm: melting temperature). Under such conditions, the stress exponent for creep rate is usually greater than 3, suggesting that dislocation creep is the relevant deformation mechanism of engineered steels.1 Creep deformation occurs both above and below athermal yield stress. The athermal yield stress is the yield stress of a material at a high strain rate at a elevated temperature, for example during loading of the creep test. Below the athermal yield stress, creep (time-dependent plastic deformation) starts after elastic deformation upon loading. Above the stress, time independent plastic deformation (dislocation glide) occurs upon loading before timedependent plastic deformation (creep). Since plastic deformation upon loading alters the major obstacle to creep deformation from the inherent obstacles to dislocation substructures, creep deformation behavior is different above and below the athermal yield stress. Engineering materials are used below the athermal yield stress. Their creep properties should be evaluated from shortterm creep data obtained below their athermal yield stress.
8.9
References
1 Maruyama K, Sawada K, Koike J, Sato H and Yagi K, ‘Examination of deformation mechanism maps in 2.25Cr–1Mo steel by creep tests at strain rates of 10–11s–1 to 10– 6 –1 s ’, Mater Sci Eng, 1997, A224, 166–172. 2 Evans R W, Parker J D and Wilshire B, ‘An extrapolation procedure for long-term creep strain and creep life prediction with special reference to 1/2Cr1/2Mo1/4V ferritic steels’, in Recent Advances in Creep and Fracture of Engineering Materials and Structures, Wilshire B and Evans R W (eds), Pineridge Press, Swansea, 1982, 135–184. 3 Maruyama K, Tanaka C and Oikawa H, ‘Long-term creep curve prediction based on the modified θ projection concept’, Trans ASME, J Press Vess Technol, 1990, 112, 92– 97. 4 Maruyama K and Nakashima H, Materials Science for High Temperature Strength – Creep Theories and their Application to Engineering Materials, Uchida-Rokakuho, Tokyo, 1997. 5 Frost H J and Ashby M F, Deformation Mechanism Maps, Pergamon Press, Oxford, 1982.
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9 Strengthening mechanisms in steel for creep and creep rupture F. A B E, National Institute for Materials Science (NIMS), Japan
9.1
Introduction
Design stress of elevated-temperature components used under creep conditions is usually determined on the basis of a 100 000 h creep rupture strength at the operating temperature, and sometimes also a 200 000–300 000 h creep rupture strength. Therefore, creep-resistant steels must be reliable over long periods exceeding 100 000 h at elevated temperature and the characteristic parameters for the strength of creep-resistant steels are long-term creep and creep rupture strength. In this chapter, the strengthening mechanisms in creep-resistant steels are described with emphasis on long-term creep strength, mainly for tempered martensitic 9–12Cr steels.
9.2
Basic ways of strengthening steels at elevated temperature
The basic ways in which creep-resistant steels can be strengthened are by solute hardening, precipitation or dispersion hardening, dislocation hardening and boundary hardening.1–4 These should be helpful in examining the behavior of engineering creep-resistant steels at elevated temperature. It is possible to combine several strengthening mechanisms but it is often difficult to quantify each contribution to the overall creep strength.
9.2.1
Solid solution hardening
Substitutional solute atoms such as Mo and W, which have much larger atomic sizes than those of solvent iron, have been favored as effective solid solution strengtheners for both ferritic and austenitic creep-resistant steels. It should be noted that the contribution of solid solution hardening by Mo and W to the overall creep strength of engineering creep-resistant steels is practically superimposed on other strengthening mechanisms, for example precipitation hardening. As will be described later, the addition of Mo and W sometimes 279 WPNL2204
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causes precipitation of the Fe2(Mo,W) Laves phase and enhances fine distributions of M23C6 carbides during exposure at elevated temperature. Taking the Hume–Rothery size effect and large solid solubility in iron into account, Ir5 and Re,6 which are located in the lower part of periodic table, are also promising for effective solid solution hardening. In terms of interstitial solute atoms, it is well known that nitrogen is beneficial for long-term creep strength of engineering ferritic and austenitic creep-resistant steels through solid solution hardening as well as precipitation hardening by fine nitrides. Figure 9.1 shows the minimum creep rate of α-Fe and solid-solution αFe–Mo–W alloys (FE, MH, MWH, WH) with a ferrite matrix and precipitation hardened α-Fe–C–Mo–W–V–Nb alloys (PO, PSM, PSW) at 600°C.7 The precipitation hardened α-Fe–C–Mo–W–V–Nb alloys contained MX carbonitrides of (Nb, V)(C, N) which were 40 nm in size in the matrix. The minimum creep rate of solid solution α-Fe–Mo–W alloys is three orders of magnitude lower than that of α-Fe, indicating an effective solid solution hardening by Mo and W. On the other hand, the minimum creep rate of precipitation hardened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) is only one-half lower than that of precipitation hardened α-Fe–C–V–Nb alloys (PO). This indicates that the strengthening mechanisms in precipitation hardened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) come mainly from precipitation strengthening by fine MX. It should be noted that the contribution 1
Minimum creeprate (h–1)
10–1
10–2
600°C α – Fe
: FE : MH : MWH : WH : PO : PSM : PSW
Solid solution α – Fe–Mo–W
7
10–3 9 10
11
–4
10–5 10–6 20
α – Fe–V–Nb
40
α – Fe–Mo–W –V–Nb
60 100 Stress (MPa)
200
9.1 Minimum creep rate of α-Fe and solid-solution α-Fe–Mo–W alloys with a ferrite matrix and precipitation hardened α-Fe–C–Mo–W–V–Nb alloys at 600°C. MH: Fe–0.120Mo (mass%), MWH: Fe–0.34Mo–1.63W, WH: Fe–2.27W, PO: Fe–0.06C–0.2V–0.08Nb, PSM: Fe–0.06C–0.51Mo– 0.2V–0.08Nb, PSW: Fe–0.06C–0.15Mo–0.70W–0.2V–0.08Nb. FE is pure iron.
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of solid solution hardening by Mo and W to the overall creep strength of precipitation strengthened α-Fe–C–Mo–W–V–Nb alloys (PSM, PSW) is much smaller than that of solid solution Fe–Mo–W alloys (MH, MWH, WH). This suggests that an additive rule for solid solution hardening and precipitation hardening does not hold. Transmission electron micrograph (TEM) observations show that the creep deformation of the alloys shown in Fig. 9.1 is controlled by the mobility of dislocations within sub-grains and at sub-grain boundaries.
9.2.2
Precipitation or dispersion hardening
Precipitation or dispersion hardening is one of the important strengthening mechanisms in creep-resistant steels at elevated temperature. To achieve enough strengthening using this effect, engineering creep-resistant steels usually contain several kinds of precipitate particles in the matrix and at grain boundaries: carbonitrides such as M23C6, M6C, M7C3, MX and M2X, where M denotes the metallic elements, C are the carbon atoms and X are the carbon and nitrogen atoms, intermetallic compounds such as the Fe2(Mo,W) Laves phase, Fe7W6 µ-phase, χ-phase and so on, and a metallic phase such as Cu. In a special case of oxide dispersion strengthened (ODS) steels, fine particles of alloy oxides such as Y2O3 are dispersed in the matrix by mechanical alloying. A dispersion of fine precipitates stabilizes free dislocations in the matrix and sub-grain structure, which enhances dislocation hardening and sub-boundary hardening. Several mechanisms have been proposed for the threshold stress, corresponding to the stress needed for the dislocation to pass through precipitate particles, for example, in the Orowan mechanism, local climb mechanism, general climb mechanism and Srolovitz mechanism, see Fig. 9.2.8 The Orowan stress σor is given by: σor = 0.8MGb/λ
[9.1]
where M is the Taylor factor (= 3), G is the shear modulus, b is the Burgers vector and λ is the mean interparticle spacing.3 Typical values of the volume fraction, diameter and spacing of the major particles contained in tempered martensitic high Cr steels after tempering are listed in Table 9.1, together with the Orowan stress estimated from the values of interparticle spacing.3 The coarsening of fine precipitates of M23C6, MX and Fe2(W, Mo) Laves phase and the dissolution of fine MX to form massive precipitates of Z phase, which have been observed in 9–12Cr steels during creep, cause an increase in λ in Equation (9.1) and hence a decrease in Orowan stress over long periods of time.3,4 The coarsening and dissolution of fine precipitates sometimes takes place preferentially in the vicinity of grain boundaries during creep, which promotes the formation of localized weak zone and promotes
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Dislocation
Dislocation Particle
Particle
1
2
3
1
2
3
4
(b)
(a)
Dislocation
x
x
Dislocation (d)
(c)
9.2 Schematic drawings of a dislocation passing through particles. (a) Orowan mechanisms, (b) Srolovitz mechanism, (c) general climb mechanism and (d) local climb mechanism. Table 9.1 Volume fraction, diameter and spacing of each kind of precipitate in high Cr ferritic steel, together with Orowan stress estimated from the values of interparticle spacing Particle
Volume fraction V(%)
Diameter dp(nm)
Spacing σp(nm)
Orowan stress σor (MPa)
Fe2(W, Mo) M23C6 MX
1.5 2 0.2
70 50 20
410 260 320
95 150 120
localized creep deformation near grain boundaries.9,10 This results in premature creep rupture and is time and temperature dependent. The strengthening mechanisms caused by a dispersion of oxide particles were examined for a 13Cr–3W–0.5Ti–0.4Y2O3 ODS steel with ferrite matrix at 650°C, by comparing the threshold stress measured by a stress abruptly loading test (SAL test) with the calculated Orowan and void-hardening stresses, σor and σV.11 Figure 9.3(a) shows the relationship between the creep stress and strain upon loading for the ODS steel at 650°C, using the time elapsed after applying the stress as a parameter. The Orowan and void-hardening stresses are calculated to be 135–192 and 114–163 MPa, respectively, from the histogram for size distribution of Y2O3 particles in the steel (Fig. 9.3(b)). These values are also shown in Fig. 9.3(a). The threshold stress, caused by
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the dislocation passing through the oxide particles, just after loading (at t = 100 ms) is measured to be 175 MPa as shown by arrow A in Fig. 9.3(a), which agrees with the calculated Orowan stress. As the time elapses to 4 s, the threshold stress decreases to 150 MPa as shown by arrow B, which agrees with the calculated void-hardening stress. This suggests that the originating mechanisms of the threshold stress come from the Srolovitz mechanism in this steel. According to the Srolovitz mechanism,12 when the matrix-particle interface is incoherent, the normal traction of dislocation stress field on the particle surface is relaxed by interface sliding and volume diffusion and the dislocation is attracted to the particle, see Fig. 9.2(b). After 16
Strain (× 10–3)
14 12 10 Elastic line 8 6 4
t = 100 ms t=1s t=2s t=3s t=4s
A B
2
σv
σv
0 0
100
200 Stress (MPa)
300
400
(a) 50
Relative frequency (%)
Total number: 1243 40
30
20
10 0 0
1
2
3
4
5 6 7 Radius (nm)
8
9
10
(b)
9.3 (a) Relationship between creep stress and instantaneous strain at 650°C and (b) size distribution of Y2O3 particles in a 13Cr–3W–0.5Ti– 0.4Y2O3 ODS steel.
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the relaxation is completed, the particles are felt by dislocations as voids and the threshold stress should be equal to the void-hardening stresses. In Fig. 9.3(a), the time of 1–4 s required for the change from the Orowan stress to the void-hardening stress corresponds for the time necessary to the full relaxation. The Srolovitz mechanism was also confirmed for the threshold stress in high-temperature deformation of Al–1.5Be and Al–0.7Mn alloys containing incoherent precipitate particles.13
9.2.3
Dislocation hardening
Dislocation hardening given by: σρ = 0.5 MGb (ρf)1/2
[9.2]
where ρf is the free dislocation density in the matrix, is an important means of strengthening steel at ambient temperature. Tempered martensitic 9–12Cr steels usually contain a high density of dislocations even after tempering, usually in the range 1–10 × 1014 m–2 in the matrix.3 The density of dislocations produced by martensitic transformation during cooling after austenitization can be controlled by changing the tempering temperature. Tempering is usually carried out at low temperatures of 700°C or lower for turbine steels to ensure enough tensile strength at ambient temperature by the dislocation hardening, whereas it is as high as 750–800°C for boiler applications. At elevated temperature, on the other hand, cold working enhances softening by promoting the recovery of excess dislocations and the recrystallization of deformed microstructure. A comparison of creep rupture strength of 12Cr– 1Mo–1W–VNb steel at 600°C and 650°C between the two tempering treatments at 750°C and 800°C shows that the low temperature tempering gives higher creep rupture strength than high temperature tempering for short times below about 15 000 h and 6000 h at 600°C and 650°C, respectively.14 However, the stress versus time to rupture curve is crossed over during long peiods of time by that of the steel subjected to high temperature tempering at 800°C. This is because excess dislocations accelerate recovery and recrystallization during creep with the aid of stress. The dislocation density after tempering at a higher temperature of 800°C is too low to promote recovery and recrystallization during creep. This effectively suppresses a rapid decrease in creep rupture strength. Also in austenitic steels, the creep strength of cold worked materials is typically higher for short times than that of solution annealed materials but it is reversed for long periods of time.15,16 These results indicate that the dislocation hardening is useful in the creep strength only for short times but that it is not useful for long-term creep strength at elevated temperature. Figure 9.4 shows the effect of cold rolling on the creep rate versus time curves and creep rate versus strain curves of a tempered martensitic 9Cr–
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Creep rate (h–1)
10–2
10–3
10–4
Standard QT QT = 20% CR QT + 40% CR 10–1
100
101 Time (h) (a)
102
103
Creep rate (h–1)
10–2
10–3
10–4
Standard QT QT + 20% CR QT + 40% CR 10–3
10–2 True strain (b)
10–1
9.4 Effect of cold rolling on (a) creep rate versus time and (b) creep rate versus strain curves for a tempered martensitic 9Cr–1W–0.1C steel at 600°C and 78 MPa.
1W–0.1C steel at 600°C and 78 MPa.17 V, Nb and nitrogen were not added to this steel and hence only M23C6 carbides were distributed along grain boundaries and lath boundaries after tempering. The decrease in creep rate with time and strain in the transient creep region is less pronounced with increasing cold rolling level. The onset of acceleration creep significantly shifts to shorter times with increasing cold rolling level. This results in a much higher minimum creep rate and hence a shorter time to rupture with
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increasing cold rolling level. The strain to reach a minimum creep rate increases with increasing cold rolling level. The cold rolling provides a high density of free dislocations, which accelerates the creep deformation at elevated temperature. The movement and annihilation of excess dislocations within the lath is considered to be the major process in the transient creep region. In austenitic steels, cold working is sometimes employed after solution annealing to accelerate the nucleation of fine carbides such as NbC at dislocations, resulting in a homogeneous distribution of a high density of fine NbC carbides in the matrix.18 This enhances the precipitation hardening.
9.2.4
Sub-boundary hardening
Tempered martensitic high Cr steels subjected to normalizing and tempering are usually observed to have a lath martensitic microstructure consisting of lath and block with a high density of dislocations and a dispersion of fine carbonitrides along the lath and block boundaries and in the matrix. The lath and block can be regarded as elongated sub-grains. The lath and block boundaries provide the sub-boundary hardening given by: σsg = 10Gb/λsg
[9.3]
where λsg is the short width of elongated subgrains.3 The subgrain width λsg, corresponding to the width of lath and block, is in the range 0.3–0.5 µm in martensitic high Cr steels after tempering. Using the values of G = 64 GPa at 650°C, b = 0.25 nm, λsg = 0.3–0.5 µm, we obtain σsg = 530–320 MPa, which are much larger than the Orowan stress in Table 9.1 for Fe2(W, Mo), M23C6 and MX. Therefore, the sub-boundary hardening gives an important means of strengthening the creep strength of tempered martensitic high Cr ferritic steels. Many people have considered that M23C6 carbides precipitate preferentially along the lath, block and prior austenite grain boundaries in tempered martensitic high Cr steels, while MX carbonitrides are distributed in the matrix. However, modern energy-filtered transmission electron microscopy (FE-TEM) studies have revealed that a large number of fine MX carbonitrides are distributed along the lath, block and prior austenite grain boundaries, as well as in the matrix after tempering,19 as shown in Fig. 9.5. The fine distributions of M23C6 carbides and MX carbonitrides along the lath and block boundaries stabilize these boundaries and exert a pinning force against the migration of lath and block boundaries and the coarsening of lath and block during creep. This suggests that the strengthening achieved by subboundary hardening is further enhanced by a dispersion of fine precipitates along sub-grain boundaries. The coarsening of lath and block with creep strain, which takes place mainly in the tertiary or acceleration creep region20,21 and causes an increase
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Bright field image
287
V map
500nm
9.5 Bright field image and elemental mapping of Cr and V for P92 in as tempered steel, by energy-filtered TEM.
in λsg in Equation (9.3), indicates the mobile nature of lath and block boundaries under stress. It is well known that polygon and sub-grain boundaries free from precipitates in pure metals and solid solution alloys are highly mobile under applied stress.1 The movement of lath and block boundaries can absorb or scavenge excess dislocations inside the lath and block. This corresponds to a dynamic recovery process, resulting in softening. Thus lath and block boundaries act as sink sites for the recovery of excess dislocations as well as act as barrier sites exerting back stress given by Equation (9.3) for the movement of dislocations inside the lath and block during creep. The role of lath and block boundaries in the strengthening mechanisms for creep and creep rupture is not yet fully understood.
9.3
Strengthening mechanisms in modern creepresistant steels
9.3.1
Bainitic low-Cr steels
The most recent advancement in bainitic low Cr steels is the improvement in the creep rupture strength of 2.25Cr and 3Cr steels.4 Figure 9.6 shows the extrapolated 100 000 h creep rupture strength for 2.25Cr–1Mo (T22) and 2.25Cr–1.6W–VNb (T23), 2.25Cr–1Mo–VTi (T24) and 3Cr–1.5W–0.75Mo– 0.25V without Ta (grade A) and with 0.1Ta (grade B) as a function of temperature, compared with martensitic 9Cr–1Mo–VNb steel (T91).22 The creep rupture strength of the grade B steel is higher than T23 for the entire test temperature range and also higher than T91 up to 615°C. The higher oxidation rate in air during creep testing of the grade B steel, owing to its 3%Cr at temperatures exceeding 615°C, as opposed to T91 which contains 9%Cr, is the reason for its lower creep rupture strength at high temperature. The grade A steel has higher creep rupture strength than T23 and T91 up to 600°C. The strengthening mechanisms in 2.25Cr–1.6W–VNb steel (T23) are found
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Creep-resistant steels 100 000 h creep rupture strength (MPa)
350 T22 (2.25Cr–1Mo) T23 (2.25Cr–1.6W–0.25V–0.05Nb) T24 (2.25Cr–1Mo–0.25V–0.07Ti) T91 (9Cr–1Mo–0.2V–0.05Nb) Grade A (3Cr–0.75Mo–1.5W –0.25V) Grade B (3Cr–0.75Mo–1.5W –0.25V–0.1Ta)
300 250 200 150 100 50 0 520
540
560
580 600 Temperature (°C)
620
640
660
9.6 100 000 h creep rupture strength as a function of temperature of T23, T24, Grade A and B of 3Cr–1.5W–0.75Mo–0.25V steel, compared with T22 and T91.
to be the combination of solid solution hardening due to W and precipitation hardening due to fine (V, Nb)C carbides in a fully bainitic matrix.23 Figure 9.7 shows the creep rupture data for 2.25Cr–1Mo–0.25V–0.05Nb steel (Mosteel) and 2.25Cr–1.6W–0.25V–0.05Nb steel (W-steel) at 650°C, comparing with those for a conventional 2.25Cr–1Mo steel. The creep rupture strength is increased by the addition of V and Nb and also by the substitution of W for Mo. M23C6 forms along prior austenite grain boundaries, while MX (M = V and Nb, X = C and N) forms along lath boundaries and in the bainitic matrix after tempering. M7C3 is occasionally observed along lath boundaries. In the crept specimens, no M23C6 and M7C3 are observed and instead blocky M6C is formed along prior austenite grain boundaries and fine MX remains along lath boundaries and in the bainitic matrix. The pronounced lath structure is kept in 2.25Cr–1Mo–0.25V–0.05Nb steel even after long term creep, suggesting the availability of sub-boundary hardening in addition to precipitation hardening for long periods of time. The substitution of W for Mo also constrains the precipitation of M6C during creep as shown in Fig. 9.8, which constrains the evolution of the bainitic microstructure.23 The precipitation of M6C loosens W and or Mo in solution, which is the key factor for the solid solution strengthening in this steel. The growth rate of M6C is 10 to 100 times slower in the W steel than in the Mo steel. It is also found that the MC carbides in Mo steel have already lost coherency with the matrix after heat treatment, while those in W steel have kept coherency even after long-term exposure. This suggests that W is superior to Mo in terms of precipitation hardening. In 2.25Cr–1Mo–
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Stress (MPa)
2.25Cr–1Mo–0.25V –0.05Nb 100 90 80 70
289
2.25Cr–1Mo
60 650°C
50 40 102
103 Time to rupture (h)
104
1
Fraction of MW precipitation (–)
Fraction of Mo precipitation (–)
9.7 Creep rupture data for 2.25Cr–1Mo–0.25V–0.05Nb and 2.25Cr– 1.6W–0.25V–0.05Nb steels at 650°C, compared with those for conventional 2.25Cr–1Mo steel.
0.8 923K 873K
0.6 0.4
823K 0.2 0 1
10
100 1000 104 Aging time (h) (a)
105
1 0.8 0.6 923K
873K
0.4 0.2 0 1
823K 10
100 1000 104 Aging time (h) (b)
105
9.8 Concentration of Mo and W in extracted residues of (a) 2.25Cr– 1Mo–0.25V–0.05Nb and (b) 2.25Cr–1.6W–0.25V–0.05Nb steels, showing the precipitation of Mo and W as M6C. The curves represent Johnson–Mehl–Avrami precipitation curves.
VTi (T24), the precipitation of fine TiC carbides is one of the strengthening mechanisms. Both 3Cr–3W–0.25V and 3Cr–3W–0.25V–0.1Ta steels are observed to have an acicular bainitic structure.24 The addition of 0.1Ta to 3Cr–3W– 0.25V steel substantially decreases the prior austenite grain size. Fine TaC particles are precipitated and dispersed within the grains and also along prior austenite grain boundaries. The carbides along prior austenite grain boundaries in the Ta-containing steel are smaller than those in the steel without Ta. This is one of the major reasons for the improvement of creep rupture strength of 3Cr–3W–0.25V steel by the addition of 0.1Ta.
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9.3.2
Creep-resistant steels
Tempered martensitic 9–12Cr steels
Figure 9.9 shows the creep rupture data for high strength martensitic 9–12Cr steels at 650°C.4 In this figure, two 9Cr–3W–3Co–0.2V–0.05Nb steels with 0.05N–0.002C (MARN: Martensitic 9Cr steel strengthened by MX nitrides)25,26 and with 0.08C–0.008N–0.014B (MARBN: Martensitic 9Cr steel strengthened by boron and MX nitrides)10,27 were developed by the National Institute for Materials Science (NIMS), 12Cr–2.6W–2.5Co–0.5Ni–0.2V–0.05Ni steel (NF12)28 and 12Cr–3W–3Co–0.2V–0.05Nb–0.1Ta–0.1Nd–0.05N steel (SAVE12)29 were developed by Japanese steelmaking companies as upgrade versions of P92 and P122, respectively, and oxide dispersion strengthened (ODS) 9Cr steel, 0.13C–9Cr–2W–0.2Ti-0.35Y2O3, with tempered martensitic microstructure30 was developed for fast breeder reactor cladding materials. The strengthening mechanisms in MARN steel come mainly from precipitation hardening by fine MX nitrides, which enhances the sub-boundary hardening. In order to achieve a dispersion of fine MX nitrides alone, which are thermally stable particles for prolonged periods of exposure at elevated 300
Stress (MPa)
200
100 80
60 1 10
102
103 Time to rupture (h)
104
105
9Cr–3W–3CO–VNb–0.05N–0.002C (MARN) 9Cr–3W–3Co–VNb–0.08C–0.008N–0.014B (MARBN) P92 (9Cr–0.5Mo–1.8W–VNb) T91 (9Cr–1Mo–VNb) ODS–9Cr (Ukai et al) NF 12 (12Cr–2.6W2.5Co–NVNb) SAVE 12 (12Cr–3W–3Co–VNbTaNdN)
9.9 Creep rupture data for 9Cr steels of MARN, MARBN, P92, T91, ODS-9Cr and 12Cr steels of NF12 and SAVE12 at 650°C.
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temperatures, it is crucial for 9Cr steels to reduce carbon content to very low amounts of less than 50 ppm, because the addition of carbon to a 9Cr steel causes the formation of a large amount of M23C6 carbides rich in Cr. The time to rupture significantly increases in the lower carbon region below 0.02%, where fine MX nitrides are dominant, Fig. 9.10. A large number of fine precipitates which have a size less than 10 nm are distributed not only in the matrix but also along lath, block, packet and prior austenite grain boundaries in the 0.002C steel after tempering, stabilizing the lath–block structure for long periods of time. The major component of the MX nitrides was identified as vanadium nitride. The strengthening mechanisms in MARBN steel come mainly from the combination of the boron effect and of precipitation hardening by fine MX nitrides, which enhances the sub-boundary hardening. The nitrogen content in the MARBN steel (0.008% N) is much lower than that in the MARN steel (0.05% N), because excess addition of boron and nitrogen causes the formation of massive boron nitrides during normalizing at high temperature.31 The formation of boron nitrides offsets the benefit given by boron and nitrogen. The carbon content for MARBN is increased in comparison with that for MARN. Addition of boron reduces the rate of Ostwald ripening of M23C6 Number of moles
0.04 M23C6 0.02 MX 0 0.00
0.05
0.10 (a)
0.15
0.20
12 000
Time to rupture (h)
9Cr–3W–3Co–0.2V–0.05Nb–0.05N
8000
4000
0 0
0.05 0.1 0.15 Carbon concentration (wt%) (b)
0.2
9.10 (a) Amount of M23C6 and MX after tempering and (b) time to rupture of 9Cr–3W–3Co–VNbN steel at 650°C and 140 MPa, as a function of carbon concentration.
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carbides in the vicinity of prior austenite grain boundaries during creep at 650°C. The M23C6 carbides are mainly distributed along prior austenite grain boundaries and along lath, block and packet boundaries inside the grain. Therefore, the fine distribution of M23C6 carbides and fine lath-block structure is maintained in the vicinity of prior austenite grain boundaries for long periods of time in the steel with about 100 ppm boron during creep. A fine distribution of M23C6 carbides is also observed in the steel without boron after tempering, but extensive coarsening of M23C6 and lath-block takes place during high-temperature exposure. With increasing boron content, the transient or primary creep region continues for longer periods of time and hence the onset of acceleration or tertiary creep retards, which results in a lower minimum creep rate and longer rupture time, as shown in Figure 9.11. The effect of boron relates to the stabilization of the fine lath-block structure in the vicinity of prior austenite grain boundaries for long periods of time through the stabilization of fine M23C6 carbides. Most of early studies on ODS steels concentrated on ODS ferritic steels with a ferrite matrix.30 The main problem in ODS ferritic steels that kept them from being used was the anisotropy in creep strength between the circumferential hoop and the longitudinal direction in tubes. One improvement in this anisotropy has been achieved by using martensitic transformation which produces an equi-axed grain structure. The ODS martensitic steel exhibits lower creep strength than the ODS steel with a ferrite matrix in the longitudinal direction but without anisotropy. TEM observations show fine distributions of Y–Ti oxide particles with a size of about 3 nm in the matrix of tempered martensite, together with Ti oxides of hundreds of nanometers and M23C6 carbides of several hundred nanometers in size. Oxide particles are introduced by complicated mechanical alloying. New attempts have also been demonstrated in using fine precipitates of FePd–L10 intermetallic compound in tempered martensitic 9Cr steel32,33 and fine intermetallic compounds in 15Cr steel with a ferrite matrix.34
9.3.3
Austenitic steels
Austenitic steels are usually solid solution hardened prior to service and are capable of significant precipitation hardening by fine carbonitrides and sometimes by fine intermetallic compounds during service at elevated temperature. The creep strength of austenitic steels for super heater boiler tubes has been enhanced from conventional 18Cr–8Ni steel to 20Cr–25Ni steel and then to high Cr and high Ni steels combined with the addition of W, which can be used for a long time at a steam temperature of 700°C. For example, the creep rupture strength at 700°C and 100 000 h is estimated to be 86 MPa and 90 MPa for 20Cr–25Ni–1.5Mo–NbTiN steel (KASUS310J2TB)35 and 23Cr–43Ni–6W–NbTiB steel (HR6W),36,37 respectively,
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10–2
Creep rate (h–1)
10–3
10–4
10–5 10–6
10–7 10–8 –1 10
0 ppmB 48 ppmB 92 ppmB 139ppmB
100
101
102 103 Time (h)
log(creep rate)
Onset of acceleration creep by microstructure recovery near G.B. Base steel
104
105
Increase in creep life
Boron-steel
Stabilization of M23C6 and lath near grain boundary
Decrease in minimum creep rate
log(time)
9.11 Effect of boron on creep rate versus time curves of 9Cr–3W– 3Co–VNb steel at 650°C and 80 MPa and schematics of a creep deformation mechanism.
which are higher than those for high-strength stainless steels such as Super304H and TP347HFG. A cast 19Cr-12.5Ni–4Mn–MoNbN steel (CF8C-Plus),18 recently developed by Oak Ridge National Laboratory, shows a creep rupture strength similar to KA-SUS310J2TB. The strengthening mechanisms in 23Cr–43Ni–6W–NbTiB steel (HR6W) come from the combination of solid solution hardening by W and precipitation hardening by fine M23C6 carbides, fine MX carbonitrides and fine Fe2W Laves phase.36,37 Figure 9.12 shows creep rupture data for 0.08C–23Cr– 43Ni–(5–7)W–0.2Nb–0.1Ti–0.003B steel (W steel) at 700, 750 and 800°C, compared with data for 0.08C–23Cr–43Ni–(3–5)Mo–0.2Nb–0.1Ti–0.003B
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Creep-resistant steels 200 3% Mo 5% Mo 5% W 7% W
Stress (MPa)
150
700°C × 105 h
100
80 21.5
22.0
22.5
23.0 23.5 24.0 T(20 + logt) × 10–3
24.5
25.0
9.12 Creep rupture data for austenitic 0.08C–23Cr–43Ni (5–7)W– 0.2Nb–0.1Ti–0.003B steel (W steel) and 0.08C–23Cr–43Ni–(3–5) Mo– 0.2Nb–0.1Ti–0.003B steel (Mo steel) at 700, 750 and 800°C.
steel (Mo steel). These steels were solution annealed at 1200°C for 1 h followed by air cooling. The creep rupture strength is improved by the substitution of W for Mo for long periods of time. The main precipitates in the Mo steel are a coarse σ phase at grain boundaries and M23C6 carbides and a coarse Fe2Mo Laves phase inside the grain, while those in the W steel are fine M23C6 carbides but no σ phase at grain boundaries and fine Fe2W Laves phase, fine M23C6 and fine MX inside the grain, as shown schematically in Fig. 9.13. TEM observations show that the fine precipitates of M23C6, MX and Fe2W Laves phase in the W steel serve as an effective dislocation barrier, Fig. 9.14. In the 20Cr–25Ni–1.5Mo–NbTiN steel (KA–SUS310J2TB) with 0.3Nb and 0.05Ti, three types of precipitate are observed in the crept specimens: massive particles of 0.4–0.5 µm, string-like Cr–Nb nitrides of less than 0.03 µm and fine granular and needle-like M23C6 carbides of less than 0.2 µm.35 The addition of a small amount of Nb and Ti is effective in depressing of precipitate coarsening. The cast austenitic steel 19Cr–12.5Ni–4Mn–MoNbN (CF8C-Plus) with 0.8Nb was developed as an upgraded version of 19Cr–10Ni–MoNbSi steel (CF8C).18 The alloy-design philosophy for CF8C-Plus is to eliminate all of the detrimental phases such as the σ phase, leaving only carbides for strengthening. After creep rupture testing of CF8C-Plus for 23 000 h at 850°C and 35 MPa, fine NbC carbides of much less than 50 nm in diameter are closely spaced and uniformly dispersed throughout the matrix. These fine NbC carbides nucleate on dislocations very early during creep and are one of the important strengthening mechanisms.
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M23C6 0.08C-23Cr-43Ni-5/7W –0.2Nb-0.1Ti-0.003B (a)
M23C6
295
Coarse laves (little)
σ 0.08C-23Cr-43Ni-3/5Mo –0.2Nb-0.1Ti-0.003B (b)
9.13 Schematic drawings of microstructure in (a) 0.08C–23Cr–43Ni– (5–7)W–0.2Nb–0.1Ti–0.003B and (b) 0.08C–23Cr–43Ni(3–5)Mo–0.2Nb– 0.1Ti-0.003B steels at 700, 750 and 800°C.
9.14 Fine precipitation of the Fe2W Laves phase as well as M23C6 and MX, showing effective obstacles to dislocation motion in 23Cr–43Ni– 7W–0.2Nb–0.1Ti–0.003B steel (HR6W), after creep rupture testing for 58 798.4 h at 700°C.
9.4
Loss of strengthening mechanisms in 9–12Cr steels during long time periods
In recent years, efforts have been made to clarify the mechanisms of creep strength loss in tempered martensitic 9–12Cr steels at 550°C and above during creep exposure. The loss of creep strength often takes the form of a sigmoidal inflection over long periods of time in creep rupture data. The proposed mechanisms relate mainly to a loss of precipitation hardening by fine carbonitrides and also to a loss of dislocation hardening during creep exposure. These accelerate microstructure evolution such as the coarsening of lath and block, resulting in a loss of sub-boundary hardening.4
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9.4.1
Creep-resistant steels
Precipitation of new phases and dissolution of fine carbonitrides
The precipitation of Z phase, M6X carbonitrides and Fe2(W, Mo) Laves phase during creep causes a loss of creep strength over long periods of time, because they consume existing fine precipitates. The Z phase is a complex nitride of the form Cr(Nb,V)N. Figure 9.15(a) shows the sigmoidal behavior of creep rupture data for 0.1C–10.84Cr–0.14Mo–2.63W–2.86Co–0.55Ni– 0.19V–0.06Nb–0.016N–0.019B steel (TAF 650) at 650°C, compared with the data for original TAF and 9Cr–1Mo–0.2V–0.05Nb steel (Mod.9Cr–1Mo, T91).38 The synergetic effect of Z phase precipitation and tungsten depletion of the solid solution by Fe2W Laves phase formation could be the reason for the sigmoidal shape of the creep strength curve of TAF 650 steel. Precipitation of the Z phase takes place after a long time at the service temperature and it forms large particles at the expense of the previously formed fine vanadium nitrides, which leaves a vanadium nitride free zone around the Z phase. This causes a drastic loss of creep strength. Figure 9.15(b) shows an example of a Z phase formed in a 12Cr steel NF12.39 It has been suggested that the precipitation of the Z phase may have the most detrimental effect on creep strength for tempered martensitic high Cr steels. 12Cr steels are more susceptible to precipitation of the Z phase than 9Cr steels because of the high content of Cr.39 A higher content of nitrogen also accelerates the precipitation of the Z phase.40 The coarsening of M 23C 6 carbides in 12CrMo(W)VNbN steels is accompanied by dissolution of fine MX carbonitrides owing to the precipitation of coarse M6X and or the Z phase.41 Increase in Ni content in 12CrMoV steel results in accelerated microstructure degradation with more rapid coarsening of M23C6, dissolution of MX, and precipitation of coarse M6X and Fe2Mo.42 The impurities Al and Ti cause the formation of AlN and TiN in creep400
TAF 650 TAF TAF 650 Mod. 9Cr-1Mo
Stress (MPa)
300 200
Z4 Z5
100 80
Z6
60
Z phase MX
40 100
1000 10 000 Time to fracture (h) (a)
100 000 0.2µm (b)
9.15 Comparison of stress versus time to rupture curves for various tempered martensitic 9–12Cr steels at 650°C.
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resistant steels during creep which degrades the creep strength at the expense of dissolved nitrogen and fine vanadium nitrides.43 Al and Ti are strong nitride-forming elements. Figure 9.16 shows the creep rupture data for the nine heats of tempered martensitic 12Cr–1Mo–1W–0.3V steel at 600°C, where the variation in impurities Al and N was 0.007–0.040Al and 0.015– 0.032N, respectively, between the different heats. The time to rupture simply decreases with decreasing available nitrogen concentration under low stress
Stress (MPa)
300
RAA RAB RAC RAD RAE RAF RAG RAH RAJ
100 80 60 40 101
102
Time to rupture (h)
105
103 104 Time to rupture (h) (a)
105
69 MPa 98 MPa 157 MPa 216 MPa
104
103
102 0
0.04 0.08 ∆ = N – Al – Ti (atom%) (b)
0.12
9.16 (a) Creep rupture data for nine heats of tempered martensitic 12Cr–1Mo–1W–0.3V steel with different Al, Ti and nitrogen contents at 600°C and (b) time to rupture as a function of available nitrogen concentration defined as nitrogen – Al – Ti (atom%).
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and long time conditions. The available nitrogen concentration is defined as the concentration of nitrogen free from AlN and TiN, and is given by the difference ∆ = nitrogen – Al – Ti (atom%). Higher Al content results in lower available nitrogen concentration.
9.4.2
Preferential recovery of microstructure near prior austenite grain boundaries
The loss of creep rupture strength in T91 steel is due to preferential recovery of the lath martensitic microstructure in the vicinity of prior–austenite grain boundaries, as shown in Fig. 9.17.44 The preferential recovery promotes preferential and localized creep deformation in the vicinity of prior-austenite grain boundaries, resulting in a premature creep rupture. The dissolution of MX and the precipitation of the Z phase promote preferential recovery.9
9.4.3
Progressive coarsening of M23C6
Under the condition of no Z phase formation, a rapid loss of creep rupture strength was observed for a 9Cr–3W–3Co–0.2V–0.05Nb steel with 0.08% carbon but no addition of nitrogen, for long periods more than about 1000 h at 650°C as shown in Fig. 9.18.25 The residual nitrogen content of the steel was only 0.0019% (19 ppm), suggesting an extremely low content of MX carbonitrides and hence an extremely low driving force for Z phase formation during creep. Therefore, the dissolution of fine M2X and MX carbonitrides and the precipitation of the Z phase can be excluded from the main explanation for the loss of creep rupture strength. The proposed mechanism is the coarsening of M23C6 carbides and the recovery of lath martensitic microstructure in the vicinity of prior-austenite grain boundaries. The addition of a small amount of boron, about 100 ppm, effectively suppresses the coarsening of M23C6 carbides in the vicinity of prior-austenite grain boundaries and hence suppresses the rapid loss of creep rupture strength.
9.4.4
Loss of creep ductility
Maruyama et al.3 pointed out that loss of ductility is the origin of the loss of creep rupture strength in 11Cr 2W–0.3Mo–CuVNb steel at 650°C. The reduction in area measured after creep rupture was only 11% at 650°C and 100 MPa, while it was 86% at 700°C and 100 MPa. The following is their scenario of the loss of creep rupture strength. Enhanced recovery of subgrain structure takes place along grain boundaries. Strain concentration along the boundary regions forms grain boundary cracks. This results in low ductility, premature failure and loss of rupture strength.
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500 300
Stress (MPa)
550 °C 600 °C 100 80
650 °C
60 700 °C
40 Predicted (εr = 30%) 20 101
102
725 °C
103 104 Time to rupture (h)
105
(a)
1 µm (b)
9.17 (a) Stress versus time to rupture curves for Mod.9Cr–1Mo steel (T91) and (b) TEM micrograph after creep rupture testing for 34 141 h at 600°C and 100 MPa.
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Creep-resistant steels 200
Stress (MPa)
150
100
80
60 50 101
0 ppmB 48 ppmB 92 ppmB 139ppmB 102
103 Time to rupture (h)
104
105
9.18 Loss of creep rupture strength in 9Cr–3W–3Co–VNb steel without boron at 650°C and for long periods of time. The addition of small amount of boron suppresses the degradation.
9.4.5
Recovery of excess dislocations resulting from low-temperature tempering
Excess dislocations resulting from low-temperature tempering cause a rapid decrease in creep rupture strength for 12Cr–1Mo–1W–VNb steel for long periods of time at 600–650°C, as described in Section 9.2.3.14 This is because excess dislocations accelerate recovery and recrystallization during creep with the aid of stress, which also promotes microstructure evolution during creep. 12Cr turbine steels are more susceptible to a rapid loss of creep strength for long periods of time than 9–12Cr boiler steels, because the dislocation density after tempering is much higher in 12Cr turbine steels than in boiler steels.
9.4.6
Effect of δ-ferrite
Dual phase 12Cr steel (12Cr–0.4Mo–2W–CuVNb, specified as KASUS410J3DTB) exhibits a rapid decrease in creep strength at 600–650°C for long periods of time. Kimura et al.45 proposed that the degradation is caused by inhomogeneous creep deformation around δ-ferrite. Because diffusion of carbon and nitrogen is promoted by the large concentration gap across the interface between martensite and δ-ferrite, enhanced diffusion promotes the coarsening of precipitates and microstructure recovery. Igarashi et al.46 reported that heterogeneous creep deformation at low stresses is a main factor in the degradation in long-term creep strength. They pointed out that a decrease in
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the density of fine MX and non-uniform distribution of MX in δ-ferrite promotes heterogeneous creep deformation at low stresses.
9.5
Future trends
The experimental results described in this chapter suggest that a homogeneous dispersion of thermally stable and fine precipitates not only along grain boundaries but also inside the grain provides us with an ideal microstructure for creep-resistant ferritic and austenitic steels in terms of long-term stability. The inhomogeneous nature of microstructure evolution during creep at low stresses, such as preferential recovery in the vicinity of grain boundaries, should also be taken into account for the minimization of creep strength loss for long time periods. The precipitation of the Z phase and the σ phase enhances inhomogeneous microstructure evolution in tempered martensitic and austenitic steels, respectively. In tempered martensitic 9–12Cr steels, the key issues for the improvement of long-term creep strength seem to be the long-term stabilization of subboundary hardening in the vicinity of prior-austenite grain boundaries. The stabilization of precipitation or dispersion hardening enhances the sub-boundary hardening. In future, therefore, much effort should be paid to clarify the mechanisms of precipitation and coarsening of carbonitrides and intermetallic compounds at the lath-block and prior-austenite grain boundaries, the movement of lath and block boundaries and the development of lath and block in the vicinity of prior-austenite grain boundaries.
9.6
References
1 Takeuchi S and Argon A S, ‘Steady-state creep of single-phase crystalline matter at high temperature’, J Mater Sci, 1976, 11, 1542–1566. 2 Meier M and Blum W, ‘Modelling high temperature creep if academic and industrial materials using the composite model’, Mater Sci Eng, 1993, A164, 290–294. 3 Maruyama K, Sawada K and Koike J, ‘Strengthening mechanisms of creep resistant tempered martensitic steel’, ISIJ Int, 2001, 41, 641–53. 4 Abe F, ‘Bainitic and martensitic creep-resistant steels’, Curr Opinion Solid State Mater Sci, 2004, 8, 305–39. 5 Abe F, Igarashi M, Fujitsuna N, Kimura K and Muneki S, ‘Research and development of advanced ferritic steels for 650°C USC boilers’, Proceedings of the 6th Liege Conference on Materials for Advanced Power Engineering 1998, Liege, Belgium, 1998, 259–268. 6 Murata Y, Morinaga M and Hashizume R, ‘Development of ferritic steels for steam turbine rotors with the aid of a molecular orbital method’, Proceedings of the 4th International Charles Parsons Turbine Conference, Newcastle, UK, 1997, 270–282. 7 Kadoya Y and Shimizu E, ‘Effect of solute Mo, W and dispersoid carbonitride on high-temperature creep of ferritic steels’, Tetsu-to-Hagane, 1999, 85, 827–834 (in Japanese)
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8 Maruyama K and Nakashima H, Materials Science for High Temperature Strength, Uchida-Rokakuho, Tokyo, 1997 (in Japanese). 9 Kimura K, Kushima H, Abe F, Suzuki K, Kumai S, Satoh A, ‘Microstructural change and degradation behaviour of 9Cr–1MoVNb steel in the long term’, Proceedings of the 5th International Charles Parsons Turbine Conference, Cambridge, UK, 2000, 590–602. 10 Abe F, ‘Metallurgy for long-term stabilization of ferritic steels for thick section boiler components in USC power plant at 650°C’, Proceedings of the 8th Liege Conference on Materials for Advanced Power Engineering 2006, Liege, Belgium, 2006, 965– 980. 11 Yoshizawa A, Fujita T, Yoshida F and Nakashima H, ‘Dispersion hardening mechanism of Y2O3 dispersed ferritic steel at high temperature’, Tetsu-to-Hagane 1996, 82, 865– 869 (in Japanese). 12 Srolovitz D J, Perkovic-Luton R A and Luton M, J. Phil Mag A 1983, 48, 795–809. 13 Yoshida F and Nakashima H, ‘Threshold stress for high-temperature deformation of dispersion-strengthened alloys with incoherent dispersoids’, Key Eng Mater, 2000, 171–174, 261–268. 14 Iseda A, Teranihsi H and Masuyama F, ‘Effects of chemical compositions and heat treatments on creep rupture strength of 12%Cr heat resistant steels for boiler’, Tetsuto-Hagane, 1990, 76, 1076–1083 (in Japanese). 15 Grant N J, Bucklin A G and Rowland W, ‘Creep-rupture properties of cold-worked type 347 stainless steel’, Trans ASM, 1955, 48, 446–455. 16 Abe F, Kimura K, Baba E, Kanemaru O and Yagi K, ‘Creep curve analysis and creep life evaluation of 10Cr–30Mn austenitic steels’, Proceedings International Symposium on Materials Aging and Component Life Extension, Bicego V, Nitta A and Viswanathan R (eds), Milan, Italy, 1995, 1075–1084. 17 Abe F, ‘Effect of quenching, tempering and cold rolling on creep deformation behavior of a tempered martensitic 9Cr–1W steel’, Metall Mater Trans A, 2003, 34A, 913– 925. 18 Shingledecker J P, Maziasz P, Evans N D and Pollard M J, ‘Creep behavior of a new cast austenitic alloy’, Proceedings of ECCC Creep Conference, London UK, 2005, 99–109. 19 Sawada K, Kubo K, Hara T and Abe F, ‘Distribution of MX carbonitrides and its effect on creep deformation in 9Cr–0.5Mo–1.8W–VNb steel’, Proceedings of the 7th Liege Conference on Materials for Advanced Power Engineering 2002, Liege, Belgium, 2002, 1181–1188. 20 Abe F, Nakazawa S, Araki H and Noda T, ‘The role of microstructural instability on creep behavior of a low radio-activation martensitic 9Cr–2W steel’, Metall Trans 1992, 23A, 469–477. 21 Straub S, Meier M, Ostermann J and Blum W, ‘Development of microstructure and strengthening in ferritic steel X20 CrMoV 12 1 at 823K during long-term creep tests and during annealing’, VGB Kraftwerkstechnik, 1993, 73, 646–653. 22 Sikka V K, Klueh R L, Maziasz P J, Babu S, Santella M L, Jawad M H, Paules J R and Orie K E, ‘Mechanical properties of new grades of Fe–3Cr–W alloys’, Amer Soc Mech Eng Pressure Vessel and Piping, 2004, 476, 97–106. 23 Miyata K and Sawaragi Y, ‘Effect of Mo and W on the phase stability of precipitates in low Cr heat resistant steels’, ISIJ Int, 2001, 41, 281–289. 24 Chen Z, Shan Z-W, Wu N Q, Sikka V K, Hua M H and Mao S X, ‘Fine carbidestrengthened 3Cr–3WVTa bainitic steel’, Metall Mater Trans, 2004, 35A, 1281– 1288.
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25 Taneike M, Abe F and Sawada K, ‘Creep-strengthening of steels at high temperatures using nano-sized carbonitride dispersions’, Nature 2003, 424, 294–296. 26 Abe F, Horiuchi T, Taneike M and Sawada K, ‘Improvement of creep strength by boron and nano-size nitrides for tempered-martensitic 9Cr–3W–3Co–VNb steel at 650°C’, Proceedings of the 6th International Charles Parsons Turbine Conference, Dublin, Ireland, 2003, 379–396. 27 Horiuchi T, Igarashi M and Abe F, ‘Improved utilization of added B in 9Cr heat– resistant steels containing W’, ISIJ Int, 2002, 42, S67–71. 28 Masuyama F, ‘12Cr–2.6W–2.5Co–0.5Ni–V–Nb steel’, in: Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology, Group VIII: Advanced Materials and Technologies, Volume 2 Springer-Verlag, Berlin, 2004, 200–203. 29 Masuyama F, ‘12Cr–3W–3Co–V–Nb–Ta–Nd–N steel’, in: Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology, Group VIII: Advanced Materials and Technologies, Volume 2 Springer-Verlag, Berlin, 2004, 204–205. 30 Ukai S and Fujiwara M, ‘Perspective of ODS alloys application in nuclear environments’, J Nucl Mater, 2002, 307–311, 749–57. 31 Sakuraya K, Okada H and Abe F, ‘Influence of heat treatment on formation behavior of boron nitride inclusions in P122 heat resistant steel’, ISIJ Int, 2006, 46, 1712– 1719. 32 Igarashi M, Muneki S, Hasegawa H, Yamada K and Abe F, ‘Creep deformation and the corresponding microstructural evolution in high-Cr ferritic steels’, ISIJ Int, 2001, 41, S101–105. 33 Okada H, Muneki S, Yamada K, Okubo H, Igarashi M and Abe F, ‘Effects of alloying elements on creep properties of 9Cr–3.3W–0.5Pd–V, Nb, N, B steels’, ISIJ Int, 2002, 42, 1169–1174. 34 Toda Y, Iijima M, Kushima H, Kimura K and Abe F, ‘Effect of Ni and heat treatment on long-term creep strength of precipitation strengthened 15Cr ferritic heat resistant steels’, ISIJ Int, 2005, 45, 1747–1753. 35 Takahashi T, Sakakibara M, Kikuchi M, Ogawa T, Araki S and Fujita T, ‘Elevatedtemperature strength and hot corrosion resistance of 20Cr–25Ni steel for tubes in ultra-supercritical power boilers’, Tetsu-to-Hagane, 1990, 76, 1131–1138 (in Japanese). 36 Semba H, Igarashi M, Yamadera Y, Iseda A and Sawaragi Y, ‘High temperature strength and microstructure of the 23Cr–43Ni–6W steel for USC boilers’, Report of the 123rd Committee on Heat Resisting Materials and Alloys, Japan Society for the Promotion of Science, 2003, Volume 44, 119–127. 37 Igarashi M, Semba H and Okada H, ‘Development of high strength austenitic steels for 700°C USC plant’, Proceedings of the 8th Ultra-Steel Workshop, National Institute for Materials Science (NIMS), Tsukuba, Japan, 2004, 194–199. 38 Sklenicka V, Kucharova K, Svoboda M, Kloc L, Bursik J and Kroupa A, ‘Long-term creep behavior of 9–12%Cr power plant steels’, Mater Character, 2003, 51, 35–48. 39 Danielsen H K and Hald J, ‘Behaviour of Z phase in 9–12%Cr steels’, Energy Mater, 2006, 1, 49–57. 40 Sawada K, Taneike M, Kimura K and Abe F, ‘Effect of nitrogen content on microstructural aspects and Creep behavior in extremely low carbon 9Cr heat-resistant steel’, ISIJ Int, 2004, 44, 1243–1249. 41 Vodarek V and Strang A, ‘Effect of nickel on the precipitation processes in 12CrMoV steel during creep at 550°C’, Scripta Mater, 1998, 38, 101–6. 42 Vodarek V and Strang A, ‘Compositional changes in minor phases present in
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12CrMoVNb steels during thermal exposure at 550°C and 600°C’, Mater Sci Technol, 2000, 16, 1207–13. Abe F. ‘12Cr–1Mo–1W–0.3V steel’, in: Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology, Group VIII: Advanced Materials and Technologies, Volume 2, Springer-Verlag, Berlin, 2004, 161–169. Kushima H, Kimura K and Abe F, ‘Degradation of Mod.9Cr–1Mo steel during longterm creep deformation’, Tetsu-to-Hagane, 1999, 85, 841–847 (in Japanese). Kimura K, Sawada K, Kushima H and Toda Y, ‘Degradation behaviour and long-term creep strength of 12Cr ferritic creep resistant steels’, Proceedings of the 8th Liege Conference on Materials for Advanced Power Engineering 2006, Liege, Belgium, 2006, 1105–1116. Igarashi M, Yoshizawa M, Iseda A, Matsuo H and Kan T, ‘Long-term creep strength degradation in T122/P122 steels for USC power plants’, Proceedings of the 8th Liege Conference on Materials for Advanced Power Engineering 2006, Liege, Belgium, 2006, 1095–1104.
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10 Precipitation during heat treatment and service: characterization, simulation and strength contribution E . K O Z E S C H N I K and I. H O L Z E R Graz University of Technology, Austria
10.1
Introduction
Precipitation hardening is one of the most prominent ways of strengthening materials. Precipitates can effectively hinder dislocation and subgrain movement and thus increase the resistance of the material microstructure against plastic deformation. In industrial processes, size and number density of precipitates are controlled by the chemical composition of the alloy as well as the thermomechanical processing route. Owing to the significant influence of precipitates on the mechanical properties of the material, efficient characterization, modelling and simulation of precipitation processes in multicomponent alloys are of considerable relevance for industry as well as for academics. The goal of these activities is to be able to produce materials with an optimized spectrum of mechanical properties based on a fundamental understanding of the complex interactions between precipitates and microstructure. Materials with superior creep properties are characterized by a microstructure, which exhibits a superior long-term resistance against plastic deformation. This can be achieved by strong pinning forces upon dislocations and subgrain boundaries. The two major effects of precipitates on the creep properties of a material are: •
•
Increase of the creep strength by direct interaction between precipitates and dislocations. Precipitates effectively hinder dislocations in their ability to move through the material as a consequence of an external load. Thus, the creep process is considerably slowed down and the creep rate is minimized. Stabilization of the initial microstructure by pinning of grain and subgrain boundaries. The high strength of the materials in the as-received condition, i.e. the conditions in the delivery state before service, is conserved, because grain and sub-grain coarsening is minimized. The interactions between microstructure and creep properties have been 305 WPNL2204
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investigated in numerous experimental and theoretical studies.1–8 They are described in other chapters of this book in detail, e.g. Chapter 9 by Abe and Chapter 13 by Blum. Figure 10.1 shows a schematic picture of the evolution of a typical microstructure as observed in ferritic/martensitic 9–12%Cr steels. The schematic emphasizes the role of precipitates in conserving the favourable fine-grained microstructure. After solidification and cooling to room temperature, the microstructure of typical 9–12%Cr steels consists primarily of martensite, usually with coarse and elongated precipitates along the prior austenite grain boundaries (see Fig. 10.1 left). In the course of the austenitization and quality heat treatments, a dense distribution of fine precipitates is produced (see Fig. 10.1 centre). If the material is exposed to long-term thermal and mechanical loading during service, coarsening of the precipitate microstructure occurs (see Fig. 10.1 right). Simultaneously, the material softens because the mean distance between individual precipitates increases, which leads to a decrease of the effective pinning force. In the following sections, the evolution of precipitates throughout the entire manufacturing process is investigated. First, the results of a comprehensive experimental characterization of the precipitate microstructure at several stages of the casting, austenitization and quality heat treatment processes are briefly reviewed. Then, a theoretical model for simulation of the precipitation process is outlined and typical simulation results are presented. The interaction of precipitates with dislocations and subgrains and the consequences for the strength of the material are discussed. The chapter concludes with a quantitative analysis of the loss of precipitation strengthening during service for the example of a typical 9%Cr steel.
10.2
Microstructure analysis of the COST alloy CB8
The experimental test alloy CB8, which is exemplarily used in the following analyses, has been designed in the COST programme 522. The COST variant HEAT TREATMENT
CREEP LOAD
10.1 Schematic evolution of the microstructure during heat treatment and service. Note that all images relate to the same length scale.
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CB8 has been selected because the most complete experimental picture of the precipitate evolution is available for this melt. These data are finally used for verification of the computer simulations. It is a typical 9%Cr steel for cast application, which showed excellent creep properties in short-term tests. For this reason, the variant CB8 has also been extensively investigated at longer times. However, a significant drop in creep strength has been observed after approximately 10 000 h of service exposure. The most important results of the experimental characterization are summarized subsequently. The chemical analysis of the steel CB8 is given in Table 10.1.
10.2.1 Precipitate evolution during manufacturing In extensive work of Sonderegger9 and Plimon,10 the evolution of precipitates has been investigated for the first time at different positions (Pos. 1-4) in the heat treatment process. The thermal profile of this treatment is shown in Fig. 10.2. It is typical for industrial components and already applied for other 9– 12%Cr cast steels. In the ‘as-cast’ condition (Pos. 1), the specimen microstructure consists mostly of martensite. Only small amounts of retained austenite could be detected by X-ray diffraction (0.8%). The size of the primary austenite grains is between 0.5 and 2.0 mm. The martensite lath width as well as the subgrain Table 10.1 Chemical composition of steel CB8–heat 173 (in wt%) C 0.17
Si 0.27
Mn 0.2
Cr 10.72
Ni 0.16
Mo 10.40
W –
V 0.21
Nb Co 0.060 2.92
Al B(ppm) 0.028 112
N 0.0319
1400 Solidification Austenitizing 1080°C/8h
1200 1000
T (°C)
FA 800
Tempering 730°C/10 h
600
SA
Stress relieving 730°C/12 h 730°C/14 h Furnace (45K/h)
SA
400 Pos. 1
200
Pos. 2
Pos. 4
Pos. 3
0 0
50
100
150 t (h)
200
250
300
10.2 Heat treatment of COST alloy CB8 with specimen positions for experimental characterization.
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size is approximately 1 µm. These parameters remain almost constant throughout the entire heat treatment. Figure 10.3 shows an optical micrograph of the as-cast microstructure, with prior austenite grain boundaries clearly visible as thin white lines. On detailed investigation of the prior austenite grain boundaries (PAGB), elongated precipitates have been detected (see Fig. 10.4). The chemistry and crystal structure of the precipitates indicates that these are large Mo-rich precipitates (probably Mo3B2) and Cr-rich precipitates (probably M7C3 and M23C6) with diameters of approximately 200 nm. Precipitates of a similar type also occur in the strongly segregated interdendritic regions (dark regions in Fig. 10.3). The interior of the prior austenite grains is otherwise more or less free from precipitates (see Fe-jump ratio TEM image in Fig. 10.5). After austenitization (Fig. 10.6), a certain degree of homogenization of the segregated concentration peaks is observed. The Mo3B2 precipitates at the PAGB have more or less disappeared. Moreover, NbC precipitates, which are randomly distributed in the matrix, and small needle-shaped Mo-rich precipitates are identified. The latter disappear again during further heat treatment. After the first quality heat treatment (Pos. 3), a significant increase in the number of precipitates is observed (Fig. 10.7). M23C6 precipitates are found in great quantities as well as VN particles. Both types of precipitate appear preferentially at the martensite lath and subgrain boundaries. In the sample corresponding to the ‘as-received’ condition (Pos. 4), the
1 mm
10.3 Optical micrograph of microstructure in the as-cast condition. Segregated regions from solidification are clearly observed as well as prior austenite grain boundaries.
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10.4 Elongated precipitates along the PAGB (scanning electron micrograph, SEM).
0.5 µm
10.5 CB8 in as-cast condition. Fe jump-ratio image (transmission electron micrograph, TEM).
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0.5 µm
10.6 CB8 after austenitizing (TEM bright field).
0.5 µm
10.7 CB8 after first heat treatment (TEM bright field).
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first indications of slight coarsening of M23C6 precipitates are observed (Fig. 10.8). The volume fraction of VN and the mean precipitate radius have significantly increased. NbC is also still present in this heat treatment condition. According to Dimmler et al.,11 the first sparse Laves phase precipitation can be found in the as-received condition.
10.2.2 Precipitate evolution during service In the analysis of the development of the microstructural parameters in the course of long-term service of CB8 at 650°C for 2000 and 7000 hours, Sonderegger9 found that the martensite lath width and the subgrain size remain almost constant in samples without mechanical loading. In samples which have been exposed to creep loading, the subgrain size slightly increases from 0.7 µm to 1.0 µm. Moreover, clear indications of coarsening of M23C6 precipitates are observed in the heat-treated and creep-loaded samples. In contrast, the mean precipitate radius of VN remains almost constant in both samples, while the number density increases visibly. The phase fraction of Laves phase increases significantly during service. Dimmler11 found that the phase fraction increases from 0.4 to 0.8% for samples analysed after 50 and 16 000 h, respectively. Simultaneously, the mean radius also increases. In the first 1 000 h, an increase of the number
0.5 µm
10.8 CB8 in as-received condition (TEM bright field).
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density of Laves phase precipitates is observed, indicating significant nucleation of new precipitates of this phase. After 1 000 h, the number density remains constant. After 16 000 hours of service, the number density of the VN precipitates suddenly reduces drastically. This tendency is stronger in the creep-loaded samples, which indicates that the external load enhances the microstructure evolution. The reason for the VN dissolution is found in the appearance of the modified Z-phase, which has been detected in the sample with 16 000 hours’ service time. The Z-phase has a higher thermodynamic stability compared to VN, however, nucleation of this phase is very difficult, which is the reason that the Z-phase was not observed earlier. The role of the Zphase in the drop of creep resistance in many 9–12%Cr steels is also discussed in Chapter 9 by Abe.
10.3
Modelling precipitation in complex systems
Since the interaction of the various different precipitates in the different stages of the thermal treatment of advanced materials is rather complex, a theoretical approach has recently been developed for simulation of precipitation in these multi-component, multi-phase systems, taking into account all thermodynamic and kinetic interactions of the different alloying elements. This approach is briefly reviewed here and parameters relevant for the actual precipitation simulation in the test alloy CB8 are discussed.
10.3.1 The precipitation kinetics model Consider a unit volume of a multi-component alloy. Allow an arbitrary number of spherical precipitates to nucleate and grow on random locations in this volume. The corresponding situation is sketched in Fig. 10.9. The total Gibbs energy of this thermodynamic system can be written as:12 n
m
i =1
k =1
G = Σ N0i µ 0i + Σ
m n 4 πρk3 Σ + 4 πρk2 γ k λ µ c Σ + 3 k i =1 ki ki k =1
[10.1]
where N0i is the number of moles of component i in the matrix phase, µ0i is the corresponding chemical potential, λk is the energy contribution due to volumetric misfit, ρk is the radius of the precipitate with index k, cki is the concentration of component i, µki the corresponding chemical potential and γk is the interfacial energy. In thermodynamic equilibrium, the Gibbs free energy is a minimum. Since real systems during heat treatments are in a highly non-equilibrated state, driving forces exist for evolution of the precipitate microstructure such that G is minimized. With each microstructural process that occurs in the system, part of the free energy is dissipated. In the model, three dissipative processes
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10.9 Modelling precipitation in complex materials. Spherical precipitates in a multi-component matrix.
have been considered: (i) Migration of interfaces with a mobility Mk, (ii) diffusion of atoms in the precipitates and (iii) diffusion of atoms in the matrix. Detailed expressions for these quantities are described in Svoboda et al.12 and Kozeschnik et al.13 With the total Gibbs free energy and the corresponding dissipation terms, the thermodynamic extremal principle14,15 can be applied and a linear system of rate equations for the change of radius and chemical composition of each individual precipitate is obtained. To find the evolution of the entire precipitate population, the rate equations are integrated numerically under the constraint of mass conservation. The corresponding algorithm has been implemented in the MatCalc software,16 which is used in the following sections to compute the precipitate evolution in the steel CB8.
10.3.2 Treatment of multi-component nucleation A most important step in modelling precipitation is the accurate treatment of the precipitate birth process, that is the nucleation stage. In the present model, nucleation of precipitates is dealt with in the framework of an extension of classical nucleation theory (CNT). According to this theory, the nucleation rate J, which describes the frequency of creation of new precipitates in unit time and unit volume, is given by:
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[10.2]
where Z is the Zeldovich factor, NC is the number of potential nucleation sites, β* is the atomic attachment rate, G* is the energy barrier to form a critical nucleus, k is the Boltzmann constant, T is the absolute temperature and τ is the incubation time. The term ‘extended’ CNT emphasizes that, in multi-component systems, some quantities in Equation [10.2] must be reformulated to apply to multicomponent situations. These expressions have been summarized and discussed recently by Kozeschnik et al.17
10.3.3 Linking microstructure and precipitate nucleation Most quantities in Equation [10.2] for the nucleation rate are either thermodynamic quantities or kinetic quantities related to the diffusivity of the atoms. They are global parameters and can be obtained from independent thermodynamic and kinetic databases. In contrast, the number of potential nucleation sites, NC, strongly depends on the microstructure of the material and the type of heterogeneous nucleation site which is preferred by the particular precipitate. This feature allows for a straightforward consideration of the material microstructure, that is, the grain and subgrain size and dislocation density, in the nucleation stage of the precipitation simulation. In order to obtain the correct number of available nucleation sites, a simple and yet realistic representation of the real microstructure is desirable. In the software MatCalc, the grain and subgrain structure of the matrix is approximated by an arrangement of tetrakaidecahedrons (Fig. 10.10), which are space-filling objects with 14 surface elements. In the symmetric geometry,
d
λ
d H D
10.10 : Tetrakaidecahedrons representing the matrix microstructure and determining the number of potential nucleation sites in the precipitation simulation.
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these objects resemble globular grains. If this structure is elongated, a good representation of martensite laths can be given. On the basis of the real, experimentally observed microstructure, the number of potential nucleation sites, for example, on prior austenite grain boundaries or subgrain boundaries, can be evaluated. The mathematical expressions for these quantities are described by Rajek.18 Table 10.2 presents an overview of the number of potential nucleation sites at different heterogeneous lattice positions, if each atomic position is considered to represent a potential nucleation site.
10.4
Computer simulation of the precipitate evolution in CB8
In this section, the kinetic model, which has been outlined previously, is applied to the simulation of the precipitate evolution during the entire heat treatment and service of steel CB8. First, a thermodynamic equilibrium analysis of this steel is performed, which provides an overview of the type and amount of different phases that can be expected to occur at a given chemical composition and temperature. Then, the results of the kinetic simulation are discussed.
10.4.1 Thermodynamic equilibrium analysis The thermodynamic equilibrium analysis is an important step in comprehensive material characterization. By this method, information can be obtained about: •
the type and number of phases which occur in this material under equilibrium conditions,
Table 10.2 Number of nucleation sites (m–3) at 650°C in a stretched tetrakaidecahedron with a subgrain structure. Dislocation density in austenite is assumed to be ρ = 1011m–2, the dislocation density in ferrite ρ = 1014m–2, the austenite and ferrite grain sizes are 100 µm and the ferrite subgrain size is 0.1 µm with an elongation factor s = 100. s0 is the lattice constant, afcc and abcc are the mean atomic distances Nucleation sites (m–3)
Ferrite (bcc) a0 = 2.87 × 10–10 m abcc = 2.27 × 10-–0 m
Bulk (B) Dislocations (D) Grain boundary (GB) Grain boundary edge (E) Grain boundary corner (C) Subgrain boundary (SGB)
8.27 4.36 7.63 5.13 1.68 2.00
× × × × × ×
1028 1023 1023 1018 1013 1026
bcc = body centred cubic; fcc = face centred cubic
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× × × × ×
1028 1020 1023 1018 1013
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chemical composition of the equilibrium phases, equilibrium solution temperatures of the individual phases, giving indications of temperatures for optimum heat treatment conditions, equilibrium transformation temperatures at which allotropic transformations (e.g. fcc (face centred cubic)–bcc (body centred cubic)/ bct (body centred tetragonal) transition) occur.
Figure 10.11 shows calculated phase diagrams for the steel CB8 as a function of carbon and nitrogen content. All simulations are based on the thermodynamic database TCFE3 and the diffusion database Mobility_v21 from ThermoCalc AB, Stockholm, Sweden. To account for the stabilizing effect of silicon on the Laves phase, the corresponding parameters have been modified as suggested by Dimmler.11 Moreover, although not considered in the phase diagrams of Fig. 10.11, a revised thermodynamic description for the modified Z-phase19 has been added to this database, which is a further development of the initial assessment of Danielsen and Hald.20 These values have been used in the kinetic simulations presented in the following section.
10.4.2 Precipitation simulation in CB8 With the theoretical model for multi-component precipitation kinetics and the thermodynamic and kinetic data described in the previous section, the entire heat treatment of the alloy CB8 has been studied. Figure 10.12 summarizes the results of the simulation. When looking at the temperature profile of the heat treatment (top image (a) in Fig. 10.12), several individual steps can be distinguished. The simulation starts at 1400°C, which is just below the solidus temperature of this alloy. It is assumed that all elements are homogeneously distributed in the matrix at this time and no precipitates exist. The material then cools linearly to a temperature of 350°C. This temperature corresponds to the observed austenite to martensite transformation start temperature. In the simulation, it is assumed that this transformation occurs instantaneously and the parent and target phases have identical chemical composition. It is further assumed that no diffusive processes and, consequently, no precipitation occurs below this temperature. At this point, the matrix phase is changed from a face centred cubic (fcc) austenite to a body centred cubic (bcc) ferrite structure. Since no separate thermodynamic description is available for the bct martensite phase, the bct phase is substituted by the bcc phase in the simulations. In the next step, the material is reheated for austenitization. At the experimentally observed transformation temperature of 847°C, the ferrite matrix is changed to austenite again. After austenitization, the transformation to martensite/ferrite is performed again at 350°C. After three quality heat treatments, service at 650°C for 100 000 h is simulated.
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(a)
317
1 2
3
1400
4
5
6 7
1200 8
10
9 11
T (°C)
1000 12
13
14
19 15
17 800
16 22
24
23
25
20
28
21 600
27
26
29
400 0
0.02
0.04
0.06
0.08
0.1
XN(%) 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.
Liquid Liquid, δ Liquid, δ, γ Liquid, δ, γ, M2B tetr. δ, γ, M2B tetr. γ, M2B tetr. γ, M2B tetr., AIN γ, M2B tetr., NbC γ, M2B tetr., AIN, NbC γ, M2B tetr., AIN, VN γ, M2B tetr., AIN, NbC, VN γ, M2B tetr., NbC, M23C6 γ, M2B tetr., AIN, NbC, M23C6 γ, M2B tetr., AIN, NbC, VN, M23C6 γ, M2B tetr., AIN, VN, M23C6
16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29.
α, γ, M2B tetr., NbC, M23C6 α, γ, M2B tetr., AIN, NbC, M23C6 α, γ, M2B tetr., AIN, NbC, VN, M23C6 α, γ, M2B tetr., AIN, VN, M23C6 α, γ, M2B tetr., AIN, VN, M23C6, Cr2N α, M2B tetr., NbC, M23C6 α, M2B tetr., AIN, NbC, M23C6 α, M2B tetr., AIN, NbC, VN, M23C6 α, M2B tetr., AIN, VN, M23C6 α, M2B tetr., AIN, VN, M23C6, Cr2N α, M2B tetr., AIN, NbC, M23C6, Laves α, M2B tetr., AIN, NbC, VN, M23C6, Laves α, M2B tetr., AIN, VN, M23C6, Laves α, M2B tetr., AIN, VN, M23C6, Cr2N, Laves
10.11 Phase diagrams for steel CB8 obtained from computational thermodynamics as function of (a) C and (b) N, respectively. The modified Z-phase, which replaces VN and, eventually, NbC precipitates at long times, has been neglected; the boride phases are included.
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1600
(b)
1 5 2
1400
6 3
4
7
8
9 10 1200
11 12 16
1000
13
17
14
15
T (°C)
18 20
19
21
22
23
800
24
25 26
600
27
400 0
0.05
1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.
0.1
Liquid Liquid, δ Liquid, δ, γ Liquid, δ, M2B tetr. Liquid, γ Liquid, γ, M2B tetr. δ, M2B tetr. Liquid, δ, γ, M2B tetr. δ, γ, M2B tetr. γ, M2B tetr. γ, M2B tetr., NbC δ, γ, M2B tetr., AIN γ, M2B tetr., AIN γ, M2B tetr., NbC, AIN
0.15
0.2 XN(%) 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27.
0.25
0.3
0.35
0.4
γ, M2B tetr., AIN, VN α/δ, γ, M2B tetr., AIN, VN γ, M2B tetr., NbC, AIN, M23C6 γ, M2B tetr., NbC, AIN, VN γ, M2B tetr., NbC, AIN, VN, M23C6 α/δ, γ, M2B tetr., NbC, AIN, VN α/δ, M2B tetr., AIN, VN α/δ, γ, M2B tetr., NbC, AIN, VN, M23C6 α/δ, M2B tetr., NbC, AIN, VN α/δ, M2B tetr., NbC, AIN, VN, M23C6 α/δ, M2B tetr., AIN, VN, Laves α/δ, M2B tetr., NbC, AIN, VN, Laves α/δ, M2B tetr., NbC, AIN, VN, M23C6, Laves
10.11 (Continued)
The three other plots in Fig. 10.12 display the evolution of the phase fraction, Fig. 10.12(b), the mean precipitate radius, Fig. 10.12(c), and the number density, Fig. 10.12(d), of each precipitate type. The phase fractions of M23C6 and Laves-phase are multiplied by a factor of 1/10 to give a better visual representation of the results. During cooling from 1400°C, various precipitate phases nucleate at the
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Precipitation during heat treatment and service M23C6(×0.1) M 7 C3 Laves-phase(×0.1) NbC VN
(a)
1200
T (°C)
1000
319
Z-PhaseP Z-PhaseM M23C6exp. (×0.1) MXexp. VNexp.
800 600 400 1
10
100
1000
10 000
100 000
1000
10 000
100 000
1000
10 000
t (h) 0.4
(b)
f (%)
0.3 0.2 0.1 0 1
10
100
t (h) (c)
R (nm)
100
10
1
0.1 1
10
100
100 000
t (h) 1×1024
(d)
22
N(m–3)
1×10
1×1020 1×1018 1×1016 1×1014 1
10
100
1000
10 000
100 000
t (h)
10.12 Kinetic simulation of the precipitate evolution in CB8 during (a) heat treatment and service; (b) f, phase fraction (c) R, mean precipitation radius; and (d) N, number density. The phase fractions of M23C6 and Laves-phase are multiplied by a factor of 1/10 to give a better visual representation of the results.
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austenite grain boundaries. At room temperature, NbC, VN, M23C6, M7C and Laves-phase precipitates are observed. After changing the matrix phase from austenite to ferrite and reheating, severe precipitation of all phases sets in again. Most phases, except NbC, dissolve again during further heating to and holding at the austenitization temperature of 1080°C. During subsequent cooling and the first of the three quality heat treatments, nucleation of various precipitates continues. Slight coarsening of some precipitates, particularly M23C6 is already observed. During service, the simulation predicts significant precipitation of the Laves-phase, which is in accordance with the observations of Dimmler.21 After several thousand hours of service, the phase fraction of the modified Z-phase gradually increases and, simultaneously, the thermodynamically less stable VN precipitates start to dissolve. Again, this is consistent with experimental observation.9 It should be noted that, owing to the different possibilities of Z-phase nucleation (Z-phase can be formed by heterogeneous nucleation in the matrix as well as by direct transformation of VN into Zphase),22 two Z-phase populations are introduced in the simulation. Generally, the simulation results are in reasonable agreement with the experimental data for CB8. In view of the complexity of the problem and the high degree of abstraction of the theoretical model, the overall performance of the simulation is excellent. It is particularly important to emphasize that the simulations have been performed on the basis of independent thermodynamic and diffusion databases and no general fit parameters have been used. The few necessary modifications of the original thermodynamic database and the correction for the estimated interfacial energy for the Lavesphase are well founded and described in detail by Rajek.18 In the next section, the interaction between precipitates and microstructure is analysed and in the last section, a prediction of the loss of precipitation strengthening over the lifetime of a component made from CB8 is attempted.
10.5
Microstructure–property relationships
As already pointed out in the Introduction (Section 10.1), precipitates act as microstructure-stabilizing components and as efficient obstacles for dislocation movement. In this section, the interaction forces are analysed on a quantitative basis.
10.5.1 Precipitate–dislocation interaction Precipitates and mobile dislocations can interact in one of the following ways (see e.g. McLean):23 1. A dislocation can pass coherent precipitates by cutting (breaking) the precipitate. A stacking fault is left in the precipitate;
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2. A dislocation can pass precipitates by bending between them and closing the bent lines to loops. A dislocation loop is left around the by-passed precipitate (Orowan mechanism); 3. A dislocation can pass the precipitates by climbing; 4. A dislocation can drag the precipitates with it. This mechanism is possible only for very small precipitates. The velocity-determining factor in this case is the mobility of the dragged precipitates. The precipitate–dislocation interaction mechanism in operation depends on a number of factors, among which the availability of glide planes, the height of the local forces and the hardness of the precipitates are most important. Owing to the physical nature of the four processes, mechanisms 1 and 2 are considerably faster than mechanisms 3 and 4. If the latter are the operating creep mechanisms, the creep rate is significantly lower than the creep rate based on mechanisms 1 and 2. For details, the interested reader is referred to References (23)–(27). If the precipitates are sufficiently hard, such that dislocations cannot bypass them by cutting, the upper limit of the pinning force is determined by the Orowan stress. According to Ashby,28 the Orowan stress τ0 is given by: ξ [10.3] τ 0 = C Gb ln r0 λ where C is a constant (C = 0.159 for screw dislocations and C = 0.227 for edge dislocations), G is the shear modulus, b is the Burgers vector, λ0 is the mean particle distance, r0 is the ‘inner cut-off radius’ and ξ the ‘outer cut-off radius’ of the dislocation.28
10.5.2 Precipitate–subgrain boundary interaction The pinning force of precipitates upon subgrain boundaries can be described based on a suggestion of Zener (reference (3) (private communication)) in Smith,29 which originally describes the drag force on grain boundaries during grain growth in the presence of precipitates. The basic idea is that a precipitate, which is located on a grain boundary, reduces the effective grain boundary area. On leaving the particle behind, this area must be re-established. This process requires energy and thus acts against boundary migration. A number of modifications of the original theory have since been developed, which have been reviewed by Manohar et al.30 McLean31 has pointed out that Zener’s ideas can likewise be applied to subgrain pinning. Accordingly, the mobility of subgrain boundaries is strongly reduced in the presence of precipitates, which has been observed experimentally, for instance, in the TEM investigations of Eggeler.32 The critical subgrain radius Rcrit during subgrain coarsening, above which the coarsening process comes to a stop, is given by:
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Rcrit = 0.846
1 n⋅r
[10.4]
where n is the number density of the precipitates and r is the mean precipitate radius. The effective retarding force of subgrains upon dislocation movement can be estimated according to Gladman33 and McElroy and Szkopiak34 based on the grain size-dependent contribution to the Hall–Petch relation. Taking into account Equation [10.4], the strength contribution ∆τsubgr from subgrains is given as:
∆τ subgr =
kd d sm
[10.5]
where kd represents the subgrain strengthening coefficient,33 ds is the subgrain size and m is an exponent, which is typically in the order of m = 1/2. According to Gladman,33 kd is usually much smaller than the coefficient for grains. Kosik et al.35 have considered the effect of subgrains as being comparable to cold-working. However, it has been found9 that the subgrain size changes only slightly during heat treatment and thermal ageing. Consequently, the stress contribution of subgrains remains more or less constant during service and it is therefore not considered further here.
10.5.3 Precipitate–grain boundary interaction Precipitates similarly affect the mobility of grain boundaries; however, the mechanism of pinning is different from that for subgrain boundaries. The latter are small angle boundaries and thus represent an array of dislocations, which accommodate the small lattice misfit between the two subgrains. A grain boundary is a randomly oriented high-angle boundary. In the case of 9–12%Cr steels, the fraction of grain boundaries is very small compared with subgrain boundaries because typical grain sizes for cast materials are in the order of millimetres, whereas subgrain sizes are typically in the order of micrometres or less. The influence of grain boundaries is therefore also neglected in the further analysis.
10.6
The back-stress concept
If an external force is acting on a microstructure, it is frequently assumed that the external force acts on each representative volume of the microstructure simultaneously. If the external force is high enough, plastic deformation of the material occurs via movement of dislocations and grain (subgrain) boundaries. At elevated temperatures, and if only a small external load is applied, which are conditions that are typical for creep deformation, the
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situation is different. In this case, part of the external driving pressure σe is counteracted by heterogeneous internal microstructural constituents, such as precipitates and interfaces. Consequently, the entire external load cannot be assumed to represent the driving force for the creep process; only the part of the external stress σex, which exceeds the amount of inner stress σi from the counteracting microstructure, effectively contributes to the creep process. Since the inner stress reduces the effect of the external stress, this approach is commonly denoted as the back-stress concept. The effective creep stress σeff can be expressed as: σeff = σex – σi
[10.6]
In a recent treatment by Dimmler,21 the inner stress σi has been expressed as a superposition of individual contributions from dislocations and precipitates. When also taking into account the contribution from subgrain boundaries, the inner stress is: σi = Mτi = M(τdisl + τprec + τsgb)
[10.7] 21
where M is the Taylor factor (usually between 2 and 3, see Dimmler) and τ is the shear stress. The subscripts in the terms in the parentheses denote contributions from dislocations, precipitates and subgrain boundaries, respectively. When taking into account the inner stress, the general Norton creep law (see e.g. Čadek)36 can be rewritten as: ε˙ = A ⋅ ( σ ex – σ i ) n = A ⋅ σ neff
[10.8]
where A and n are constants. When examining the individual contributions of the different mechanisms to the back-stress based on Equation [10.7], according to Taylor (cited by Dimmler),21 the part stemming from dislocations can be expressed as:
τ disl = α ⋅ G ⋅ b ⋅
ρm
[10.9]
where ρm denotes the density of mobile dislocations and the value of α is between 0.84 and 1 (see e.g. Weinert).37 The quantities G and b are defined in Equation [10.3]. The contribution of precipitates to the total back stress has already been discussed in Section 10.5.1 and is described by the critical Orowan stress τ0 (equation [10.3]). This quantity denotes the maximum back stress caused by a random distribution of precipitates using a mean distance λ between the precipitates. The latter can be estimated using the assumption that each precipitate consumes approximately the same bulk volume. In this case, λ is given by: λ=
3
6 π ⋅ ni
[10.10]
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Creep-resistant steels
where ni represents the number density of precipitates in units of m–3. When combining Equations [10.3] and [10.10], the total back-stress contribution from a bulk distribution of precipitates reads:
τ prec = CGb 3
π ⋅ ni ξ ln 6 r0
[10.11]
The quantity τprec represents the maximum back-stress caused by precipitates. If the external load reduced by the back-stress contribution of the other strengthening mechanisms is below this threshold, the dislocations are effectively pinned and can only pass the precipitates by the climb mechanism. Since dislocation climb is a diffusional process, the effective creep rates are usually low. If the threshold stress is exceeded, the dislocations can bypass the precipitates by the Orowan mechanism, which is a much faster process compared to climb. When this change in mechanism occurs, the exponent in the Norton creep law increases significantly and creep deformation is strongly enhanced. The selection of the operative creep mechanism is mainly determined by the height of the external load. However, a transition from dislocation climb to the Orowan mechanism can also be caused by a decrease in back-stress during service. If this transition occurs, for instance, owing to coarsening of precipitates or thermodynamic instability of a precipitation strengthening phase, extrapolation of the creep strength from short-term experiments to long-term service behaviour can result in fatal overestimation of the residual lifetime of components. This aspect has been discussed in detail by Dimmler and co-workers.21,38
10.7
Loss of precipitation strengthening during service of CB8
In this section, the theory and methodology described previously is applied to a prediction of the evolution of the strength contribution of precipitates during service of the steel CB8. The same calculation as shown previously in Section 10.4.2 is used as a basis and evaluated in terms of the maximum Orowan stress, that is the back-stress contribution from precipitates. For the evaluation of the back-stress during service at 650°C, the constants in Equation [10.11] are assumed to be C = 0.19 and G = 62.3 GPa (converted from the data in Guntz et al.39 The outer cut-off radius ξ is assumed to have a value of twice the precipitate radius ξ = 2rprec, the inner cut-off radius is taken as twice the Burger’s vector with r0 ~ 2b ~ 0.5 nm. Figure 10.13 shows the predicted evolution of the back-stress during service as calculated by the model described in this chapter. The graphs show the predicted contributions for each phase separately as well as the total back stress including the combined effect of all precipitates with and without the effect of the Z-phase.
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Back-stress (N mm–2 )
100 Totalwith Z-Phase
VN
80
Laves-phase
Totalwithout Z-Phase
60 40
M23C6 Z-PhaseM NbC
20
Z-PhaseP
0
Heat treatment
Service
1000
10 000
100 000
t (h)
10.13 Predicted degradation of the back stress during service of CB8 including the modified Z-phase contribution (solid bold line) and artificially suppressing this phase (dashed bold line). Subscript P denotes precipitates of Z-phase nucleated on existing VN particles, the subscript M denotes Z-phase precipitates nucleated in the matrix.
In the as-received condition, after all heat treatments, the total strength contribution from precipitation hardening is estimated to be in the order of 90 MPa. This quantity varies strongly with selection of input values and should therefore not be considered in terms of absolute values. Owing to the inevitable effect of Ostwald ripening (coarsening), the density of precipitates is reduced, which is reflected in the gradual decrease of the total back-stress up to times of approximately 10 000 h. When the modified Z-phase is included, the total back stress shows a drastic depression between 10 000 and 20 000 h. This effect is due to the enhanced nucleation and subsequent growth of the Z-phase, which causes dissolution of the VN precipitates. In the later stages, the decrease in back stress continues in a steady manner again, however, at a much lower level than before. Comparison of the curves for the integrated back stress in Fig. 10.13 indicates that Z-phase formation causes an additional back stress reduction at 100 000 h approximately 20 MPa. This effect is assumed to be responsible for the drop in creep strength of various different ferritic/martensitic creep-resistant 9–12%Cr steels during long-term creep exposure (compare with Chapter 9).
10.8
Summary and outlook
Precipitation strengthening is a key mechanism for improving mechanical properties of creep resistant materials. To capture the evolution of the precipitate microstructure during heat treatment and service in these complex materials, experimental characterization must be performed with state-of-the-art techniques and methodologies. If these are complemented with suitable
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modelling and simulation approaches, a profound understanding of the interactions of all microstructural constituents can be achieved and predictions of the influence of variations of chemical composition and process parameters can be attempted. A corresponding methodology is introduced in this chapter. With the simulation software MatCalc, which is based on independent thermodynamic and kinetic databases and on a novel theoretical approach to model multicomponent multi-phase precipitation kinetics, the strengthening contribution from precipitates is predicted for the entire life time of a sample for the example of the 9–12%Cr steel experimental test alloy COST CB8. The influence of Z-phase precipitation on the degradation of the creep strength is quantified and a significant drop of the creep strength is confirmed at the time when the VN precipitates dissolve. In further optimization of the mechanical properties of creep-resistant materials, improved understanding of the thermodynamic and kinetic interactions between the elements and phases in the microstructure of these complex alloy systems is necessary. Based on these improved input data and provided that the theoretical models describing the physical processes are realistic and accurate, kinetic modelling and simulation can aid even better in optimizing existing alloy concepts and production routes or even help to identify promising concepts for new materials for the next generation of creep-resistant materials.
10.9
References
1 F. Kauffmann, G. Zies, D. Willer, C. Scheu, K. Maile, K.H. Mayer and S. Straub, ‘Microstructural investigation of the boron containing TAF steel and the correlation to the creep strength’, 31. MPA-Seminar Werkstoff- und Bauteilverhalten in der Energie& Anlagentechnik, Stuttgart, 2005, 13–14 October. 2 A. Kostka, K.-G. Tak, R.J. Helling, Y. Estrin and G. Eggeler, ‘On the contribution of carbides and micrograin boundaries to the creep strength of tempered martensite ferritic steels’, Acta Mater., 2007, 55, 539–550. 3 J. Eliasson, A. Gustafson and R. Sandström, ‘Kinetic modelling of the influence of particles on creep strength’, Key Eng. Mater., 2000, 171–174, 277–284. 4 R. Lagneborg, ‘Effect of grain size and precipitation of carbides on creep properties in Fe–20%Cr–35%Ni alloys’, J. Iron and Steel Institute, 1969, 1503–1506. 5 F.R.N. Nabarro, ‘Grain size, stress, and creep in polycrystalline solids’, Phys. Solid State, 2000, 42, 1456–1459. 6 K.E. Amin and J.E. Dorn, ‘Creep of a dispersion strengthened steel’, Acta Metallurgica, 1969, 7, 1429–1434. 7 K. Maruyama, K. Sawada and J. Koike, ‘Strengthening mechanisms of creep resistant tempered martensitic steel’, ISIJ Int., 2001, 41, 641–653. 8 K. Sawada, K. Kubo and F. Abe, ‘Creep behaviour and stability of MX precipitates at high temperature in 9Cr–0.5Mo–1.8W–VNb steel’, Mater. Sci. Eng., 2001, A, 319–321, 784–787.
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9 B. Sonderegger, Characterisation of the Substructure of Modern Power Plant Steels using the EBSD-Method, PhD Thesis, Graz University of Technology, 2005 (in German). 10 S.W. Plimon, Simulation of an Industrial Heat Treatment and Accompanying Microstructural Investigation of a Modern 9–12% Chromium Steel, Master Thesis, Graz University of Technology, 2004 (in German). 11 G. Dimmler, P. Weinert, E. Kozeschnik and H. Cerjak: ‘Quantification of Lavesphase in advanced 9–12% Cr-steels using a standard SEM’, Mater. Character., 2003, 51, 341–352. 12 J. Svoboda, F.D. Fischer, P. Fratzl and E. Kozeschnik, ‘Modelling of kinetics in multi-component multi-phase systems with spherical precipitates I – Theory’, Mater. Sci. Eng. A, 2004, 385 (1–2), 166–174. 13 E. Kozeschnik, J. Svoboda and F.D. Fischer, ‘Modified evolution equations for the precipitation kinetics of complex phases in multi-component systems’, CALPHAD, 2005, 28 (4), 379–382. 14 L. Onsager, ‘Reciprocal relations in irreversible processes’, Phys. Rev., 1931, 37, 405–426, 1931, 38, 2265–2279. 15 J. Svoboda, I. Turek and F.D. Fischer, ‘Application of the thermodynamic extremal principle to modeling of thermodynamic processes in material sciences’, Phil. Mag., 2005, 85 (31), 3699–3707. 16 E. Kozeschnik, J. Svoboda, P. Fratzl and F.D. Fischer, ‘Modelling of kinetics in multi-component multi-phase systems with spherical precipitates II – Numerical solution and application’, Mater. Sci. Eng. A, 2004, 385 (1–2), 157–165. 17 E. Kozeschnik, J. Svoboda and F.D. Fischer, ‘On the role of chemical composition in multi-component nucleation’, invited paper in Proceedings International Conference Solid-Solid Phase Transformations in Inorganic Materials, PTM 2005, Pointe Hilton Squaw Peak Resort, Phoenix, AZ, USA, 2005, 29.5.–3.6, 301–310. 18 J. Rajek, Computer Simulation of Precipitation Kinetics in Solid Metals and Application to the Complex Power Plant Steel CB8, PhD Thesis, Graz University of Technology, 2005. 19 H.K. Danielsen, private communication. 20 H.K. Danielsen and J. Hald, Z-phase in 9–12%Cr Steels, Internal Report 863, Värmeforsk AB, 2004. 21 G. Dimmler, Quantification of Creep Resistance and Creep Fracture Strength of 9– 12%Cr Steel on Microstructural Basis, PhD Thesis, Graz University of Technology, 2003 (in German). 22 H.K. Danielsen, J. Hald, F.G. Grumsen and M.A.J. Somers, ‘On the crystal structure of Z-phase Cr(V, Nb)N’, Metall. Mater. Trans., 2006, 37A, 2633–2640. 23 M. McLean, ‘On the threshold stress for dislocation creep in particle strengthened alloys’, Acta Metallurgica, 1985, 33, 545–556. 24 L.M. Brown and R.K. Ham, in Strengthening Methods in Crystals, A. Kelly and R.B. Nicholson (eds), Elsevier, Amsterdam 1971. 25 R. Lagneborg, ‘Bypassing of dislocations past particles by a climb mechanism’, Scripta Metallurgica, 1973, 7, 605–614. 26 E. Arzt and M.F. Ashby, ‘Threshold stresses in materials containing dispersed particles’, Scripta Metallurgica, 1982, 16, 1285–1290. 27 J.D. Verhoeven, Fundamentals of Physical Metallurgy, John Wiley & Sons, New York, 1975. 28 M. Ashby, ‘The theory of the critical shear stress and work hardening of dispersion-
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29 30 31 32 33 34 35 36 37
38
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hardened crystals’, G.S. Ansell, T.D. Cooper and F.V. Lenel (eds), Metallurgical Society Conference, Vol. 47, Gordon and Breach, New York, 1968, 143–205. C.S. Smith, ‘Grains, phases and interfaces: an interpretation of microstructure’, Trans. AIME, 1948, 175, 15–51. P.A. Manohar, M. Ferry and T. Chandra, ‘Five decades of the Zener equation’, ISIJ Int., 1998, 38, 913–924. D. McLean, ‘Resistance to hot deformation’, Trans. Metall. Soc. AIME 1968, 242, 1193–1203. G. Eggeler, ‘The effect of long-term creep on particle coarsening in tempered martensite ferritic steels’, Acta Metallurgica, 1989, 37, 3225–3234. T. Gladman, The Physical Metallurgy of Microalloyed Steels, The Institute of Materials, London, 1997. R.J. McElroy and Z.C. Szkopiak, ‘Dislocation-substructure-strengthening and mechanical–thermal treatment of metals’, Int. Metall. Rev., 1972, 17, 175–202. O. Kosik, D.J. Abson and J.J. Jonas, ‘Strengthening effect of hot work subgrains at room temperature’, J. Iron and Steel Inst., 1971, 209, (88), 624–629. J. Čadek, Creep in Metallic Materials, Elsevier, Amsterdam 1988. P. Weinert, Modelling of the Creep Behaviour of Ferritic/Martensitic 9–12% Cr Steels on a Microstructural Base, PhD Thesis, Graz University of Technology, 2001 (in German). G. Dimmler, P. Weinert and H. Cerjak, ‘Extrapolation of short-term creep rupture data – The potential risk of over-estimation’, Proceedings Creep Conference, 12–14 September 2005, London, 165–176. G. Guntz, M. Julien, G. Kottmann, F. Pellicani, A. Pouilly and J.C. Vaillant, The T 91 Book – Ferritic Tubes and Pipe for High Temperature Use in Boilers, Vallourec and Mannesmann Tubes, France, 1991.
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11 Grain boundaries in creep-resistant steels R . G . F A U L K N E R , Loughborough University, UK
11.1
Introduction
The consequences of the grain boundary topology and the grain boundary structure in creep-resistant steels have been given limited prominence in proportion to their importance. Much creep deformation takes place along grain boundaries and the importance of identifying the ideal grain boundaries from the viewpoint of grain boundary engineering is still not fully appreciated. Grain boundary precipitation is also valuable in reducing grain boundary mobility and thus stabilising the grain size during heat treatment and thermal exposure. Grain boundary engineering (Palumbo, 1998) is a phase developed in the last decade to encompass all attempts to engineer the grain boundary structure to give a preferred boundary type. This type of grain boundary is typically a ∑3 twin-type boundary and it can lead to superior mechanical properties in alloys that have been tested. These alloys still remain in the copper and nickel alloy sector. An example study on IN718 (see Table 11.1 for composition) is discussed by Randle (2003), but very little work has been done on steels. Bearing in mind that creep properties are much more dominated by boundary effects, grain boundary engineering for creep strength in high alloy steels is required to consider much more than the crystallographic misorientation properties of the boundaries. Precipitation which is so prevalent in all high alloy creep resistant steels can have substantial effects on creep strength (Sun et al., 1992). This is not only because of the roughening effect that precipitates must have on grain boundary sliding but also because the precipitates will change the vacancy sink and emissive properties of the boundaries at high temperature. Furthermore, substantial segregation effects are seen on grain boundaries (Faulkner, 1996). These can be controlled by heat treatment but are rarely taken into consideration when considering the dynamics of creep deformation at high temperatures. One element that excites attention in this respect is boron. Boron is increasingly being accepted as playing a role in creep by 329 WPNL2204
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being segregated to grain boundaries in both austenitic and ferritic steels (Abe, 2006; Williams et al., 1976) and, as a result, slowing down atom transport mechanisms along the boundary plane. Dynamic recrystallisation effects at grain boundaries have been observed in ferritic steels with the result that creep deformation around the grain boundaries is more dispersed and creep life lengthened. The final aspect of grain boundaries that can lead to creep property control is the grain size itself. Traditionally, fine grain sizes have been thought to encourage creep deformation but nanocrystalline materials studies have opened up a new arena in which many of the accepted views on grain size effects on creep are being questioned. The central issue is that for many metallic nanocrystalline materials the strength decreases with decreasing grain size. This contradicts the Hall–Petch relationship but it does offer a weaker grain option in many materials that is the requirement for high creep strength materials. The origin of the anomaly lies in the very high triple junction density in these nanocrystalline materials (Suryanarayana et al., 1992). This chapter will review studies that have attempted to take into consideration the above-mentioned grain boundary effects in determining creep-strengths in creep resistant steels. The discussion will be divided between ferritic (9– 12%Cr) steels and austenitic (18%Cr–12%Ni) steels. A last consideration must concern the control of chromium behaviour on grain boundaries in high Cr (> 12%Cr) steels. The proliferation of chromium carbide on grain boundaries in these steels results in substantial Cr depletion within the sub-micrometre region close to the boundaries. These Cr depleted zones allow accelerated oxidation to take place from the surface along grain boundary planes. This process is called intergranular cracking and is the basis of other accelerated corrosion effects seen in these steels, such as intergranular stress corrosion cracking (Faulkner et al., 2005 and irradiationassisted stress corrosion cracking (Bruemmer, 1999). The subject is extensive and occurs at high temperatures, but it is not directly connected with creep properties. For this reason the subject will not be discussed further here but the interested reader is referred to a typical review (Newman, 2001).
11.2
Ferritic steels
11.2.1 General A typical ferritic high creep strength steel contains 9–12%Cr and additions of B, W, Mo, C, V, N, Mn and Si. The alloys are described in more detail elsewhere in this book. The microstructure of such a steel is shown schematically in Fig. 11.1. It can be seen that there are various grain boundary types. First the prior austenite grain boundaries define the grain structure that existed before the alloy was quenched into the martensite regime. After the quench, martensite
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10 micrometres Lath boundary
Prior Austenite GB
Carbides, nitrides, Z phase, and Laves phase
Lath packet boundary
11.1 Schematic view of high chromium ferritic steel microstructure.
is formed in small, sub-micrometre sized plates contained in packets of 10– 20 plates. These plates are often called laths and the bundles of laths are referred to as lath packets. Whilst the mechanical strengths of these lath boundaries and prior austenite boundaries are quite different, the ability of all types of boundary to nucleate chromium carbide (M23C6) is reasonably similar. Also segregation of elements like B, Hf, P and Mo is observed to occur to all types of boundary (Morgan et al., 1992) with equal intensity.
11.2.2 Grain boundary precipitation Grain boundary precipitates observed in P92 alloy (see Table 11.1 for composition) are mainly M23C6, where M stands mainly for Cr with some Mo, Fe and W (Czyrska-Filemonowicz et al., 2003; Hattestrand and Andren, 2001). The grain boundary precipitates provide creep strength to the alloy and further contribute to creep strength by pinning grain boundaries so that the grain size is stabilised during high temperature exposure. The experimental studies of grain boundary precipitates in this alloy have been performed by transmission electron microscopy on alloys that have seen a variety of tempering treatments in the temperature range 750–760°C for up to 2 h, followed by ageing at temperatures between 550 and 700°C for times up to 10 000 h (about 1 year). These studies show that most of the grain boundary precipitate is nucleated by the end of the tempering treatment and that there is a regime of stable precipitate size and spacing out to about 10 000 h at 650°C, after which coarsening occurs. It is at this stage that the creep strength is lost and tertiary creep rapidly sets in. Although VN precipitates within the grains, the intergranular M23C6 plays a substantial role in controlling the creep strength. The thermodynamics and kinetics of grain boundary precipitate evolution in ferritic steels has been modelled using a dedicated
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E911 P92 IN718 MA957 Eurofer 1.4914 20/25 PE 16 304LN 316LN
C
Mn
Si
P
S
Cr
Ni
0.12 0.11 0.04
0.35 0.35 0.06
0.19 0.19 0.40
0.007 0.008
0.003 0.003
0.22 1.00 0.06 0.47 Bal 3.00 0.3
0.11 0.11 0.049 0.08 0.027 0.021
0.42 0.35
0.06 0.45
0.05 1.46 0.44
0.72 0.28
9.10 8.96 18.2 14.0 8.9 11.3 20.0 17.1 18.8 17.7
0.012 0.012
0.012 0.019
Mo
0.7 25.0 42.5 10.5 12.1
0.5 3.1 0.06 2.2
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V
Al
Nb
W
N
0.23 0.20
0.006
0.069 0.07 5.20
0.98 1.84
0.07 0.05
0.50 0.2 0.3
B
Ti
Y
1.0
0.25 0.3
0.001 0.004
1.1 0.25 0.52
0.029
0.007
0.177 0.25
0.002 0.003
1.3
1.2
Creep-resistant steels
Table 11.1 Alloy compositions (wt%: Fe bal except where specified)
Grain boundaries in creep-resistant steels
333
Monte Carlo modelling method (Yin and Faulkner 2005). The predictions of the model have been shown to give good fits with the grain boundary precipitate experimental data for P92 (See Fig. 11.2). Various workers have used Avrami-DICTRA type relationships to forecast the evolution of M23C6 (Golpayegani et al., 2003), and to particularly demonstrate that the effect of B is to reduce the rate of coarsening in P91 type compositions by slowing down the transport of Cr between the M23C6 particles during coarsening.
11.2.3 Grain boundary segregation
Equivalent circle radius, L (nm)
There are two main types of grain boundary segregation (Faulkner, 1996). The first is Gibbs equilibrium segregation, which is caused by misfitting 100
10
1
Equivalent circle radius, L (nm)
10–2
10–1
100
101 102 Time, t (h) (a)
103
104
105
10–1
100
101 102 Time, t (h) (b)
103
104
105
100
10
1
10–2
11.2 Precipitation kinetics in P92 at 600°C (a) and 650°C (b) (see Table 11.1 for composition). Lines are model predictions (dotted, intragranular; solid, intergranular). Symbols are experimental data.
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impurity atoms seeing energy benefits from sitting on a disordered grain boundary plane as opposed to residing in the grain. Atoms diffuse under this driving force towards the boundary and the kinetics are controlled by the diffusion of impurity atoms in the ferritic steel matrix (McLean, 1957). The net result is that at high temperatures (> 600°C in ferritic steel) the energy benefit seen by the impurity becomes very limited, whereas at lower temperatures, diffusion limits the process. Therefore there is an optimum temperature where the most prevalent equilibrium segregation takes place. This is typically around 500°C in ferritic steels. The second type of grain boundary segregation is non-equilibrium segregation (Aust et al., 1967). This process occurs when the steel is rapidly quenched from a temperature like the normalising or solution treatment temperature. In these circumstances, the grain boundaries are able to preserve their equilibrium vacancy concentrations because of the good vacancy sink efficiency of the grain boundary. On the other hand, in regions away from the boundaries, in the grain centres, the vacancy concentration is maintained at high, non-equilibrium levels. Consequently, in quenched material, there exists a positive vacancy concentration gradient towards the grain boundary plane, down which the vacancies diffuse during subsequent heat treatment after the quench. It appears that in ferritic steels this gradient is similar for lath, lath packet and prior austenite grain boundaries. Impurity atoms with large misfit will be preferentially attracted to these vacancies diffusing towards the grain boundaries and will be dragged with them towards the grain boundary. This results in an accumulation of impurity on the boundary. It is a nonequilibrium process and so if the system is allowed to equilibrate by heating at an intermediate temperature the segregating impurity atoms will diffuse back down their concentration gradients to even out the segregation concentration profile surrounding the boundary. Another way of looking at this is to represent the return to equilibrium conditions by very slow quench rates. At the other end of the timescale, very fast quench rates will not allow sufficient time for the impurity drag mechanism to accumulate large amounts of impurity on the boundary. This leads to the conclusion that there is a critical quench rate at which the maximum non-equilibrium segregation to the grain boundary is expected. This effect is displayed for Nb and Si in a German ferritic steel 1.4914 (see Table 11.1 for composition) in Fig. 11.3, showing the effect of quench rate and starting temperature on the impurity enrichment expected on the grain boundary (Faulkner, 1987).
11.2.4 Dynamic recrystallisation at grain boundaries Abe (2006) has observed recrystallisation to occur near to prior austenite boundaries in ferritic steels (see Fig. 11.4). He has ascribed this phenomenon to the onset of tertiary creep in such materials. Consequently, any elemental
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Si in 1.4914 1.6 1.4 1.2 1.0 0.8 0.6 0.4 0.2
1.6 1.4 1.2 1.0 F 0.8 0.6 0.4 0.2 15.0
9
Te m
14.5 pe 14.0 rat ure (K × 1 13.5 00 )
7 6
–1
5 (s te ra ing
8 ×1
0 00
)
4 3 13.0
1
2
Co
ol
Nb in 1.4914
7 6 5 4 3 2 1
7
F × 10
6 5 4 3 2 1 15.0
Te m
14.5
pe
rat
ure
14.0
(K
× 1 13.5 00 )
13.0
1.3 ) 1.1 00 0.9 1 × 10 – 0.7 (s te 0.5 ra ng 0.3 oli Co 0.1
11.3 Extent of grain boundary segregation, F, as a function of starting temperature and cooling rate for Si and Nb in 1.4914 (see Table 11.1 for composition) ferritic/martensitic steel.
addition or prior working treatment that can be applied to postpone the onset of this transformation will be valuable in improving the creep strength. B has been recognised as one element that can do this because B additions to P91,
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1 µm
11.4. Dynamic recrystallisation in modified 9Cr–1Mo ferritic steel tested at 100 MPa and 600°C for 34 141 h. After Abe (2006).
together with some adjustments to the W/Mo ratios to create P92, have resulted in improved creep strength. Presumably the effectiveness of the addition element must be contained in its ability to restrict grain boundary movement at high temperature. For this reason, elements and thermal treatments which produce excessive precipitate or segregation on grain boundaries could be useful in increasing creep life.
11.2.5 Effect of boron The ability of B to interact with diffusing atoms on grain boundary planes such that grain boundary diffusion is curtailed is one of the attractive features of this element. Abe (2006) has shown that grain boundary diffusion rates in an Fe–9Cr–3Co–3W alloy with small amounts of nitrogen are reduced, with the effect being manifested in a reduced M23X6 precipitation rate, and a corresponding increase in creep life is obtained. Boron additionally assists in retarding the formation of fine grain size regions in welds (Abe, 2006). These are responsible for type IV cracking and when B is added to the above steel, the occurrence of type IV cracking can be eliminated. This effect is presumably connected with the influence of boron on the dynamic recrystallisation in the boundary neighbourhood mentioned in the previous section. To complete the boron story, the intragranular precipitation may be
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stabilised against coarsening by limitation of diffusion between coarsening particles (Golpayegani et al., 2003).
11.2.6 Ductile to brittle transition temperature DBTT One of the important distinctions between ferritic steels and austenitic steels is the difference in their crystal structure. Ferritic steels possess the body centred cubic crystal structure, whereas the austenitic steels are face centred cubic. Body centred cubic systems have fewer slip systems available and so at lower temperatures their capability to deform plastically becomes severely limited. This is represented on a curve of yield strength, σy, versus temperature as shown in Fig. 11.5. Reviewing the temperature dependence of the cleavage or grain boundary failure strength indicates a line with a much lower slope, so that the picture for grain boundary fracture stress, óF, as opposed to matrix fracture is summarised as in Fig. 11.5. If segregation, dynamic recrystallisation, or any change in grain boundary structure occurs, this will act to reduce the level of the grain boundary fracture line in Fig. 11.5. The temperature at which the crossover between matrix yield and grain boundary or cleavage fracture is known as the ductile to brittle transition temperature (DBTT). Any change in the grain boundary state which lowers the grain boundary line or any strengthening that increases the matrix strength will increase the DBTT. To explain more fully, an increase in the matrix yield stress, ∆σy, will cause a shift in the DBTT of ∆Ty, assuming Matrix yield stress Grain boundary fracture stress
Stress
∆σy
∆σF ∆TF ∆Ty
Temperature
11.5 Stress-temperature diagram, showing the contributions of cleavage, grain boundary fracture stress and matrix failure stress to the ductile to brittle transition temperature (DBTT) shift. An increase in the matrix yield stress, ∆σy, will cause a shift in the DBTT of ∆Ty, assuming that the grain boundary fracture stress remains the same. If the grain boundary fracture stress decreases by ∆σF, while the matrix yield stress remains the same, then the DBTT shift will be ∆TF. If both the matrix yield stress and grain boundary fracture stress change simultaneously, the total DBTT shift will be ∆Ty + ∆TF.
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that the grain boundary fracture stress remains the same. If the grain boundary fracture stress decreases by ∆σF, while the matrix yield stress remains the same, then the DBTT shift will be ∆TF. If both the matrix yield stress and grain boundary fracture stress change simultaneously, the total DBTT shift will be ∆Ty + ∆TF. Materials possessing a low DBTT are more favourable because the grain boundary failure is commonly associated with fast, uncontrollable fracture. If the DBTT is above room temperature, it is generally recognised that the material will be unsafe. Therefore it is especially important for creep-resistant ferritic steels to possess a low DBTT as well as a high creep strength. The most important factor controlling the grain boundary strength line in Fig. 11.5 is the segregation. Elements like S and P present in quantities greater than about 0.01% combined in an ferritic alloy with no strong carbideforming elements like Ta or Nb are likely to have low grain boundary strengths. The grain boundary structure is also important: high angle random grain boundaries being the worst. The total contributions from all these effects have recently been calculated and summed in a theory of grain boundary fracture (Faulkner, 2005). Thus, although DBTT is not a creep-dependent property, it is a property that must be taken into consideration for a creepresistant ferritic steel and there are a set of criteria which will enable the steel designer to decide whether a chosen creep-resistant ferritic steel will be fit for purpose in a creep application.
11.2.7 Hafnium One of the success stories of modelling of creep-resistant ferritic steels is the forecast that was made about the beneficial properties of hafnium. The addition of this element in quantities of up to 1% has been shown to convey resistance to intergranular segregation effects concerning phosphorus and chromium (Lu et al., 2006). Furthermore the Hf addition forms HfC preferentially to chromium carbide on the grain boundaries (and in the grains), thus replacing a relatively coarsening-prone phase by a very stable grain boundary precipitate phase. The initial work has been done on E911 (Yin and Faulkner, 2005 (see Table 11.1 for composition). The net result is to prolong the time at which the microstructure will remain stable at a given temperature, thus producing longer creep life or similar creep life at higher temperatures (see Fig. 11.6). Currently alloy development is underway to test the predictions of the modelling.
11.2.8 Oxide dispersion strengthening Attempts to improve the strength of ferritic steels at high temperature usually become unsuccessful when the temperature rises above 625°C. This is because
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M2N
Strain, ε (%)
80
VN
339
HfC
50 years
60
40
20
0 0
200 000
400 000 600 000 Time, t (h)
800 000
11.6 Model predictions for creep strain rate using various phases as particle strengtheners. Microstructural predictions have been linked with the continuum damage mechanics models to produce this figure. Alloy P91 at 600°C.
the strengthening components in the alloys become insignificant above this temperature. The poor oxidation resistance above this temperature should also be borne in mind. However, if the application is in protected circumstances, like vacuum operation or if coatings are being considered, it should be possible to produce a ferritic steel that will have good creep strength up to 800°C. The means of achieving this objective are through oxide dispersion strengthening. The most common ferritic alloy is MA957, a 14%Cr ferritic steel (see Table 11.1 for composition) with a 0.25 wt% yttria dispersion. More recently the international fusion reactor programme has developed a special ferritic 9%Cr oxide dispersion strengthened (ODS) steel called Eurofer ODS (see Table 11.1 for composition) with microstructure and properties similar to the 9–12%Cr steels discussed so far in this chapter, but with creep resistance up to 800°C. The microstructure consists of the typical tempered martensite grain structure with yttria particles situated on intragranular sites. Studies are being made of the grain boundary structure in this material (Lu and Faulkner, 2007) and the field emission gun scanning electron microscope (FEGSEM) with electron back scatter diffraction (EBSD) is proving an invaluable technique in determining the real grain size and the distribution of grain boundary misorientations as a function of processing parameters. For example, there is a clear effect of quench rate from the solution treatment temperature on the percentage of low angle boundaries seen in Eurofer97 ODS (see Fig. 11.7). The prevalence of low angle boundaries provides more ductility in general yielding and this leads to lower DBTTs, which, as discussed earlier, is an obvious design advantage for these brands of ferritic steel.
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30 µm (a)
WQ AC FC
0.25
Number fraction
340
0.20
0.15
0.10
0.05
0.00 10
20 30 40 50 Misorientation angle (degrees) (b)
60
11.7 (a) EBSD picture and grain boundary misorientation plot of Eurofer97 ODS as a function of cooling rate from the solution treatment temperature. Note the lack of low angle boundaries in material that has been furnace cooled (FC).
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11.3
341
Austenitic steels
11.3.1 General Referring to Fig. 11.1, the main difference between the microstructures of austenitic and ferritic steels is that the remanents of the martensitic transformation are absent. This is because the Ni stabilises the austenite phase at all temperatures. The only grain boundaries are the so-called prior austenite boundaries of the ferritic microstructure. Other differences between the two types of steel are that austenitic steels possess a face centred cubic crystal structure and this results in a more compact lattice than that of the ferritic steel. The consequences of this are reduced diffusion rates and thermal conductivity, and higher thermal expansion coefficients. The higher solubility of carbon in austenite means that there is a greater potential for forming carbides but this is offset because the coarsening rates of such carbides are much greater than those contained within a ferritric matrix. Consequently carbide strengthening for creep purposes must rely on other phase strengthening mechanisms than carbide precipitation. Nevertheless the precipitation hardening component of creep life is still high because of the lower diffusion rates and austenitic steels can usually operate in creep environments with higher temperatures than those typically allowed for ferritic steels.
11.3.2 Precipitation The main type of grain boundary precipitate in austenitic steels is M23C6. TiC and NbC can form in higher temperature windows than those typical of service life for the austenitics. TiC will predominate on grain boundaries in the 800–1000°C regime, and at temperatures from 1000–1300°C, NbC will preferentially form. A typical set of isothermal precipitation curves fo NbC and M23C6 in a typical stainless steel based on the austenitic composition (see Table 11.1) are shown in Fig. 11.8 (after Faulkner, 1979). These results come from earlier modelling exercises based on analytical calculations of grain boundary precipitate behaviour (Carolan and Faulkner, 1988). Examples of grain boundary precipitation in austenitic steel are given in Fig. 11.9.
11.3.3 Segregation The austenitic steel composition has probably received more attention in relation to intergranular segregation than any other except perhaps Cu–Bi. Studies of boron segregation have been explored theoretically and autoradiographically by Williams et al. (1976) and by atom probe field ion microscopy by Karlsson and Norden (1988). Boron clearly segregates and it is thought that the main mechanism is non-equilibrium segregation. This is demonstrated by a series
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Precipitate size = 2.5 × 100nm diameter NbC (0.1 vol%C)
1450
1200 NbC (0.01 vol%C)
Temperature (°C)
1100
1350
1000 1250 900
(0.1 vol% M23C6 (0.05 vol% C) C)
1150 M23C6(0.01 vol%C)
800
1050
700 600
Theory Experimental points from Adamson, 1972 10
100 Time (s)
Temperature (K)
342
950
1000
10 000
11.8 Isothermal precipitation curves for 20C/25Ni austenitic steel (see Table 11.1 for composition), showing the ranges of temperature over which NbC and M23C6 are expected.
of autoradiographs showing B grain boundary segregation as a function of starting temperature and cooling rate (Fig. 11.10). The trends shown strongly support the proposed mechanism of non-equilibrium segregation discussed in the segregation section applied to ferritic steels in that the segregation enrichment passes through a maximum at an intermediate cooling rate. Another interesting grain boundary effect seen in austenitic steels concerns the segregation behaviour of chromium. If sensitised in the 600–700°C regime then chromium depletion will occur caused by the grain boundary carbide precipitation. If air cooled and not subsequently aged, chromium segregates to the boundaries (Flewitt and Vorlicek, 1993). The dependence of the depletion effect on grain boundary structure has been most intensively studied by Laws and Goodhew, (1991), who looked at 51 boundaries in the transmission electron microscope after sensitising Type 316 steel. Generally those boundaries with low fitting, high ∑ values, lead to wide depletion zones and, generally, lower minimum chromium concentrations ∑ = 3,11,13a,13b and 29a have narrow depletion zones and higher minimum boundary chromium concentrations. The depletion is also correlated with the extent of chromium carbide precipitation on the boundary.
11.3.4 Boron, hafnium and zirconium There is evidence (Hattestrand and Andren, 2001) that creep-resistant austenitic steels like Type 316 (see Table 11.1 for an alloy of similar
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0.5 µm (a)
(b)
11.9 Examples of intergranular precipitation in austenitic steels. (a) M23C6 and M6C in Type 304 LN (see Table 11.1 for composition), after 28 840 h at 650°C (after Vodarek et al., 1998); (b) TiC in Nimonic PE16 (see Table 11.1 for composition), after 1 min at 825°C (after Faulkner and Caisley, 1977).
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Ti = 1150°C
Ti = 1250°C (b)
(a)
Ti = 1200°C
(c)
Ti = 1200°C (d)
(e)
11.10 Autoradiographs of 316 LN steel (see Table 11.1 for composition) showing the effects of solution treatment temperature (a) 1350°C, (b) 1250°C, (c) 1150°C, and cooling rate from the solution treatment temperature of 1200°C at (d) 500 s–1 and (e) 5000 s–1 on B intergranular segregation. WPNL2204
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Ti = 1350°C
Grain boundaries in creep-resistant steels
345
composition) are improved if B is added. The reason for this is that the diffusion of chromium between grain boundary carbides is limited by the presence of boron in the neighbourhood of the boundaries. This reduces the coarsening rate with a resultant improvement in long-term creep life. Heavy elements like Zr and Hf are thought to limit grain boundary sliding during high temperature deformation. The most evidence for this is found in superalloy compositions.
11.4
Grain boundary properties and constitutive creep design equations
Constitutive equations exist for forecasting the damage accumulation in a variety of forms, and features connected with the grain boundary properties are extremely important parameters. The creep rupture life is usually calculated on the basis of the Hull–Rimmer equation, with adjustments made by Raj and Ashby (1975). These models calculate the vacancy flux to grain boundaries and elucidate how rapidly a given grain size structure will cavitate at inclusions to the point, or time tr, that cavity linkage will occur and the material will fail by grain boundary decohesion. The Raj and Ashby equation is given by:
tr =
FV ( α ) 3 π kT ∞ 3/2 32 ΩδDb σ ρ FB3/2 ( α )
∫
Amax
Amin
dA f ( A)
[11.1]
where σ∞ is the applied tensile stress, δDb is the grain boundary diffusivity, k is Boltzmann’s constant, T is the absolute temperature, α is the angle that defines the shape of the cavity (the angle between the tangent to the cavity and the grain boundary plane), FB is the shape factor determining the area of void at the inclusion–matrix interface, FV is a similar shape factor determining the volume of the void adjacent to the inclusion, Ω is the molar volume and ρ is the number of voids per unit area of boundary. The problem is in determining the integral involving the fraction of cavitated boundaries needed to cause fracture in the Raj expression, equation 11.1. It is found to be dependent upon whether the supply of vacancies is constant or limited. Therefore a more understandable form of predictive equation remains the Hull–Rimmer formulation: tr =
kT η 3 ΩδDb σ F
[11.2]
where k is the Boltzmann constant, T is the absolute temperature, η is the interparticle or cavity nucleation site spacing on the grain boundary, Ω is the atomic volume, δ is the grain boundary width, Db is the grain boundary selfdiffusion coefficient and σF is the applied stress. In both of the approaches,
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the grain boundary properties like diffusion, width and grain boundary nucleant spacing are all central to the time to rupture problem solution. Furthermore, grain boundaries feature strongly in damage-based parametric equations for secondary creep rate, dε/dt, where ε is strain and t is time. Traditionally a Monkman–Grant type analysis is used. This can be used to relate the minimum secondary creep rate to the rupture life. Abe (2006) uses an expression of the following form:
tr =
const ε˙ d ln ε˙ min dε
[11.3]
where ε˙ min is the minimum creep rate and the differential is the acceleration of creep rate in the acceleration creep regime. There are several versions of the equation depending upon which regime of temperature–stress space is occupied. The grain boundary properties dominate in the high temperature–low stress domain. When in this region Coble creep operates and the working equation commonly used is as follows:
ε˙ min =
ADb 4 σ kTd 3
[11.4]
where πδDb D = Dl 1 + dDl
[11.5]
where Dl is the lattice diffusion coefficient and Db is the grain boundary diffusion coefficient. As before, δ is the grain boundary width, b is the operational dislocation Burgers vector, σ is the applied stress, k is the Boltzmann constant, T is the absolute temperature and d is the grain size. Clearly the grain size and grain boundary diffusion coefficient must be known in order to forecast properly the minimum creep rate in operating conditions where Coble creep is dominant.
11.5
Future trends
As mentioned earlier it is normally accepted that fine grain sizes lead to poor creep properties because there is more opportunity for grain boundary sliding and cavity nucleation. This is the normally accepted reason for making single crystal superalloys for high creep strength applications in gas turbines. There is however evidence that the normal grain boundary controlled creep behaviour is altered when very small, nanosized grains are present. Under these circumstances the nature of the grain boundary becomes much less planar and the intergrain boundary spacing becomes less than the average dislocation pile-up length. Both of these factors could have a large effect on the ability
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of the materials to deform around the boundary plane. But there is in principle a larger effect concerned with the very high density of triple junctions expected in nanocrystalline materials. Galina et al. (1987), have shown that when the topological class of boundaries, according to the Von Neumann–Mullins criterion, is less than 6, the triple junction mobility reaches a rather low equilibrium state, implying that there is a strong limit to grain growth in a grain structure characterised by the n<6 criterion when the density of triple junctions is high, that is, in a nanocrystalline material. This means that many nanocrystalline materials could possess the kind of microstructure that would contain triple junction drag up to high temperatures and thus produce higher creep strength. Ultra-fine grain size steels are currently being explored from this viewpoint. Reference has already been made to the effects of Hf on the potential creep properties of ferritic high creep strength alloys. Further work is in progress exploring the beneficial properties of Hf in austenitic steels and superalloys. In all these types of materials, Hf seems to slow down migration kinetics in the neighbourhood of the boundary and to produce more stable carbide and nitride phases on the prior austenite and lath boundaries. Both of these characteristics lead to higher creep strengths. The removal of chromiumrich phases and substitution of hafnium-rich phases on the grain boundaries furthermore reduces chromium depletion effects that commonly lead to accelerated intergranular oxidation at high temperatures. This means that Hf will reduce the intergranular corrosion susceptibility in many austenitic steels. Finally, the modelling of grain boundaries and associated precipitation, deformation and segregation effects is set to undergo a considerable transformation in the near future. Molecular dynamics (MD) modelling, which goes back to first principles associated with the interatomic potentials, is becoming more powerful with the advent of parallel processing computers and the use of tricks to extend the timescale over which MD calculations operate. This means that atom movements will be simulated from first principles calculations to predict the approach of a non-equilibrium material to its equilibrium state. This, with a superimposed stress, is effectively a description of the creep process. It is anticipated that soon all analytical equation approaches to describing creep will be superseded by the MD simulation method. It is already being used successfully to predict the effects of neutron irradiation damage and dislocation–obstacle interaction. High temperature microstructural evolution and the associated creep property prediction could well be the next major application of the MD method.
11.6
References
Abe F. (2006), ‘Metallurgy of long term stabilisation of ferritic steels for thick section boiler components in USC power plant at 650°C’, Proceedings Conference on ‘Materials
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for Advanced Power Engineering, J. Lecompte-Beckers, M. Carton, F. Schubert and P.J. Ennis (eds), Scriften des Forschungszentrum Juhlich, Germany, 965. Adamson J.P. (1972), Precipitation Behaviour in Austenitic Steels, D.Phil Thesis, Oxford University. Aust K.T., Armijo S.J., Koch E.F. and Westbrook J.A. (1967), ‘Vacancy-driven grain boundary segregation’, Trans. ASM, 60, 360. Bruemmer S.M. (1999), ‘Grain boundary composition and its effects on environmental degradation’, Mater. Sci. Forum., 294–296, 75. Carolan R.A. and Faulkner R.G. (1988), ‘Grain boundary precipitation in alloy 800’, Acta. Metall, 36, 257. Czyrska-Filemonowicz A., Bryla K., Spiradek-Hahn K., Firganek H., Zieliunska-Lipiec A. and Ennis P.J. (2003), ‘Role of boron in 9% cr steels for steam power plant’, Proceedings Parsons 2003: Engineering Issues in Turbine Machinery, Power Plant and Renewables, Strang A., Conroy R.D., Banks W.M., Blackler M., Leggett J., McColvin G.M., Simpson S., Smith M., Starr F. and Vanstone R.W. (eds), Maney, London, p. 365. Faulkner R.G. (1979), ‘Inter-granular precipitation in austenitic alloys’, J. Mater. Sci. Lett., 14, 2249. Faulkner R.G. (1987), ‘Combined grain boundary equilibrium and non-equilibrium segregation in ferritic-martensitic steels’, Acta. Metall, 35, 2905. Faulkner R.G. (1996), ‘Segregation to boundaries and interfaces in solids’, Int. Mater. Rev., 41, 198. Faulkner R.G. (2005), ‘Grain boundary segregation and fracture’, Zeitschrift fur Metallk, 96, 1213. Faulkner R.G. and Caisley J. (1977) ‘Kinetics of grain boundary precipitation in nimonic PE16’, Metal Sci. J, 11, 200. Faulkner R.G., Yin Y., Cintas J. and Montes J.M. (2005), ‘Modelling and experimental studies of inter-granular corrosion in austenitic steels used in light water reactor systems’, 12th Conference on Environmental Degradation of Materials in Nuclear Power Systems, Allen T., King P. and Nelson L. (eds), Salt Lake City, TMS publication, Warrendale, PA, USA p. 135. Flewitt P.E.J. and Vorlicek V. (1993), ‘Cooling induced segregation of impurity elements to grain boundaries in Fe–3Ni alloys, 2 1/4Cr 1Mo steel, and submerged arc weld metal’, Nuclear Electric Report, TD/SEB/REP/2050/93. Galina A.V., Fradkov V.E. and Shvindlerman L.S. (1987), ‘Grain boundary mobility and its relation to topological class’, Phys. Met. Metalloved., 63, 165. Golpayegani A., Hattestrand M. and Andren H-O. (2003), ‘Effect of boron on precipitation, growth and coarsening in martensitic Cr steels’, Proceedings Parsons 2003, Strang A., Conroy R.D., Banks W.M., Blackler M., Leggett J., McColvin G.M., Simpson S., Smith M., Starr F. and Vanstone R.W. (eds), Maney, London, p. 347. Hattestrand M. and Andren H-O. (2001), ‘Influence of strain on precipitation reactions during creep of an advanced 9% Cr steel’, Acta Materialia, 49, 2123. Karlsson L. and Norden H. (1988), ‘Non-equilibrium segregation of boron in austenitic steels, Acta. Metall., 36, 13. Laws M.S. and Goodhew P.J. (1991), ‘Grain boundary structure and Cr segregation in a 316 stainless steel’, Acta. Metall., 39, 1525. Lu Z. and Faulkner R.G. (2008), to be published in J. Nucl. Mater. Lu Z., Faulkner R.G., Sakaguchi N., Kinoshita H., Takahashi H. and Flewitt P.E.J. (2006), ‘Effect of hafnium on radiation-induced segregation in ferritic steel’, J. Nuclear Mater., 351, 155.
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McLean D. (1957), Grain Boundaries in Metals, Oxford University Press, p.131. T.S. Morgan, E.A. Little, R.G. Faulkner and J.M. Titchmarsh (1992), ‘Interfacial segregation in fast reactor-irradiated 12% Cr martensitic steel’, Effects of Radiation on Materials: 15th International Symposium, Nashville, Stoller R.E., Kumar A.S. and Gelles D.S. (eds), (1992) ASTM STP 1125, p. 633. Newman R.C. (2001), ‘Environmentally assisted corrosion’, Corrosion, 57, 1030–1041. Palumbo G., LeHockey E.M. and Lin P. (1998), ‘Grain boundary engineering’, J. Metals, 50, 40. Raj R. and Ashby M.F. (1975), ‘Inter-granular fracture at elevated temperature’, Acta. Metall, 23, 653. Randal V. (2003), ‘Grain boundary engineering in IN718’, Scripta Metall., 44, 2789. Sun W.P., Militzer M. and Jonas J.J. (1992), ‘Grain boundary precipitation and high temperature deformation in steels’, Metall. Trans., 23A, 821. Suryaranayana C., Mukohopadhyay D., Patankar S.N. and Froes F.H. (1992), ‘The HallPetch relationship in nanocrystalline materials’, J. Mater. Res., 7, 25. Vodarek V., Sobotkova M. and Sobotkova J. (1998), ‘Effect of microstructural evolution on the creep rupture behaviour of CrNi (Mo)N austenitic steels’, Microstructural Stability of Creep Resistant Alloys for High Temperature Applications, A. Strang (ed.), Institute of Materials, London, p. 69. Williams T.M., Stoneham A.M. and Harries D.R. (1976), ‘Grain boundary segregation of boron in 316 steel’, Metal Sci., 10, 14. Yin Y. and Faulkner R.G. (2005), ‘Creep damage and grain boundary precipitation in power plant steels’, MST, 21, 1239.
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12 Fracture mechanism map and fundamental aspects of creep fracture K . M A R U YA M A , Tohoku University, Japan
12.1
Introduction
Suppose a structural component made of steel is loaded at room temperature. The component deforms elastically or plastically depending on the stress level, but the deformation under constant load (or stress) always stops within a short period of time. At elevated temperature, on the other hand, the component deforms continuously even under constant stress, and finally breaks after a certain period of time. This is the consequence of creep deformation and creep fracture. The time-dependent deformation and fracture are characteristics of high temperature deformation. Resistance to the deformation and fracture is an important property to be considered in the structural design of elevated temperature plants made of steel. The stress to cause fracture in 105 h is especially important, since it usually gives the allowable stress of steels. The value is found from short-term creep data by formulating creep data based on, for example, the Larson–Miller or Orr– Sherby–Dorn equations. The formulation should be done on the sound physical basis of creep fracture. This chapter deals with the physical basis of creep fracture. The creep fracture mechanism often changes from transgranular fracture in short-term creep to intergranular fracture in long-term creep. The changes in fracture mechanisms are summarized in fracture mechanism maps. The transition of fracture modes is often accompanied by a decrease in stress exponent and activation energy, namely breakdown of creep strength. Since the change in activation energy causes serious problems in the prediction of long-term properties, one should be careful about this change. This chapter will explain why and how creep fracture mechanisms and creep rupture properties change with creep testing conditions. Several examples are provided of how the changes bring about overestimation of the long-term rupture life. A multi region analysis of creep rupture data is proposed to prevent overestimation. 350 WPNL2204
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351
Fracture mechanisms and ductility of materials
A material exhibits several modes of fracture depending on creep testing conditions, namely stress and temperature. The fracture modes are classified into three types as shown in Fig. 12.1:1 (a) brittle intergranular fracture, (b) intergranular or transgranular fracture with some ductility, and (c) rupture after 100% reduction in area. The fracture ductility changes substantially with the fracture mechanisms. There are two types of brittle intergranular creep fracture (Fig.12.1 (a)). One is brought about by the coalescence of cavities formed on grain boundaries aligned perpendicular to the applied stress (see Section 12.3.1). The other is wedge cracking formed at triple grain boundary junctions by stress concentration that is created by grain boundary sliding. The intragranular creep cavities are nucleated at inclusions (solid circles) as shown in the lower part of Fig. 12.1 (b) and their coalescence by deformation brings about the final fracture. Recrystallization occurs in the necking part during the rupture (Fig. 12.1 (c)). Ashby and his co-workers have compiled fracture mechanism maps and Fig. 12.21 provides an example. The map can foretell the fracture mode operative under a given creep condition. The ductile fracture appearing at the high stress is typical of tensile deformation at room temperature. In Fig. 12.2, testing temperature is given as a fraction of melting temperature Tm. The rupture takes place at very high temperature only. Structural materials Brittle
Ductile σ
σ (a)
(b)
(c)
12.1 Representative fracture modes at high temperature. (a) Brittle intergranular fracture by wedge cracking and diffusional cavity growth, (b) intergranular and transgranular cavity growth due to plastic deformation and (c) rupture after 100% reduction of area. The solid circles in (b) represent inclusions.
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0.50 Tm
Nimonic 80A
Normalized stress, σ/E
Wedge crack Trans– granular fracture
0.59
10–3
0.62
W+C 0.66 Cavity
Rupture 10–4 100
0.53
102
0.78 Intergranular fracture 104 106 Rupture life, tr (s)
108
1010
12.2 Fracture mechanism map of Nimonic 80A. Creep stress σ is normalized with Young’s modulus E. Testing temperatures are given as a fraction of melting temperature Tm.
are usually used around 50% of their melting temperature, and transgranular creep fracture and intergranular creep fracture caused by the coalescence of cavities are important mechanisms under such creep conditions. The former appears at higher stress (shorter term creep) and the latter at lower stress (longer term creep). Wedge cracking is sometimes observed in the transition region between the two fracture mechanisms.
12.3
Stress and temperature dependence of rupture life
12.3.1 Cavity growth controlled by grain boundary diffusion The rupture life tr of a nickel based alloy is plotted in Fig. 12.2 as a function of creep stress σ. The rupture life is expressed as: tr = t0 σ–n exp (Q/RT)
[12.1]
where t0 is a material constant, n is the stress exponent, Q is the activation energy for rupture life, R is the universal gas constant and T is the absolute temperature. As evident in the figure, the stress exponent (n = dlntr/dlnσ) and activation energy for rupture life change with the fracture modes: from higher values of the transgranular fracture to low values of the intergranular
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fracture controlled by diffusional cavity growth. The different lengths of the horizontal arrows drawn in the figure confirm the decrease in activation energy. Section 12.4 will examine how the fracture modes change and why the change brings about decreases in the stress exponent and activation energy. This section is focused on the stress and temperature dependence of rupture life, and provides physical bases for the discussion in Section 12.4. A typical process of intergranular fracture consists of three stages: (1) nucleation of cavities at grain boundaries, (2) their growth and coalescence, resulting in a facet crack of single grain size, and (3) propagation of the crack to the final fracture. The rupture life tr of a material can be expressed as: tr = tn + tg + tp ≅ tg
[12.2]
where tn, tg and tp are the time taken for cavity nucleation, cavity growth and crack propagation, respectively. The second term is usually the longest in creep of heat-resistant steels and determines their rupture lives. Suppose creep cavities are nucleated on a grain boundary at an interspace of 2S (see Fig. 12.3), then the duration tg of cavity coalescence is given by: tg =
∫
S
(dr /dt ) –1 dr
[12.3]
r0
where r0 is the initial radius, r is the current radius of curvature of cavities and (dr/dt) is the growth rate of cavities. The cavity diameter is usually close to 2r. One can easily calculate rupture life, when (dr/dt) is known. Several formulae for the growth rate have been proposed to calculate tg and an example is given below. σ
2S Cavity
Cavity surface diffusion
Grain boundary diffusion
σ
12.3 Vacancy flow paths during diffusional growth of cavities on grain boundaries aligned perpendicular to the applied stress direction.
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Suppose we have cavities on a grain boundary aligned perpendicular to the applied stress (see Fig. 12.3). Concentrations of vacancies on the grain boundary and on the cavity surface with radius of curvature of r are given by: Cgb = Cv exp(Ω σ/k T)
[12.4]
Ccavity = Cv exp(2Γs Ω/r k T)
[12.5]
where Cv is the concentration of vacancies in matrix, Ω is the volume of a vacancy, Γs is the surface energy and k is Boltzmann’s constant. Since σ to be used in Equation [12.4] is the stress component normal to the grain boundary plane, vacancy concentration is highest on the boundaries aligned perpendicular to the applied stress. Therefore, creep cavities are more frequently formed on grain boundaries aligned normal to the applied stress direction. When σ > 2 Γs/r
[12.6]
vacancies flow from grain boundary (high vacancy concentration of Cgb) to cavity surface (low concentration of Ccavity) by diffusion through the grain boundary and then the cavity surface, resulting in cavity growth. When grain boundary diffusion controls the cavity growth rate, the growth rate is expressed as:2 dr/dt ∝ (Dgb δgb/r2)(Ω σ/kT)
[12.7]
where Dgb is the grain boundary diffusion coefficient and δgb is the thickness of the grain boundary zone. Substituting this equation into Equation [12.3], one can obtain the following equation describing stress and temperature dependence of rupture life: tr = tg ∝ (S3/Dogb δgb)(k T/Ω σ) exp(Qgb/R T)
[12.8]
where Dogb is a material constant and Qgb is the activation energy for grain boundary diffusion. This equation predicts a stress exponent of n = 1 and an activation energy of Q = Qgb.
12.3.2 Cavity growth controlled by cavity surface diffusion and by creep deformation Growth rates of grain boundary cavities have been formulated in other cases: cavity growth controlled by cavity surface diffusion and by creep deformation. Rupture lives of these cases can easily be derived by adopting a similar procedure to the one mentioned in Section 12.3.1. When the cavity surface diffusion controls cavity growth, the rupture life is given by:2,3 tr ∝ (S Γ s2 /Dos δs)(k T/Ω σ n ) exp(Qs/R T)
[12.9]
where Dos is a material constant, δs is the thickness of surface zone and Qs is the activation energy for surface diffusion. The stress exponent n for
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rupture life is 1.5 or 3 depending on the detailed assumptions about the cavity growth process. At higher stress or higher temperature, creep deformation of the matrix surrounding a cavity controls the growth rate of the cavity in grain boundaries. In such a case, the rupture life is given by:4 tr ∝ (1/ ε˙ ) ∝ (1/Dol σ n) exp(Ql/RT)
[12.10]
where ε˙ is the creep rate of the matrix, Dol is a material constant and Ql is the activation energy for lattice diffusion. The stress exponent n is greater than 3. The value of Ql is about twice those of Qgb and Qs. The growth rate of the wedge crack at triple grain boundary junctions is also controlled by creep deformation in the interior of grains and their rupture lives take the same form of Equation [12.10]. The same equation is applicable to transgranular fracture.
12.4
Fracture mechanism maps
As mentioned in Section 12.3, the growth rate of creep cavities is controlled by (1) grain boundary diffusion (2) cavity surface diffusion or (3) creep deformation in the body surrounding cavities. The stress and temperature dependences of the rupture life differ among cavity growth mechanisms as represented by Equations [12.8]–[12.10]. It can be postulated that the fastest mechanism among them, namely the one giving the shortest rupture life governs creep fracture of a material under a given creep condition. This assumption provides the stress and temperature dependences of rupture life depicted in Fig. 12.4 (a) and (b), respectively. The cavity growth mechanism with low stress exponent (cavity growth controlled by grain boundary diffusion, n = 1) appears at low stress, whereas the cavity growth mechanisms controlled by creep deformation within grains (cavity growth within grains (transgranular fracture), wedge cracking and grain boundary cavity growth controlled by creep deformation, n > 3) are dominant at high stress. The cavity growth controlled by cavity surface diffusion takes place in the medium stress range, and may be absent at high temperature.5 Since the activation energies for grain boundary diffusion Qgb and surface diffusion Qs are about a half that of lattice diffusion Ql, the cavity growth mechanisms controlled by these short-circuit diffusion are operative at low temperatures. As aforementioned, creep fracture theories predict a decrease in the stress exponent with decreasing stress (Fig. 12.4 (a)). The transgranular fracture (with a large n value) and intergranular fracture (with a low n value) should operate at high stress and low stress, respectively. The fracture mechanism map shown in Fig. 12.2 clearly confirms these predictions. Under a given stress, the fracture mechanism changes from transgranular fracture to intergranular fracture with decreasing testing temperature. The fracture
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Grain boundary cavity growth
n
log σ
By cavity surface diffusion GB cavity growth by creep deformation
By grain boundary diffusion
3–1.5
Wedge cracking
1
Transgranular fracture log tr (a)
Qs or Qgb
log tr
GB cavity growth by creep deformation
Grain boundary cavity growth
Wedge cracking
Transgranular fracture
By cavity surface diffusion
Q1 By grain boundary diffusion
1/ T (b)
12.4 (a) Stress and (b) temperature dependence of rupture life. The stress exponent and the activation energy are different depending on the fracture mechanisms that are operative.
mechanism change is accompanied by a decrease in the activation energy from Ql to Qgb or Qs (Fig. 12.4 (b)). The different length of the arrows in Fig. 12.2 confirms the change in activation energy.
12.5
Influence of fracture mechanism change on creep rupture strength
A set of creep rupture data for type 316 stainless steel is given in Fig. 12.5,6 together with its fracture mechanism map. Three regions with different n and
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Q values appear in the figure, and the dash–dot lines are the boundaries between the three regions: n = 9.7 and Q = 440 kJ mol–1 in the high stress region H, n = 5.8 and Q = 390 kJ mol–1 in the medium stress region M and n = 2.1 and Q = 280 kJ mol–1 in the low stress region L. In practical materials precipitates are often the major strengthener and the amount of precipitate usually decreases with increasing temperature. Therefore, the apparent activation energy for creep rupture life is often greater than the one predicted by creep fracture theories. This is the case in type 316 stainless steel. The activation energy for lattice self-diffusion in γ iron is Q1 = 285 kJ mol–1. The value of Q in region L (280 kJ mol–1) can be below Q1, if the apparent contribution of the microstructural change is eliminated. In Fig. 12.5 the stress exponent and the activation energy for rupture life decrease with increasing rupture life. This fact is consistent with the fracture mechanism map discussed in Section 12.4. The stress exponent is less than three in region L in addition to the low activation energy, suggesting intergranular fracture controlled by diffusional growth of grain boundary cavities. The stress exponents in regions M and H are greater than three, suggesting transgranular fracture or intergranular fracture caused by deformation controlled growth of grain boundary cavities. Fracture mechanisms of the steel are indicated by the letters T, W, C and σ in Fig. 12.5. Transgranular fracture (T) occurs in the high stress region, 500 Type 316 steel 600 °C
300
650 °C 200
H T M
Stress (MPa)
W C
700 °C 100 70
L
750 °C σ
50
30 20 102
103 104 Rupture life (h)
105
12.5 Fracture mechanism map for type 316 stainless steel together with its creep rupture data. T: transgranular fracture, W: wedge cracking, C: cavity nucleation at the grain boundary carbides, σ: cavity nucleation at the grain boundary σ phase.
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corresponding to region H. Wedge cracking (W, controlled by creep deformation) and grain boundary cavity formation at carbides (C) correspond to the medium stress range M. Grain boundary cavities are formed by σ phase particles and intergranular fracture occurs in the σ region. This region corresponds to the low stress region L. The changes in fracture modes coincide very well with the changes in n and Q values explained above. Grain boundary cavity is usually formed at the intersection of a particle surface and a grain boundary, and σ phase particles are responsible for cavity formation and the consequent appearance of region L. Figure 12.5 suggests that rupture life can be extended up to the extension of the dashed lines, if we can prevent the formation of the σ phase.
12.6
Influence of microstructural degradation on creep rupture strength
Advanced high Cr ferritic steels for boiler applications are normalized and then tempered at 750–800°C for a few hours and have a tempered martensitic lath structure. Figure 12.67 is an example of 9Cr–2W–0.4Mo–1Cu–VNb steel tempered at 770°C for 2 h. Vickers hardness values were measured in a grip portion of crept specimens at room temperature and they are plotted in Fig. 12.6 as a function of rupture life (in other words, aging time). The constant hardness for short-term exposure demonstrates that the microstructure is partly stabilized and remains as it is up to 4 × 103 h exposure at 650°C. Since the microstructure is not in equilibrium, agglomeration of precipitates and recovery of the lath structure become evident after long-term exposure. A drop in hardness starts at 4 × 103 h at 650°C and 600 h at 700°C. The softening curve can be expressed as: ta = to exp(–Hv/Hvo) exp(Qa/R T)
[12.11]
where ta is the aging time, to is a constant, Hv is the Vickers hardness, Hvo is a material constant, and Qa is the activation energy. The softening curves provides Qa = 295 kJ mol–1, being in agreement with the activation energy for lattice diffusion in the temperature range of interest (in ferromagnetic region). The dashed curve drawn in Fig. 12.6 (a) was predicted at 600°C based on the activation energy. Creep rupture data for the same steel are given in Fig. 12.6 (b).7 There are three regions with a different stress exponent n and an activation energy Q: n = 15 and Q = 775 kJ mol–1 in the high stress and short-term region H1, n = 10 and Q is the same in the medium stress and short-term region H2, and n = 5.3 and Q = 565 kJ mol–1 in the low stress and long-term region L. The different lengths of the horizontal arrows confirms the decrease in activation energy. The dash–dot line is the boundary between regions H and L. The decrease in n and Q namely the breakdown of creep strength in long-term
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(a)
60
220
0 °C
Hardness, Hv
240
600°C 650°C 700°C
200
Stress (MPa)
180
H1
200
H2 L 100 60 (b) 40 100
101
102 103 Rupture life (h)
104
105
12.6 (a) Vickers hardness in the grip portion and (b) creep rupture data for advanced high Cr ferritic steel (9Cr–2W–0.4Mo–1Cu–VNb steel). Downward arrows: onset of hardness drop at each temperature.
creep, is a serious problem in advanced high Cr ferritic steels. Several causes have been proposed to explain the breakdown: formation of the Z phase at the expense of MX precipitates, enhanced recovery of the lath structure along grain boundaries, a change from transgranular fracture to intergranular fracture, and so on. The downward arrows in the figure indicate the onset of hardness drop. There is a close correlation between the hardness drop and the breakdown of creep strength. A good correlation has been confirmed in several advanced high Cr ferritic steels.7 This fact suggests that the microstructural degradation can be a major cause of breakdown and the accompanying decrease in n and Q values. If we can postpone microstructural degradation (hardness drop) by some means, we may be able to improve creep strength in region L up to the dashed lines in Fig. 12.6 (b).
12.7
Change in creep rupture properties at athermal yield stress
As mentioned in Chapter 8, Section 8.4, instantaneous plastic deformation takes place upon loading when a material is creep tested above the athermal
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yield stress. Since instantaneous plastic deformation introduces dislocations into a specimen and alters its dislocation substructure, creep deformation behavior above the athermal yield stress should be different from that below it (see Chapter 8, Section 8.6). Creep tests are always performed below the athermal yield stress in the case of high strength steels such as advanced high Cr ferritic steels with a tempered martensitic lath structure. On the other hand, the creep test conditions for austenitic stainless steels are often above the athermal yield stress, since the steels are creep-tested after solution treatment. A substantial amount of plastic deformation takes palce upon loading above the athermal yield stress. Structural materials are used below the athermal yield stress in engineering plants. We should know what the changes in creep rupture properties are introduced at the athermal yield stress. In this section creep rupture data for 2.25Cr–1Mo steel are taken as an example. The steel was normalized at 930°C and then tempered at 720°C. Its creep deformation behavior is described in Chapter 8, Section 8.6, and creep rupture data for the steel are plotted in Fig. 12.7.8 The dotted line indicates the athermal yield stress level σa. The creep stress is normalized by Young’s modulus E, since σa is truly independent of the testing temperature when normalized by E.9 Creep tests were conducted both above and below the athermal yield stress. Usually the slope of the log tr versus log(σ/E) curve decreases with decreasing stress. However, it increases discontinuously below σa, indicating threshold-like behavior. The same threshold-like behavior of the creep rate is recognized (see Fig. 8.10). When the curves determined above σa are extrapolated to lower stress (dashed line), the extrapolation underestimates the creep rupture strength below σa. Therefore, we should
Normalized stress, σ/10–3 E
4 450°C 475°C 500°C
2
1 σa 0.6 0.4
0.2 0.1 101
525°C 550°C 600°C 650°C
103 Rupture life (h)
105
12.7 Creep rupture life of normalized and tempered 2.25Cr–1Mo steel. Creep stress σ is normalized by Young’s modulus E. The dotted line represents the athermal yield stress of the steel.
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carry out accelerated creep tests below the athermal yield stress in order correctly to evaluate long-term creep properties under service conditions. Despite the change in the stress exponent, the activation energy for rupture life does not change in this case: Q = 360 kJ mol–1 both above and below σa. The value is close to the activation energy for lattice diffusion of α iron in the ferromagnetic temperature range a little below the Curie temperature.
12.8
Multi-region analysis of creep rupture data
The long-term creep strength of a material is evaluated from short-term creep data by means of a time–temperature parameter (TTP) method, such as the Orr–Sherby–Dorn (OSD) and Larson–Miller parameters:10 log tr – Q log e/R T = f(σ)
[12.12]
(log tr + C ) T = f(σ)
[12.13]
where C is the Larson–Miller constant, f(σ) is a function of creep stress σ and e has its usual meaning. The parameters assume a linear relation between log tr and 1/T as shown in Fig. 12.8 (a). The conventional TTP analyses suppose that the temperature dependence of rupture life, in other words, the activation energy Q for rupture life or the Larson–Miller constant C is unchangd in the rupture data set. Therefore, they adopt the linear extrapolation and predict the rupture life given by the open circles in Fig. 12.8 (a). However, the activation (a)
(b) Q′ R σ3
II
ln tr
σ3
I σ2
σ1
T1
σ2
Q R Experimental data
T2
σ1
T1
T2
1/T
12.8 (a) The basic assumption of the TTP methods and (b) the decrease in activation energy from Q to Q ’ sometimes observed in real creep rupture data. The square symbols and the solid lines represent available experimental data, and circles and dotted lines are predictions.
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energy sometimes decreases to a smaller value of Q′ in long-term creep (Fig. 12.8 (b)), as demonstrated experimentally in Figs. 12.2, 12.5 and 12.6 (b). If this is the case, the open circles overestimate the true lives given by the solid circles in Fig. 12.8 (b). Such overestimation is a serious problem in austenitic stainless steels11 and advanced high Cr ferritic steels.12,13 Multi-region analysis of creep rupture data14 has been proposed to avoid overestimation of long-term creep properties. In the analysis, a set of creep rupture data is divided into several data sets, so that the value of Q is unique in each divided data set. Then conventional TTP analysis based on Equations [12.12] or [12.13] is applied to each divided data set. Creep rupture data for Gr.122 steel are given in Fig. 12.9.12 The data set is divided into two regions H and L at the boundary represented by the thin dash–dot curve. Multiregion analysis (OSD analysis of each divided data set) gives the regression curves represented by thick solid lines. The solid symbols in Fig. 12.9 were not used in the regression analyses. The activation energies are QH = 710 kJ mol–1 and QL = 370 kJ mol–1 in regions H and L, respectively. Conventional TTP analysis assumes a unique Q value for the whole data set. This conventional single regional analysis provides the regression curves represented by the dashed lines. The slope of the straight lines in Fig. 12.9 (b) corresponds to an activation energy of 620 kJ mol–1. The standard error of estimate in terms of log tr is 0.05 in multi-region analysis and 0.15 in single region analysis. In the multi-region analysis regression curves provide a better fit to the data points than those of conventional single region analysis. In the multi-region analysis regression curves describe very well the trends of data point in Fig. 12.9, and can predict the solid symbols adequately. This fact assures proper evaluation of long-term properties by multi-region analysis. On the other hand, the regression curve at 60 MPa in Fig. 12.9 (b) given by conventional single region analysis (dashed line) overestimates the solid symbol measured at 650°C. This is the overestimation often made by conventional TTP analyses. In the example given in Fig. 12.7 the activation energy does not obviously change up to 105 h. In this case one can correctly predict long-term creep rupture life from short-term data by conventional linear extrapolation.
12.9
Summary
Creep tests are carried out over a wide range of stress and temperature conditions and the creep rupture behavior of a material changes with the test conditions. For example stress exponent n and activation energy Q for rupture life decrease from high values for short-term creep to low values for longterm creep at low stress or low temperature. The change in Q values is often accompanied by a change in the fracture mechanism from transgranular fracture in short-term creep to intergranular fracture in long-term creep.
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400 600°C 625°C
Stress (MPa)
200 H 100 60 40 650°C 675°C 20 101
L
700°C
750°C 102
103 Rupture life (h) (a)
104
105
105
Rupture life (h)
QL = 370 kJ/mol
L
104
103
H
σ (MPa) 40
620 kJ/mol 102
QH = 710 kJ/mol
Gr. 122
60 90 140
1
10 0.95
1
1.05 1000/T (K–1) (b)
1.1
1.15
12.9 (a) Stress versus rupture life plot and (b) temperature dependence of rupture life of Gr.122 steel. Solid regression curves: multi-region analysis taking account of the change in activation energy Q between regions H and L. Dashed regression curves: single region analysis neglecting the change in Q.
Creep fracture theories can explain rationally the change in fracture mechanism and the consequent decrease in Q and n values. Of course, these changes may not appear in some materials within an appropriate range of test duration. Microstructural degradation during high temperature exposure is another probable cause of the decrease of Q in high Cr ferritic steels. The decreases in Q and n values result in a quick drop in creep strength and are recognized as breakdown of creep strength. Breakdown is a serious problem in austenitic stainless steels and advanced high Cr ferritic steels. When the onset of intergranular fracture is a cause of the breakdown, one
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can achieve significant improvement in creep strength by preventing the fracture mode from changing. Stabilization of the microstructure is a proper way to postpone breakdown caused by microstructural degradation. We usually evaluate long-term creep rupture life by linear extrapolation of a log tr versus 1/T relation assuming a constant Q value. The decrease in activation energy is the major cause of the overestimation of long-term creep strength evaluated by conventional time–temperature parameter methods. Multiregional analysis of creep rupture data is necessary to avoid overestimation. Creep rupture behavior above the athermal yield stress is different from that below it. We should carry out accelerated creep tests below the athermal yield stress to evaluate properly long-term creep strength under service conditions.
12.10 References 1 Ashby M F, Gandhi C and Taplin D M R (1979), ‘Fracture mechanism maps and their construction for F.C.C metals and alloys’, Acta Metall, 27, 699–729. 2 Nix W D, Yu K S, and Wang J S (1983), ‘The effects of segregation on the kinetics of intergranular cavity growth during creep condition’, Metall Trans, 14A, 563–70. 3 Chuang T J, Kagawa K I, Rice J R and Sills L B (1979), ‘Non-equilibrium models for diffusive cavitation of grain interfaces’, Acta Metall, 27, 265–84. 4 Beere W and Speight M V (1978), ‘Creep cavitation by vacancy diffusion in plastically deforming solid’, Metal Sci, 12, 172–6. 5 Cocks A C F and Ashby M F (1982), ‘On creep fracture by void growth’, Prog Mater Sci, 27, 189–244. 6 Nakakuki H, Maruyama K, Oikawa H and Yagi K (1995), ‘Collective evaluation of temperature and stress dependence of creep rupture life in austenitic stainless steels’, Tetsu-to-Hagane, 81, 220–224. 7 Ghassemi Armaki H, Maruyama K, Yoshizawa M and Igarashi M (2007), ‘Prediction of breakdown transition of creep strength in advanced high Cr ferritic steels by hardness measurement of aged structures at high temperature’, Key Eng Mater, 345– 346, 553–556. 8 Maruyama K, Kushima H and Watanabe T (1990), ‘Prediction of long term creep curve and rupture life of 2.25Cr–1Mo steel’, ISIJ Int, 30, 817–22. 9 Frost H J, and Ashby M F (1982), Deformation Mechanism Maps, Pergamon Press, Oxford. 10 Viswanathan R (1989), Damage Mechanisms and Life Assessment of High Temperature Components, ASM International Metals Park. 11 Maruyama K, Ghassemi Armaki H and Yoshimi K (2007), ‘Multiregion analysis of creep rupture data of 316 stainless steel’, Int J Press Vess Piping, 84, 171–176. 12 Maruyama K and Yoshimi K (2007), ‘Influence of data analysis method and allowable stress criterion on allowable stress of Gr.122 heat resistant steel’, Trans ASME, J Press Vess Technol, 129, (3), 449–453. 13 Maruyama K and Yoshimi K (2007), ‘Methodology of creep data analysis for advanced high Cr ferritic steel’, Proc of 8th int conf on Creep and fatigue at elevated temperatures, San Antonio, USA, 2007, ASME, New York, Paper No. 26510, 1–6. 14 Maruyama K, Baba E, Yokokawa K, Kushima H and Yagi K (1994), ‘Errors of creep rupture life extrapolated by time-temperature parameter methods’, Tetsu-to-Hagane, 80, 336–341.
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13 Mechanisms of creep deformation in steel W . B L U M, University of Erlangen-Nuernberg, Germany
13.1
Introduction
The mechanisms of creep can be grouped into diffusive flow (matter transport by atomic diffusion), sliding on high-angle grain boundaries and crystallographic slip. As the latter is dominant in producing strain under most conditions of creep,1 we will focus attention on this mechanism. Crystallographic slip is conveniently described by the motion of dislocations, representing the borderlines of slipped regions. Owing to their stress fields, dislocations interact with each other, with solute atoms and with particles of foreign phases. These interactions are the clue to understanding creep mechanisms. To a good approximation it is sufficient to describe the microstructure by frequency distributions and the mean values of its characteristic spacings, in particular the spacings between grain and subgrain boundaries, free dislocations inside the subgrains, dislocations in low-angle boundaries, solute atoms and particles along dislocation lines. Even such a simplified view of the microstructure leads to considerable complexity. It should be taken before working on the lower scale of discrete dislocations with much higher computational effort. While complementing observations on various length scales from the atomic to the macroscopic level are necessary to obtain a consistent picture, the microstructural level with spatial averages of characteristic spacings is indispensable to achieving a basic understanding of the observed creep behavior and providing a basis for a phenomenological description of creep properties. The present treatment begins with the initial state of steels after production under typical conditions (Section 13.2). The fundamental difference between austenitic and ferritic steels with regard to the solid state phase transformation from austenite to ferrite leads to significant differences between their initial structures. These are reflected in the variation of creep resistance with strain in monotonic creep at constant stress as displayed in the creep rate–strain curves (Section 13.3). Section 13.4 summarizes information on creep mechanisms which can be obtained from the transient material response to 365 WPNL2204
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sudden changes in stress. This information is also useful in understanding the acceleration and deceleration of creep resulting from cyclic variation of load (Section 13.5). In Section 13.6, observations on creep behavior presented in the preceding sections are given a semi-quantitative microstructural interpretation. Section 13.7 reports on the quantitative description of creep by dislocation models coupling laws of deformation kinetics and structure evolution. Section 13.8 compares the model to in situ observations of dislocation processes during straining in a transmission electron microscope (TEM) at elevated temperature. The prospects for dislocation models in providing a quantitative fit of the observed macroscopic creep properties on a microstructural basis and in guiding alloy development are briefly discussed in Section 13.9.
13.2
Initial microstructure
13.2.1 Austenitic steels Austenitic steels do not undergo a phase transformation during their production. Their initial structure is typical of a recrystallized material with a low content of dislocations within the grains. The high solubility of interstitial elements in the matrix reduces the tendency for precipitation of carbonitrides. However, depending on the concentrations of the interstitial elements C and N, there is heterogeneous precipitation of carbonitrides, preferably at grain boundaries and dislocations.2,3 High contents of alloying elements like Cr and Ni in the matrix lead to solid solution strengthening.
13.2.2 Ferritic steels During cooling from the melt, the matrix of ferritic steels transforms from austenite to ferrite. The transformation is accompanied by a strong reduction in solubility of foreign atoms. The martensitic transformation is a particularly important case generating a class of tempered martensite steels. It comprises cooperative shearing of the crystal lattice (see e.g. Wayman4 and Haasen5) and thus causes intensive local deformation of the matrix, resulting in strong work hardening due to a cellular dislocation structure with high dislocation density. Tempering of the martensitic structure leads to precipitation of solute atoms and to recovery of the dislocation cell structure6–8 resulting in a subgrain structure, characterized by the frequency distributions of misorientations and of boundary spacings to be determined along test lines. The subgrains are bounded by the boundaries of the prior austenite grains, of blocks of martensite laths of similar orientation, and of martensite laths and of subgrains within the laths.9,10 The first two kinds of boundaries are of the high-angle type with misorientations across the boundaries generally
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lying above 10–15°;10–12 the latter ones are of the low-angle type with lower misorientations. In contrast to low-angle boundaries constituting planar dislocation networks, high-angle boundaries interrupt the coherency of the lattices of the neighboring crystallites, cannot in general be penetrated by dislocations, allow grains to slide relative to each other, and provide a particularly effective short circuit path for diffusion of atoms. The subgrain size w (mean linear intercept) yields the overall boundary area per volume, 2/w.13 If the subgrains are not equiaxed, one may differentiate between the mean subgrain intercepts w|| parallel and w⊥ perpendicular to the direction of subgrain elongation. It was found that 2/w ≈ 1/w|| + 1/w⊥.14 Other methods of quantifying the subgrain size, for instance, from the mean subgrain area, lead to results which differ to a limited extent from the mean subgrain intercept length w, and vary in proportion to it.15 It is known that w approaches a steady state value in the course of deformation which scales in inverse proportion to stress: w∞ = kw b G/σ
[13.1]
Here b is the length of the Burgers vector, G is the elastic shear modulus and kw is a numerical factor which for steels was reported to lie at about 10.16–18 A typical value of the initial subgrain size w0 in tempered martensite is 0.4 µm.6,17,19 According to Equation [13.1] this value corresponds to a stress level 380 MPa, representing an estimate of the maximal local stresses which have acted in the steels during martensitic transformation. Owing to the qualitative differences in the properties of high- and lowangle boundaries mentioned above, the areal fraction fhb = w/d of high-angle boundaries (d is the spacing of high-angle boundaries) in the subgrain structure of tempered martensites may be of importance. In the initial state fhb is too small to cause a significant influence of high-angle boundaries on the creep resistance. The situation may change, however, when the subgrain structure coarsens during creep at low σ so that fhb increases significantly along with w. The relatively low solubility of interstitial atoms in the ferritic grains drives precipitation of carbonitrides during and after martensitic transformation. Nucleation of most carbonitrides is heterogeneous and preferentially takes place at boundaries of relatively high misorientation.6,20 These precipitates grow to large sizes dp corresponding to the low number of competing nuclei. Some carbonitrides, in particular those of type MX (where M are metal and X interstitial atoms), also precipitate more homogeneously inside the subgrains, probably owing to lattice coherence facilitating nucleation. They remain distinctly smaller than the heterogeneous ones and therefore have a high strengthening potential in spite of their low volume fraction fp. This follows from the fact the strengthening potential scales in inverse proportion to the spacing of precipitates along dislocation lines which in turn is proportional to d p / f p for an isotropic distribution of particles (see e.g. Reppich21).
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Creep at constant stress
Creep tests require a specimen to be heated to test temperature T and the stress to be raised from zero to the intended final value. During loading the rate ε˙ of mechanical (elastic plus inelastic) strain is relatively high (typically ≈ 10–3 s–1). To sustain the applied stress the specimens work harden. The inelastic strain accumulated during loading is similar to that measured in tests at constant ε˙ = 10–3 s–1 when the flow stress has reached the level of creep stress. After loading the specimen creeps at constant stress. As the elastic strain remains constant, the creep strain is purely inelastic in nature. Continuing work hardening at constant stress causes ε˙ to decrease in the so-called primary stage of creep down to a minimal value ε˙ min in the secondary creep stage. As metals in general do not fracture in uniaxial compression, a tertiary stage with ε˙ increasing by tensile fracture, coupled with external and internal necking, does not occur here. Investigation of compressive creep is thus a means of identifying loss of creep resistance owing to global changes (coarsening) of the microstructure (see Section 13.3.2).
13.3.1 Austenitic steels Being relatively soft in their coarse-grained, recrystallized initial state, austenitic steels may display large inelastic loading strains in creep tests at elevated temperature. These may be suppressed by prior work hardening (increase in dislocation density) in predeformation.22,23 Figure 13.1 displays a set of creep rate ε˙ –strain ε curves for the alloy 800H at 1073 K. The creep stress of 156 MPa exceeds the yield stress of ≈ 100 MPa at ε˙ ≈ 10–3 s–1 and induces a loading strain ≈ 0.03 which hardens the material at an average rate ∆σ/∆ε ≈ 0.04 G, where G ≈ 50 GPa24 is the elastic shear modulus of alloy 800H at the test temperature. This rate is similar to the work hardening rate M2 × G/300 ≈ 0.03G expected for polycrystals (M = 3: the Taylor factor) from the work hardening rate of single crystals at Stage II, resulting from storage of dislocations in the absence of pronounced dynamic recovery.25,26 Subsequently the creep rate decreases steeply at an average rate | d ln ε˙ /dε| > 950. Using the stress exponent n = d log ε˙ /d log σ ≈ ∆log ε˙ /∆log σ ≈ 11 of the creep rate derived from Fig. 13.4, the work hardening rate at the beginning of creep at constant σ is estimated as dσ/dε = σ d ln σ/dε = (σ/n) d ln ε˙ /d ε˙ > 0.21 G, exceeding the former one by a factor > 5. Within a subsequent small strain interval ≈ 0.01, the work hardening rate abruptly drops to low values near zero before work hardening at a low rate < 0.004 G recommences. The sequence of an extraordinary relative maximum in work hardening rate immediately followed by a relative minimum at the beginning of primary
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10–3
10–4 156 10–5
ε˙ (s –1)
138 125 103
10–6
76 10–7 10–8 0.0
800H 1073 K
σ/MPa = 58 0.1
0.2 ε
0.3
0.4
13.1 Creep rate versus strain curves for alloy 800H (German designation X10NiCrAlTi32 20, composition in mass %: ≈ 32 Cr, ≈ 20 Ni, 0.1 C) in tension. From Portella.24
creep is even more pronounced for the austenitic Cr–Ni steel shown in Fig. 13.2; here the minimal work hardening rate even becomes negative so that a transient period of work softening develops before work hardening is resumed. These features can hardly be explained in terms of storage and annihilation of dislocations alone, which in pure materials lead to a comparatively slow, monotonic transition to the steady state.27 They will be discussed in Section 13.6.3 in terms of interaction between dislocations and foreign atoms. Beyond the extrema of work hardening rate, primary creep proceeds at a gradually decreasing rate and enters the secondary stage of creep at a minimum creep rate ε˙ min which is close to the rate of steady state creep, where work hardening and recovery of dislocations compensate and the dislocation structure is in a state of dynamic equilibrium. The tensile creep curves of Fig. 13.1 end with an extended tertiary stage where creep is strongly accelerating at increasing rate d log ε˙ /dε. As this stage is absent in compression, it must be related to tensile fracture.
13.3.2 Ferritic steels From their thermal-mechanical production history described above, ferritic steels have a relatively high yield stress owing to hardening by precipitates and dislocations and therefore do not generally need significant work hardening to carry the creep stress (Fig. 13.2). This is particularly pronounced in tempered martensites where hardening by fine dislocation structures is particularly strong. Here the primary stage of creep relatively quickly leads into the
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ε˙ (s –1)
X20 CrMoV 12 1 915 K, 235 MPa
10–4
X3 CrNi 18 9 923 K, 250 MPa 10–5 0.0
0.1
0.2
0.3
0.4
0.5
ε
13.2 Creep rate versus strain curve for tempered martensite 12Cr– 1Mo–V steel in uniaxial compression. Thin line: austenitic 18Cr–9Ni steel for comparison. From work by Hofmann,92 and his collegue Straub.
secondary stage with a relative minimum ε˙ min of creep rate which is immediately followed by a tertiary stage with increasing ε˙ . In contrast to the typical behavior of austenitic steels, a tertiary stage appears not only in tension, but also in compression of tempered martensites, and displays a rate d log ε˙ /dε of increase of log ε˙ that decreases with compressive strain ε (Fig. 13.2). As mentioned above, the reason for tertiary creep in compression is microstructural coarsening. When the coarsening rate has sufficiently declined in relation to the rate at which strain is progressing, the compressive creep rate ε˙ attains a constant steady state value at relatively large strains (Fig. 13.2), which by far exceeds the minimum creep rate ε˙ min . The latter must therefore be kept quite distinct from the steady state creep rate.28 It simply characterizes the maximum deformation resistance at the point where the effects of initial work hardening and softening by microstructural coarsening happen to compensate. In tension of tempered martensites the steady state of creep is generally not attained owing to interfering fracture processes.16
13.4
Transient response to stress changes
13.4.1 Small transient strains Sudden changes in stress during creep cause materials to respond by elastic (time-independent, reversible), anelastic (time-dependent, reversible) and
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plastic (time-dependent, irreversible) deformation. Figure 13.3 shows these reactions in a test on a tempered martensite steel where the uniaxial compressive stress σ was reduced in steps. As creep decelerates with decreasing σ, a logarithmic time scale was chosen to display the progress of ε. Each change ∆σ of the initial stress σ0 to a new level σ = σ0 + ∆σ leads to an instantaneous change in elastic strain ∆εel. If ∆εel is measured free from machine contributions, the ratio E′ = ∆σ/∆εel equals the static elastic (Young’s) 0.1915 εmech,0
σ0 = 294 MPa
ε˙ 0 = 1.3 · 10–4/s
0.1910
εmech
0.1905
0.1900
0.1895
0.1890
300
σ (MPa)
250
200
150
100
50 102
103
104
105
(t – 5200 s) (s)
13.3 Mechanical strain εmech as function of time t for an experimental tempered martensite 12Cr–2W–5Co steel at 923 K in response to stepwise stress reduction from the initial state after creep for εmech,0 at stress σ0 to creep rate ε˙ 0 . Measurement in uniaxial compression on a specimen of 6 mm height and 5 mm width. Horizontal dashes mark end of elastic reaction to unloading. From Backes.31
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modulus Estat of the steel specimen. In the test for Fig. 13.3, E′ is lower than the dynamic modulus E of the steel by a factor of 2, which is probably mainly due to reversible machine contributions to the measured length changes. Tensile tests on long specimens are better suited for the correct determination of Estat than tests on short compression specimens (see e.g. An et al.).29 The plastic response to a stress reduction | ∆σ | < σ0 is forward straining at a rate depending on acting stress σ and the evolving structure. For small | ∆σ |, the creep rate immediately after stress reduction is positive as plastic strain accumulation dominates. This allows one to measure the σ-dependence of the plastic deformation rate at nearly constant dislocation structure. As the latter is directly proportional to the average velocity of gliding dislocations, these tests give access to the dislocation glide velocity at the given structure. The anelastic response to sudden stress changes is forward or back straining for stress increases and stress reductions, respectively, mainly owing to reversible motion of dislocations towards or away from the obstacles exerting a back stress on them. At the lowest σ-levels shown in Fig. 13.3 the dominance of anelastic back straining shortly after the stress reduction is clearly seen. One notices that the anelastic reaction is smaller than the elastic reaction (this holds also when the strain signal is corrected for elastic machine contributions). At a certain intermediate value of stress reduction, plastic forward strain and anelastic back strain exactly compensate for some time and the average inelastic (anelastic plus plastic) strain rate is zero immediately after the stress reduction. This stress reduction is an estimate of the effective stress σ* for glide needed for modeling (see Section 13.7.1). However, the test in Fig. 13.3 is not optimal for this task, as the constancy of the microstructure in the sequence of stress reductions is not guaranteed. It was designed to measure the (positive) creep rate ε˙ at approximately constant total strain ε as function of σ; this function allows one to calculate the stress relaxation behavior.30,31 Stress reductions for estimating σ*32–34 must be performed when a defined microstructural state with a certain deformation resistance, given by the rate ε˙ 0 of plastic deformation at the stress σ0, is established. Therefore such stress reductions have to start from a point where σ = σ0 and ε˙ = ε˙ 0 .
13.4.2 Large transient strains When creep is monitored for large strain intervals after a σ-change, structural changes following the change in σ become visible in the changes of ε˙ with ε as the dislocation structure evolves towards the new steady state corresponding to the new stress. Figures 13.4 and 13.5 show that ε˙ varies in a non-monotonic manner with ε after large stress reductions and, for 800H, also after large stress increases. The transients after stress reduction, with work hardening followed by work softening, resemble the initial primary creep reaction of
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10–3 800H 1073 K 10–4 σ/MPa = 10–5
ε˙ (s –1)
123
10–6 80 10–7
10–8 0.0
0.1
0.2
0.3
ε
13.4 Transient ε˙ –ε response of alloy 800H to sudden reductions in tensile stress at 1073 K. Thin lines: reference curves at constant stress. After Portella.24
10–4 X20 CrMoV 12 1 873 K
ε˙ (s –1)
10–5 σ/MPa = 196 10–6 150
10–7
10–8 0.00
0.05
0.10
0.15 ε
0.20
0.25
0.30
13.5 Transient ε˙ –ε response of tempered martensite 12 C–1Mo steel to changes in compressive stress. Thin lines: reference curves at constant stress. From Straub.17
austenites described above. In addition, one sees that the steels creep faster at the higher stress after an intermediate period of slow creep over a relatively long time and, analogously, more slowly at the lower stress, when creep time is saved by an intermediate period of relatively fast creep at the higher stress. This indicates that there are also time-dependent softening processes occurring in both steels during creep.
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Creep-resistant steels
Cyclic creep
The progress of inelastic strain under cyclic variation of stress is called cyclic creep. In a simple case a time period ∆tc of cyclic creep consists of two phases where a maximum stress σ and a reduced stress Rσ (0 ≤ R ≤ 1) are acting in time intervals ∆tc,σ and ∆tc – ∆tc,σ > 0, respectively. Owing to the strong stress dependence of creep rate, progress of creep at significantly reduced stress is often negligible, so that the net creep rate per cycle may be estimated to be reduced by the ratio ∆tc,σ/∆tc. However, this is not generally found. Rather, one observes cyclic creep which is accelerated or decelerated compared to the naive expectation.
13.5.1 Cyclic creep deceleration An example of repeated changes of creep mode of alloy 800H from monotonic creep at constant stress σ to cyclic creep and back is presented in Fig. 13.6. For the last range of cyclic creep with the largest period ∆tc = 250 s the creep strain at σ in the time interval ∆tc,σ = 0.5 ∆tc is expected to be 1.25 × 10–3 from a rate of 10–5 s–1 for monotonic creep just before the mode change. This is larger by a factor of about 1.5 than the change in elastic strain related to the stress change (taking E ≈ 150 GPa). The net cyclic creep rate shown in the figure drops by a factor of 2.1 compared to the level of monotonic creep rate. This is close to the naively expected time ratio ∆tc/∆tc,σ = 2. However, for the first range of cyclic creep with the smallest cycle period 10–3 1073 K 124 MPa
800H
10–4
R=1 1
ε˙ (s –1) 10–7
10–8 0.0
0.04 250 s
0.04 50 s
10–6
R = 0.04 ∆tc = 10 s 0.1
1
1
10–5
∆tcσ = ∆tc /2
0.2
0.3
ε
13.6 Transient ε˙ –ε response of alloy 800H to sudden change from monotonic creep (R = 1) to cyclic creep (R = 0.04, period ∆tc) at net rate ε˙ . From Portella.24
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∆tc = 10 s, where the plastic creep strain is expected to be only 5% of the elastic strain, the net cyclic creep rate drops by a factor of ≈ 100 compared to the preceding rate of monotonic creep instead of the expected factor of 2. This means that the plastic creep rate at σ under conditions of cyclically interrupted loading is much smaller than at constant stress σ. In addition, there is a pronounced transient response after the mode change which closely resembles the transient response following the first stress reduction shown in Fig. 13.4. This indicates microstructural changes after a sudden change from monotonic to cyclic creep, which are similar to those following stress reductions. We conclude that a change from monotonic to cyclic creep with intermittent unloading at sufficiently high frequency may act like a stress reduction. This phenomenon is known as cyclic creep deceleration.
13.5.2 Cyclic creep acceleration When the frequencies of cycling are relatively low so that the creep strain at maximum stress σ is larger than the elastic (and the anelastic) strain, creep tends to be accelerated by the cyclic unloading.35 One reason for acceleration is the recovery of the dislocation structure in the reduced stress phase which may cause significant acceleration of creep at full stress in materials with normal transient creep behavior,35,27 in particular in pure materials where normal creep transients are most pronounced and anelastic deformation is relatively fast, as solute friction is missing. Another obvious possible reason for cyclic creep acceleration lies in time-dependent softening processes.36–38 As described in Section 13.4, these processes tend to accelerate creep at the maximum stress σ. However, they also tend to decelerate it at the reduced stress, so that the acceleration is partially compensated. The creep cycles reflecting loading conditions of real components, for instance of power stations, are quite complicated. In a phenomenological description, the total creep acceleration factor results from a weighted sum of acceleration factors for each phase of cyclic creep.37 The experimental effort needed in determining the acceleration and weighting factors cycle by cycle is high; an accompanying modeling of cyclic creep on a microstructural basis would be most useful in saving costs and time.
13.6
Microstructural interpretation of creep rate
13.6.1 Basic equations The dominant process of strain generation in metals is crystallographic slip by glide of free dislocations. The average rate ε˙ M , at which the resolved shear strain increases in polycrystals, is directly proportional to the product of density (length per volume) ρf and average glide velocity v of free
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dislocations, equalling the slipped area per volume and time (Orowan equation):
ε˙ M = b ρf v
[13.2]
(Note: The term ‘free dislocations’ covers all dislocations lying in the interior of subgrains, regardless of their mobility. The free dislocation structure is also addressed as the three-dimensional Frank network.) The obstacles to dislocation glide are found in the dislocation structure itself, in solute atoms and in hard particles of foreign phase. It is common to assume linear superposition of the strengthening components σp from particles and σ from dislocations: σ = σ p + σρ
[13.3]
based on the idea that particles lead to a general reduction in the stress available to move dislocations in the matrix between the particles. Following classical approaches (see Seeger39 and Nes40) the obstacles posed by the dislocation structure itself, for example forest dislocations, are assumed to consist of two parts, an athermal one, which can only be overcome by mechanical force, and a thermal peak which is susceptible to thermal activation. The athermal part requires a stress component σG,ρ, depending on temperature only through the elastic shear modulus G(T). The thermal part requires an additive thermal stress component σ*, also called the effective stress, to be overcome at a certain speed. Assuming linear superposition one obtains: σρ = σG,ρ + σ*
[13.4]
σG,ρ may be written in the simplest form as the sum of contributions from free dislocations (subscript f) and dislocations forming low-angle boundaries (subscript b): σG,ρ = σG,f + σG,b
[13.5]
σG,f varies in inverse proportion to the average spacing ρf–0.5 of free dislocations: σG,f = Cf α M b G ρf–0.5
[13.6]
α ≈ 0.3 is the dislocation interaction constant. The factor Cf increases from 0 to 1 as the dislocations move past their obstacles when the stress is applied. Similarly σG,b must be built up by motion of free dislocations against the subgrain boundaries.
13.6.2 Solid solution hardening The velocity v at which dislocations glide over the thermal obstacles depends on σ* and the nature of the obstacles. It is a monotonic function of σ* with
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v = 0 at σ* = 0. Within a sufficiently limited interval of temperature and stress the v(σ*) relation can be well approximated by a power law: v = B σ*m
Qg B = B0 exp – RT
[13.7]
with effective stress exponent m, apparent activation energy Qg for glide and a constant B0. It is important to note that the function v(σ*) differs depending on whether or not the dislocations are surrounded by a cloud of solute atoms during glide. Such clouds with enhanced solute concentration tend to form around the (edge) dislocations, as solutes are attracted by the dislocation stress field. When the dislocations start gliding, the clouds of solutes exert a drag on the dislocations. The solute drag causes glide of dislocations to be viscous with m = 1.41 The expression for the parameter B in Equation [13.7] derived by Cottrell and Jaswon42 for interaction between solute atoms and dislocations on the basis of the relative atomic size misfit εa of solutes reads.41,43
B=
9 kBT D ⋅ sol M G b 4 ln ( r2 / r1 ) c 0 ε 2a 2
[13.8]
where kB is Boltzmann’s constant and r1 and r2 are the internal and external cut-off radii of dislocation stress field. c0 and Dsol are the atomic concentration and the diffusion coefficient of the cloud of solutes respectively (see Chapter 7 for details). The temperature dependence of Dsol makes a determining contribution to Qg. If a cloud of solutes does not form or if dislocations break away from their clouds, glide occurs in a jerky manner where fixed obstacles, formed by dislocations in combination with solutes, are overcome after a certain waiting time with support by thermal activation. The exponent m is >> 1 in this case. The hardening effect of solutes is strong, if the factor B in Equation [13.7] is low so that the dislocations move slowly. As seen from Equation [13.8], solute atoms with both high concentration in the matrix and strong interaction with dislocations, for instance by large atomic misfit, are strong hardeners. In the case of viscous glide, a low diffusivity Dsol is an additional condition for strong hardening. Mo atoms turn out to be particularly effective solid solution hardeners of steel. However, Equations [13.7] and [13.2] show that the effect of solid solution hardening becomes small under conditions of slow long-term creep, because the effective stress diminishes in direct proportion to the dimishing dislocation velocity v. Therefore the long-term creep resistance must be guaranteed by particle hardening rather than by solute hardening. The schematical Fig. 13.7 illustrates the fact that jerky glide is faster than viscous glide at high σ* and correspondingly low T. The break away of the dislocations from their clouds at a certain maximum of
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effective stress for viscous glide is connected to the increase in v. The formation of clouds at a certain minimum value of effective stress for jerky glide leads to a decrease in v. Processes related to the formation of and break away from clouds are known by the term dynamic strain aging. Far from being a peculiarity, this addresses a ubiquitous phenomenon in the deformation of alloys. The manifold of different interstitial and substitutional alloying elements and the fact that the activation energies of creep and diffusion of solutes often are similar produces dynamic strain aging effects in a wide range of temperatures and strain rates. Creep tests are particularly well suited to revealing the effects of dynamic strain aging, because the creep rate for obvious reasons is much more sensitive to changes in the thermal stress component σ* than the flow stress. With this qualitative knowledge of v(σ*) and using Equations [13.2] to [13.7] it is possible to gain a qualitative understanding of the processes during creep by considering the variations in the microstructural parameters σf, ρb and σρ with strain.
Jerky 1
log ν
2 Break-away
Thinning of clouds
Cloud formation 4 Viscous
3
log σ ∗
13.7 Dislocation velocity v as function of effective stress σ* for jerky and viscous glide at given T, with jerky ↔ viscous transition (schematic); upward arrow: viscously gliding dislocations break away from their solute clouds; downward arrow: jerkily moving dislocations slow down owing to formation of a solute cloud.
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13.6.3 Loading strain and initial primary creep Austenitic steels with low initial dislocation density Primary creep transients of the type described in Section 13.3.1 have also been found in highly alloyed Ni–Cr base alloys29 and in Ti alloys.44,45 Their interpretation in terms of dynamic strain aging applies to steels as well. A simple interpretation neglecting the frequency distribution of the dislocation velocity and cooperative effects in formation and break away from clouds is as follows: During work hardening in the loading phase where stress increases, few dislocations move at high velocity v in the jerky mode without cloud (point 1 in Fig. 13.7). Increase of dislocation density leads to reduction of v at approximately constant ε˙ (see Equation [13.2]) without change of mode (point 2). When the increase of flow stress stops and creep begins, continuing work hardening makes σ* and v fall. This enables solute atoms to attach themselves to the dislocations, form a cloud and significantly reduce v to the level of viscous glide (point 3). The abrupt decrease in v reduces the rate of generation of dislocations so strongly that the rate of recovery of dislocations, which as a diffusion-controlled process is rather insensitive to solute atoms, now exceeds the generation rate. This leads to coarsening of the dislocation structure, decrease of σG,ρ and increase of σ* and v, until the steady state density of free dislocations is eventually reached at point 4. This consideration explains the pronounced relative maximum of apparent work hardening rate (see Section 13.2.1) as a direct consequence of the jerky–viscous transition (from point 2 to point 3). The subsequent minimum in work hardening rate is due to the recovery-induced transition from point 3 to 4. The present discussion of dynamic strain aging was restricted to interaction between dislocations and solutes by solute drag. It may be noted that short range ordering, which has been reported, for example for Cr in Fe (see Schönfeld),46 may cause similar effects; the destruction of short-range order by fast glide corresponds to the break away phenomenon; the rebuilding of short range order at low creep rate corresponds to cloud formation. If the interpretation of the beginning of primary creep in terms of dynamic strain is correct, the effective stress for viscous glide must be smaller than the available effective stress. This condition can be checked. Taking literature data for Ni,29,47 combining Equations [13.2], [13.7] and [13.8], and approximating ρf by (σ/(Gb))2 yields σ*/σ = 0.6 for creep at 76 MPa. This is a reasonable value in view of the numerous uncertainties and approximations involved. The effective stresses σ *v , max for break away from clouds were reported by Oikawa and Langdon41 to range between 10–4 and 10–3 G. This means that the highest creep stress ≈ 150 MPa = 3 × 10–3 G is close to the viscous–jerky transition. Thus the proposal of dynamic strain aging in creep of alloy 800H is viable.
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The preceding consideration does not cover the slow work hardening in the extended primary creep stage preceding the steady state of creep (Fig. 13.1). This work hardening is easily explained in the usual manner by gradual formation of a subgrain structure which is superimposed on the evolution of free dislocations and has been observed, for instance, in a Ni–Cr alloy,27,29,48 which is similar in behavior to alloy 800. In addition to solid solution strengthening, the strengthening by precipitates and dispersed particles reduces the primary creep rate.24 Coarsening of the precipitates with time increases the spacing of precipitates along dislocation lines and causes gradual loss of particle strengthening. This process may superimpose on the ones described above. Ferritic steels with high initial dislocation density The processes described in the preceding section tend to be suppressed in ferritic steels, which in their initial state generally are more strongly hardened by precipitates and dislocations inherited from phase transformation and in which the contents of alloying elements in solid solution tend to be smaller. Often the dislocation density need not increase during loading, as the initial density already exceeds the steady state density under creep conditions,6,17,49 and the loading strain is negligible. A primary stage of creep with decrease of ε˙ is nevertheless observed (Figs. 13.2 and 13.5). The reason is that the athermal dislocation hardening term σG,ρ is low at the beginning of creep. This can be explained in the following way: The deformation paths differ in phase transformation and creep. The high initial dislocation content inherited from martensitic transformation comprises a significant fraction of dislocations which are easily glissile under the combined influences of inherited internal stresses and applied creep stress. These dislocations glide past obstacles. As they do so, the forces exerted by the gliding dislocations on the obstacles rise. In consequence, the factor Cf in Equation [13.6] rises up to its maximum value of 1. Assuming a total length per volume, ρ, of dislocations to move by a distance in the order of an average dislocation spacing ρ–0.5, before being stopped at obstacles in the dislocation structure, means a strain of (b/M) ρ ρ–0.5 (see Equation [13.2]), amounting to 0.002 if the total dislocation density is estimated from Equation [13.6] by setting σG,f equal to the estimate of the local stress 380 MPa acting during martensitic transformation (see Section 13.2.2). This rough estimate lies in the order of magnitude of the observed primary creep strains. It is concluded that the fast creep in the primary stage can be explained in terms of build-up of dislocation interactions opposing the applied stress. The description of primary creep given above resembles the idea that sites with easy strain generation are gradually exhausted (exhaustion creep).50
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The theory of exhaustion creep predicts a decrease of creep rate with time t according to ε˙ ∝ t–p. p = 2/3 corresponds to the empirical Andrade creep law.51,52 This type of time-dependent decay of primary creep rate is in fact observed in steels.52–55
13.6.4 Transient creep response to changes in creep conditions Stress changes The interpretation of primary creep transients in terms of dynamic strain aging given in Section 13.6.3 can be transferred to the transients of Figs. 13.4 and 13.5 following abrupt changes in stress. In this interpretation the initial work hardening (decrease of ε˙ with ε) after stress reduction is connected to the densification of the clouds of solutes leading to significant reduction in dislocation velocity. In the subsequent range of work softening (increase of ε˙ ) the dislocation structure coarsens until a new balance between generation of dislocations in the course of glide and annihilation of dislocations in the course of recovery is re-established. Analogously, the reaction after stress increase is explained by thinning of the clouds and refinement of the dislocation structure. In addition to dynamic strain aging, recovery of dislocations will initially add to the observed decrease of ε˙ .56,57 This is seen from the following consideration. The effective stress was determined by stress reduction tests to be about 25 MPa in creep of 800H at 123 MPa.24 This means that forward glide associated with generation of dislocations ceases after the stress reduction in Fig. 13.4, because σ* is no longer positive. In this situation creep strain is produced as a by-product of recovery, for instance by migrating subgrain boundaries.58 However, the strain rate associated with recovery falls quickly as recovery slows down and glide of free dislocations becomes dominant again. For tempered martensite steels the situation is similar to that discussed for austenites with regard to solute effects and recovery-associated strain. An estimate similar to the one in Section 13.6.3 based on thermodynamic data47,59,60 yielded relative effective stresses σ*/σ = 0.37 for viscous drag of a cloud of Mo at 200 MPa and 873 K. The stress of 200 MPa corresponds to 3 × 10–3 G, which is again close to the break away of dislocations from clouds. A difference between Figs. 13.4 and 13.5 lies in the fact that the stress changes in the tempered martensite are performed in the tertiary stage of creep, in contrast to those in the austenitic alloy. In both cases the transients join the general trend of increase and decrease of ε˙ with ε. The fact that the transients keep a distance from the curves measured without stress changes has already been related to time-dependent softening (see Section 13.4).
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Loss of particle hardening offers itself as the most probable mechanism (see Section 13.6.6). Changes from monotonic to cyclic creep The intermediate unloading in each cycle of cyclic creep allows the dislocations to come to rest, apart from anelastic backward glide motion. In this situation more solutes have the chance to attach themselves to the dislocations. Thus, a change from monotonic to cyclic creep acts in a similar way as a limited stress reduction. The net creep rate therefore changes with strain in the mode change test of Fig. 13.6 in close similarity to the stress change test (Fig. 13.4); for instance, repeated unloading for 5 s followed by creep at full load for 5 s causes a transient creep reaction similar to lowering the stress from 124 MPa to 80 MPa.
13.6.5 Tertiary creep caused by subgrain coarsening In Section 13.6.3 formation and refinement of subgrains have been invoked to explain primary creep of steels with low initial dislocation density. Correspondingly, coarsening of a pre-existing subgrain structure will lead to softening, that is, to increase in creep rate with time, which phenomenologically is considered as tertiary creep, ending the secondary stage of creep. Subgrain coarsening occurs in creep when the initial subgrain size is smaller than the steady state subgrain size according to Equation [13.1]. This is the case in most applications of tempered martensite steels. Reduction of creep strength by subgrain coarsening was confirmed in tests where subgrains were made to coarsen within a relatively short time, where changes in the precipitate structure are negligible.61–63 The coarsening was achieved by strain-controlled cyclic straining at elevated temperature. When the maximum stress acting in cyclic deformation is sufficiently low, the subgrains grow fast with accumulating inelastic strain towards the instantaneous stress-dependent value of the steady state subgrain size (Equation [13.1]). The subgrain coarsening to 1.5 µm caused the minimum creep rate to increase by about an order of magnitude.61,62 Softening caused by subgrain coarsening, shows up in the tertiary stage of creep of tempered martensite steels in tension as well as in compression.17 In contrast to austenitic steels, the softening begins at low strains, as soon as the initial primary creep related to build-up of forces on the obstacles (see Section 13.6.3) has ceased. Subgrain strengthening is extremely important for the relatively high creep resistance of the tempered martensites not only by reducing the minimum creep rate compared to the steady state creep rate, but also by shifting the minimum to low strains. These two advantages counteract the disadvantage of ferritic steels compared to conventional
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austenitic ones regarding higher diffusivity, causing higher rates of structural coarsening and creep. It must be noted that subgrain strengthening is not a pure dislocation effect. The energy stored in fine structures provides a high driving force for structural coarsening. It needs the pinning effect of precipitates to stabilize the subgrain structure of tempered martensites. This was convincingly demonstrated by an experiment by Kostka et al.64 These authors produced a subgrain structure in a ferritic Fe–Cr alloy by severe plastic predeformation in equal channel angular pressing followed by annealing. The initial subgrain structure was quite similar to that of tempered martensites. But the lack of precipitates at the subgrain boundaries caused fast coarsening of the subgrain structure by dynamic recrystallization during creep. In effect, the unstabilized subgrain structure induced by severe plastic predeformation led to a creep curve at 100 MPa and 873 K where the relative minimum in ε˙ had nearly disappeared so that the actual value of the minimum creep rate had increased by a factor of 104 compared to the tempered martensite. This result clearly shows the essential effect of structural stabilization by particles. The particles reduce the velocity of migration of the subgrain boundaries and thus inhibit the formation of stable nuclei for recrystallization. The subgrain structure in tempered martensite steels is therefore particle stabilized.6
13.6.6 Increase of creep rate caused by degradation of particle hardening Degradation of particle hardening caused by an increase in the mean spacing of particles along dislocations is unavoidable at elevated temperatures. It generally occurs by Ostwald ripening of the existing precipitates. During ripening the average precipitate volume increases linearly with time t spent at elevated temperature T: 3 d p3 = d p,0 + kp t
[13.9]
so that the spacing of precipitates along dislocations increases and particle hardening is diminished.16,17 This time-dependent softening counteracts the work hardening in primary creep and contributes to the increase of creep rate in the tertiary stage of creep.16,17 The growth constant kp is proportional to the specific energy of the phase boundary of precipitates and depends on the atoms forming the precipitates. Atoms with a high concentration in the precipitates, a low concentration in the matrix and low diffusivity in the matrix are useful in reducing kp.65–67 The time effects on creep rate observed in stress change tests (Section 13.4) are consistent with Ostwald ripening. An example of relatively fast Ostwald ripening is given by the phase M2X in tempered martensite steels. This phase precipitates in a fine
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distribution within the volume, hardens the subgrain interior and provides a high short-term creep resistance. However, it also coarsens rapidly, preferably at subgrain boundaries, while the fine precipitates inside the subgrains dissolve. The accompanying loss of precipitation hardening of subgrains leads to relatively fast degradation of creep resistance.63,68,69 The ripening of precipitates becomes more complicated when a newly nucleating phase like the Z-phase disturbs the diffusion currents set up between the precipitates. A stability advantage of the new phase may cause existing phases to dissolve in favor of the new phase. The result may be similar to fast ripening of existing phases and lead to fast degradation of an initially high creep resistance. It is therefore of importance to discourage the nucleation of more stable phases. This is difficult to achieve, as there are no simple recipes for inhibiting formation of more stable phases, because heterogeneous nucleation is a local event depending on the properties of the individual nucleation sites. The effect of coarsening of the precipitate structure may degrade the creep resistance of an initially creep-resistant steel to an extent that its longterm creep resistance is lower than that of a steel with relatively low initial creep resistance, but more stable particles.68,69
13.6.7 Deceleration and acceleration of creep by cyclic variation of stress The large reduction in net creep rate caused by cyclic unloading at high frequency (Fig. 13.6) cannot be explained by dislocation–solute interaction alone. It needs an explanation in terms of anelastic deformation.36,70,71 We recall that cyclic creep is decelerated when the expected plastic strain in the phase of maximum stress is small compared to the elastic strain. In this case the anelastic strain may be small, too, and may have been interrupted before it reached its final value corresponding to the applied stress. Anelastic glide motion of dislocations is necessary to build up the forces on the dislocation obstacles which are necessary for their irreversible overcoming producing plastic deformation. This is described by the factor Cf in the example of Equation [13.6] and holds for other obstacles like particles as well. Interruption of anelastic deformation means that the forces on the obstacles are lower than in monotonic deformation. This explains why the plastic deformation rate may be strongly reduced by cyclic variation of stress at high frequency. Acceleration of creep by cyclic stressing has nothing to do with anelasticity, but is explained by structural coarsening in the unloading phases as already mentioned in Section 13.5.2.
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385
Dislocation models of creep
13.7.1 Kinetics of dislocation glide An important step in development of a microstructural model of creep sketched in Section 13.6.1 is the formulation of the dislocation glide velocity v as function of the effective stress σ* at given T. This requires an educated guess, as neither v nor σ* can be fixed exactly. By definition of effective stress, v(σ*) must vanish at σ* = 0. Using this condition, σ* can in principle be determined in tests where σ is suddenly reduced to a level where the strain increment ∆ε within a small time increment ∆t is zero. For technical materials like steels this is comparatively easy. The stresses and strains are relatively high so that their changes become relatively large. This eases the problems34 encountered with pure materials where the strain changes may be too small to be resolved. However, the problem of systematic errors, caused, for instance, by strains associated with recovery, remains. Nevertheless the stress reduction technique allows the order of magnitude of σ*/σ to be fixed and thus facilitates reasonable model assumptions. As an example, Fig. 13.8 shows the results of a series of stress change tests for a tempered martensite steel. δε is the difference between the maximum back strain and the elastic back strain. For stress reductions ∆σ < 0.05 σ0 the difference δε was zero, that is, no net inelastic back flow was observed, indicating that the effective stress for glide remained ≥ 0 after the stress 0 X20 CrMoV 12
δε (10–4)
–1
–2
873 K σ 0 = 320 MPa –3 0.5
0.6
0.7 0.8 (σ 0 – ∆σ )/σ 0
0.9
1.0
13.8 Difference δε between maximum back strain and elastic back strain as a function of relative reduced stress for the tempered martensite steel from Fig. 13.2. Stress reduction tests were performed starting from σ0 at strains 0.035 < ε0 < 0.44 and initial creep rates of 4 × 10–5 s < ε0 < 4 × 10–4 s in uniaxial compression. The estimated experimental error of individual δε values is ±10–4. After Goblirsch.67
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reduction. For stress reductions ∆σ > 0.24 σ0 the inelastic back flow |δε| has grown to 2 × 10–4 indicating net back flow under σ* < 0. A straightforward conclusion is that the effective stress at σ = σ0 should lie between 0.05 and 0.24 σ0 in the investigated case. However, Fig. 13.8 shows that this conclusion is reliable only, if the uncertainty in δε due to uncertainty in strain measurement and systematic errors is distinctly less than 10–4. If creep rate ε˙ , dislocation density ρf and effective stress σ* are known, it is possible to derive the dislocation velocity v as function of effective stress σ* (Equation [13.7]) from Equation [13.2]. A complete set of data was provided by Milička for alloy 800 at 975 K.22 It yields the v(σ*) relation shown in Fig. 13.9. The effective stress exponent m of v lies near 3. This value corresponds to the range of the curve in Fig. 13.7 where the solute cloud is thinning with increasing v before break away occurs. Wolf and coworkers29,53 performed a similar analysis for a Ni–Cr alloy. Their result, also shown in Fig. 13.9, is similar to that for alloy 800, if the difference in T is neglected. Polcik49 proposed the curve of Fig. 13.9 for a tempered martensite steel at 923 K. The lower and the upper branch of this curve were interpreted in terms of viscous and jerky glide, respectively. 10–6
10–7
ν (ms–1)
10–8
10–9
NiC
r22
Co
–10
10
lo Al 10–11
9-1
0 12M
y8
2%
0
Cr
73 , 10
9 0,
ste
75
, els
K
K
87
3K
10–12 10
100 σ* (MPa)
13.9 Dislocation velocity v as a function of effective stress σ* for alloy 800 derived from data of Milička,22 for typical 9–12%Cr steels49 and for a Ni-base alloy.53,29
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The overall consistency of the data displayed in Fig. 13.9 is encouraging. Even though no definite conclusion is possible at present regarding the exact form of the v(σ*) relation, investigations of the kind described are a useful guideline in fitting the v(σ*) relation, needed for modeling, to experimental data.
13.7.2 Evolution of dislocation structure It is desirable to model the evolution of dislocation structure on the basis of expressions for the rates of generation and loss of dislocations. Significant progress has been made in developing these expressions.40,43,72,73 However, there have been only few efforts in the field of steel (see e.g. Weinert)74 and a widespread direct application of the structure evolution approach to steel is still lacking. Fortunately, the exact system of differential equations for dislocation structure evolution can be approximated by a system which is based on experimental knowledge of the steady state dislocation structure.17,18 Knowing the initial values X0 and the steady state values X∞ of the characteristic spacings X, the evolution of the dislocation structure with strain can be approximated by integrating rate equations of the type dX = – X – X ∞ [13.10] dε kX with adequately chosen values of kX. At constant stress and temperature, Equation [13.10] predicts X to approach X∞ in an exponential fashion. This phenomenological law of structure evolution has successfully been applied to tempered martensite steels.16–18,49 It was shown to be applicable not only for creep, but also for subgrain size evolution in strain controlled cyclic deformation at elevated temperature.62,69 The modeling is simplified by the fact that the steady state values of the spacings X = w and ρf–0.5 can be approximated as unique functions of normalized stress σ/G and that the mean spacing s of dislocations in subgrain boundaries of low-angle type is approximately constant.17 The latter approximation is probably not valid for steels without pre-existing subgrain structures and should be modified according to proposals for strain controlled decrease of s (increase of average boundary misorientation b/s).75
13.7.3 Particle hardening As mentioned before, particle hardening and stabilization of the dislocation structure by particles may have a decisive influence on the creep resistance. In order to model creep of particle-hardened steels one needs to know the mechanism of interaction between precipitates and dislocations and the evolution of the particle structure during creep.
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The simplest approach to interaction stress σp between free dislocations and particles is to assume that the dislocations by-pass the particles by glide. In this case σp takes its maximum value given by the Orowan stress. As seen in Section 13.8, the assumption that σp is near to its maximum value during deformation at elevated temperature is supported by electron microscopic studies. The Orowan stress can be calculated for particles which are relatively homogeneously distributed in the volume, like MX-precipitates. However, it does not consider the effect of climb over particles during slow creep. In case of climb over particles, the interaction stress may well be given by the attractive interaction between particles and free dislocations.76 In this case σp would be smaller than the Orowan stress and might also have a thermal component. The stabilization of low-angle boundaries (dislocation networks) against migration under stress must be described in a different manner. Polcik49 used an approach based on Zener drag resulting from attractive interaction between boundaries and particles, which is another form of attractive dislocation– particle interaction. At present there is no safe procedure for modeling precipitation hardening. Therefore it is necessary to rely on fitting modeling approaches containing particle structure parameters to experimental results. An example will be given below (see Appendix). In any case the coarsening of the particle structure and the related softening is a fact which needs description. The parameters entering the description are the locations of the particles (at boundaries or distributed over the volume) and the sizes dp and volume fractions fp of particles. As different phases behave differently, a phase specific description of evolution is necessary. The simplest approach is to assume that particles evolve via diffusive processes with time. At the present state of knowledge this appears to be justified.68 In spite of repeated claims (e.g. Eggeler77 and Cerri et al.78) a clear influence of concurrent creep straining on ripening of particles has not been unequivocally proven.68,79 However, there are indications that particle ripening is enhanced at low-angle boundaries where short-circuit diffusion along dislocations may be important.6,68,69,79 Despite the complexity of the situation with different phases at different locations, which may be enhanced by the appearance of new phases, characterization of the particle structure appears possible on the basis of experimental observations combined with the knowledge provided by thermodynamic data banks.67,80 The next step leading from the particle structure to expressions for particle hardening needs educated guesses and fitting.
13.7.4 Composite model As an example of a statistical dislocation model of creep of steels, we consider the (iso-strain) composite model of plastic deformation proposed by Mughrabi,81
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expanded for inclusion of strain rates by Blum and co-workers82–84 and applied to steel by Straub and co-workers85 and Polcik49 and to Ni-phase alloys with dominant solute hardening by Meier and Blum.48 In this model the material is viewed as a composite of soft subgrains surrounded by hard boundaries. The equations used in fitting the model to a tempered martensite steel are given in the Appendix. The composite model comprises: • • • • •
dislocation structure evolution, particle structure evolution for individual precipitate phases in the subgrain interior and at the subgrain boundaries, hardening by free dislocations and subgrain boundaries, particle hardening of subgrain interiors and boundaries, solute hardening influencing the dislocation glide velocity.
In the case of inhomogeneous formation of subgrains in solid solution hardened fcc metals, the deformation resistances of the volume fractions with and without subgrain structure were averaged using an iso-stress composite approach.48 The composite model applied to tempered martensite steels provides ε˙ – ε curves which reproduce experimentally observed features in that the creep rate reaches a minimum with strain and then increases owing to straindependent coarsening of the dislocation structure and time-dependent coarsening of the particle structure.49,85,86 There is a sufficient number of parameters in the model to fit the quantitative result to experimental observations. Anelastic deformation caused by interaction between soft subgrain interiors and hard boundaries is implicit in the model. The softening with both strain and time is reproduced by the model. The stress dependence of the minimum creep rate can also be reproduced to a reasonable extent. Little effort has been spent so far to improve the fit between model and experiment leaving opportunity for improvement.
13.8
In situ transmission electron microscope observations of dislocation activity
As mentioned in the introduction, creep is all about dislocation activity. One would therefore like to see dislocations moving in a steel under stress at elevated temperature. This is in principle possible by deforming specimens of steels inside a transmission electron microscope (TEM) and recording the motion of dislocations. Such in situ TEM observations have recently been made by Messerschmidt et al.87 on an experimental tempered martensite steel. The subgrain structure in this steel had been coarsened to w ≈ 1.5 µm by prior cyclic deformation at 873 K.62 This facilitated the observations. Figure 13.10 shows the interior and some of the boundaries of a subgrain. Most of the boundaries appear to be pinned by precipitates. A free dislocation
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13.10 In situ observations in TEM with 1 MV acceleration voltage of a thin area of an experimental tempered martensite 12Cr–2W–5Co steel during slow elongation at 923 K by Messerschmidt et al.87 reported by Chilukuru:69 subgrain with boundaries at top, bottom and right (see dislocation networks) and free dislocations bowed at precipitates (arrows). Still picture.
is held up at two precipitates. It is relatively strongly bowed. This means that the stress needed for the dislocation to overcome the precipitates in the subgrain interior is not far from the maximum given by the Orowan stress. Figure 13.11 illustrates why subgrain boundaries in tempered martensite steels are regarded as relatively hard regions with relatively high densities of dislocations and particles. The subgrain boundary in the lower half of the figure identifies itself by the linear black–white contrast (fringes of equal thickness) and the sharply delineated intersection with the foil surface. It is pinned by large precipitates; note the slight bowing of the boundary between precipitates. Below the boundary dislocations are out of contrast. Above the boundary bowed dislocation are seen which apparently come out of the boundary. This means that dislocations are able to cross boundaries by moving into it from one side and, some time later, leaving it towards the other side. This is in line with the notion that the mean free path of free dislocations in
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2 µm
13.11 As Fig. 13.10: free dislocations leave the subgrain boundary. Still picture.
subgrain structures is distinctly larger than the subgrain size, see Blum27 and Nes.40 It is unlikely that the boundary is completely dissolving, because the subgrain structure is already relatively coarse and essentially maintains its average boundary spacing during deformation at the elevated temperature.16,49,62,88 Figure 13.12 shows video frames of a subgrain taken at different times. Dislocation loops are generated at the bottom subgrain boundary. They expand into the subgrain. Some dislocations overtake others. Many dislocations are resting. These dislocations are relatively straight, probably because they feel only a low resolved shear stress acting on them. When bowed dislocations are released from strong obstacles, probably precipitates, they jump forward at high speed. In situ TEM observations can be used to test the validity of the microstructural model of creep. The lower part of the subgrain is crossed by about six dislocations within a time interval ∆t = 50 s. This means a shear strain of ∆γ = 6b/tfoil and a strain rate ε˙ = M–1 ∆γ/∆t ≈ 10–5 s–1 (tfoil ≈ 0.5 µm for 1 MV-TEM). This rate corresponds to a stress of about 174 MPa (data from Dubey et al.62 for steel 2A extrapolated with a stress exponent of 8). From the data given in the Appendix one expects an average spacing of free
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(a)
(b)
0.5 µm
(c)
(d)
13.12 In situ observation of configuration of gliding dislocations in subgrains of the steel in Fig. 13.10 (frames from video of Messerschmidt et al.)87: (a) dislocations have been emitted from a source in the subgrain boundary at the bottom of this TEM picture, (b) within less than 0.1 s from (a), dislocation 4 has jumped a distance of 0.2 µm, much faster than average, (c) 3 s later dislocation 5 is held up at an obstacle; bowing has reached a maximum, (d) within ≈ 0.1 s from (c) dislocation 5 has moved for 0.1 µm after release from the obstacle into a new position where it slows down again.
dislocations of 0.35 µm at this stress. This is compatible with the spacings seen in Fig. 13.12. Assuming an effective stress fraction σ*/σ = 0.2 gives σ* = 35 MPa. The curve for the tempered martensite steel in Fig. 13.9(b) shows that the average velocity vg of viscously gliding dislocations should be 2 × 10–4 µm s–1 at 873 K and, with the activation energy given in the Appendix, 0.6 µm s–1 at 823 K. The dislocation seen jumping between the positions in Fig. 13.12(a) and (b) moves particularly fast, at a velocity above 2 µm s–1. On average, the velocities seen in situ are lower by about a factor of 10. This is consistent with the average velocity estimated from Fig. 13.9. The radius r ≈ MbG/σloc of curvature of resting dislocations gives information about the locally acting stress σloc.81 In the subgrain interior typical values of r lie near
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0.7 µm corresponding to a local stress σloc = 33 MPa which is distinctly less than the estimated stress of 174 MPa acting in the investigated area. At the subgrain boundaries, r is found to be as small as 0.08 µm. This corresponds to a local stress of 290 µm in the immediate vicinity of the subgrain boundaries which significantly exceeds the acting stress. These observations are consistent with the concentration of stress at the subgrain boundaries postulated in the composite model. The strong curvature of the dislocation in Fig. 13.11(c) shortly before release from the obstacle (probably a precipitate) also indicates locally enhanced stress which corresponds to the relatively high velocity of the dislocation shortly after release. We conclude that the in situ observations support the validity of the microstructural model of creep.
13.9
Discussion and outlook
As shown in the preceding, rather simple expressions for structure evolution and dislocation velocity are available which explain a variety of creep features and even yield a quantitative formulation for creep after fitting to experimental results. Applying them to specific steels gives a feeling of the relative importance of hardening contributions. It is common to ask for the rate-determining mechanism of creep. Creep control by either climb or glide of dislocations are popular alternatives. However, the model described above teaches that creep is a complex plasticity phenomenon where a number of subprocesses interact as illustrated by Blum et al.43 Being a special case of crystal plasticity, creep must be described by plasticity models which include the description of dislocation motion as well as the evolution of the dislocation structure, which is connected to dislocation motion, and the evolution of the hardening phases.89 Glide and climb of dislocations enter the description of the subprocesses of creep, but do not control creep in an autonomous manner. As another example of coupling of subprocesses, we consider the motions of free dislocations and of low-angle boundaries. Subgrain boundary migration makes a certain constant relative contribution to strain lying in the order of 10% for pure metals.58,90 In this sense, subgrain boundary migration may control the rate of creep by glide of free dislocations. However, according to the composite model it is not an autonomous process, as the glide of free dislocations leads to stress concentration at the subgrain boundaries. One of the weakest points of microstructural creep models is the quantification of the particle influence on the motion of free dislocations and low-angle boundaries under stress. Particle hardening is essential in creep of steels. It enhances the creep resistance not only directly by adding a stress component (Equation [13.3]), but also indirectly by refining the subgrain structure caused by the enhanced stress level (Equation [13.1]) and stabilizing it against strain-controlled coarsening and recrystallization. It is notable that
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the apparent activation energies of creep are much higher than those of atomic diffusion in tempered martensite steels. This indicates a deficit in the present model which may have to do with the present description of particle strengthening. Approximation of the particle hardening in the subgrain interior by the Orowan stress may not be sufficient. Thermal activation by climb of dislocations over large particles and thermally activated detachment from small particles may need to be taken into account. This will lead to a temperature dependence of the particle hardening term in addition to that following from the temperature dependence of the solubility of alloying elements. There is also the possibility that the description of particle hardening of the subgrain boundaries by the Zener stress is insufficient. The Zener stress addresses pinning of the boundaries by particles. However, the motion of free dislocations through low-angle subgrain boundaries is the primary process to be considered in the composite model. This motion may be strongly impeded by dense arrays of large particles at the boundaries which probably have to be overcome by climb. In the model proposed in Section 13.6.1, particle hardening reduces the glide velocity simply by reducing the effective stress for glide between the particles. A better way to arrive at the average glide velocity of free dislocations may be explicitly to treat the stop-and-go mode of dislocation glide, with waiting at particles and relatively fast glide between the particles. This way was followed in deriving the creep rate of a mechanically alloyed Fe-base alloy.91 The formulation of the velocity of glide of free dislocations in the hard region (Equation [A.4] of the Appendix) poses a basic problem. A concentrated stress σh is needed by the free dislocations to cross the dense dislocations array of the low-angle subgrain boundaries. Once this is achieved, the dislocations bow out between the precipitates into the subgrain interior where the local stress is relatively small. This is seen from in situ pictures like those in Fig. 13.10 which show only limited bowing under relatively low stress although the dislocations are still linked to precipitates at the subgrain boundaries. This observation indicates that the model should be modified in the sense that the precipitates at the boundaries contribute to the particle hardening experienced by the free dislocations in the soft subgrain interior. The pinning effect exerted by the boundary particles on the subgrain boundaries will still be maintained. It leads to suppression of subgrain boundary migration which explains the relatively low value of the growth constant kw. Despite the existing deficiencies, modeling of creep on a microstructural basis is quite promising to be useful in guiding alloy development and giving qualitative or even semi-quantitative explanations for material properties. The formulation of the model in terms of differential equations allows complex loading paths similar to those occurring in practical applications to be simulated, provided that the relevant subprocesses of creep and the relevant microstructural
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elements enter the model in a reasonable manner. Coupling the deformation model with a thermodynamic data base has been attempted successfully80 and is a promising further step in model development. Industrial development of steels which resist creep will have to concentrate on restricting dislocation motion by enriching the steel matrix with strongly interacting, slowly diffusing solutes and a stable structure of fine, hard, slowly coarsening particles stabilizing the dislocation structure by opposing dislocation motion. As the creep resistance increases, the danger of creep fracture increases, too. Therefore the resistance of the material against nucleation of pores and microcracks requires additional attention.
13.10 Acknowledgments Thanks are due to Dr.-Ing. S. Straub and Dr.-Ing. P. Polcik, who explored the microstructure of the steels and applied the composite model to them, to Dr.Ing. H. Chilukuru and R. Agamennone for further scientific progress reached in their works, to H. Chilukuru and W. Wranik for support in preparing the figures, to Prof. U. Messerschmidt, Dr. M. Bartsch and their co-workers from MPI für Mikrostrukturphysik Halle for performing in situ TEM observations, to Dr. J. Granacher, Dr. A. Scholz and Prof. C. Berger from TU Darstadt for providing long-term crept specimens and to the Deutsche Forschungsgemeinschaft, the Bundesministerium für Wirtschaft and the associated group of industrial companies for continuous financial support.
13.11 References 1 B. Wilshire, ‘On the evidence for diffusional creep processes’, In B. Wilshire and R. W. Evans (eds), Creep and Fracture of Engineering Materials and Structures, The Institute of Metals, London, 1990, 1–9. 2 F. R. Beckitt and B. R. Clark, ‘The shape and mechanism of formation of M23C6 carbide in austenite’, Acta Metall., 1967, 15, 113–129. 3 B. Sasmal. ‘Mechanism of the formation of M23C6 plates around undissolved NbC particles in a stabilised austenitic stainless steel’, J. Mater. Sci. 1997, 32, 5439– 5444. 4 C. M. Wayman, Introduction to the Crystallography of Martensitic Transformations, Macmillan Series in Material Science, The Macmillan Company, New York, 1964. 5 P. Haasen, Physical Metallurgy, Cambridge University Press, Cambridge, 2nd edition, 1986. 6 G. Eggeler, N. Nilsvang and B. Ilschner, ‘Microstructural changes in a 12% chromium steel during creep’, Steel Res. 1987, 58, 97–103. 7 F. Abe, H. Araki and T. Noda, ‘Microstructural evolution in bainite martensite, and δ ferrite of low activation Cr-2W ferritic steels’, Material Science and Technology, 1990, 6, 714–723. 8 P. J. Ennis, A. Zielińska-Lipiec and A. Czyrska-Filemonowicz, ‘Influence of heat treatments on microstructural parameters and mechanical properties of P92 steel’, Materials Science and Technology, 2000, 16(10), 1226–1232.
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9 F. Yoshida, D. Terada, H. Nakashima, H. Abe, H. Hayakawa and S. Zaefferer, ‘Microstructure change during creep deformation of modified 9Cr-1Mo steel’, In Advances in Materials Technology for Fossil Power Plants, Proceedings of the 3rd Conference, R. Viswanathan, W.T. Bakker and J. D. Parker (eds), held at University of Wales Swansea, 5–6 April The Institute of Materials, London, 2001, 143–151. ∨ 10 A. Dronhofer, J. Peši c ka, A. Dlouhý and G. Eggeler, ‘On the nature of internal interfaces in tempered martensite ferritic steels’, Z. Metallkd., 2003, 94(5), 511–520. 11 A. H. Cottrell, An Introduction to Metallurgy, Edward Arnold London, 1968. 12 M. Winning, G. Gottstein and L. S. Shvindlerman, ‘Migration of grain boundaries under the influence of an external shear stress. Mater. Sci. Eng., 2001, A317(1–2), 17–20. 13 E. E. Underwood, Quantitative Stereology, Addison-Wesley, Reading, MA, 1970. 14 H. Chilukuru, W. Blum, M. Schwienheer and A. Scholz, ‘Creep-fatigue interaction in martensitic tempered steels by dynamic subgrain growth’, In Langzeitverhalten warmfester Stähle und Hochtemperaturwerkstoffe, Beiträge zur 26. Vortragsveranstaltung der Arbeitsgemeinschaft für warmfeste Stähle und für Hochtemperaturwerkstoffe am 28. November in Düsseldorf, Stahlinstitut VDEh, 2003, 65–74. 15 W. Blum and G. Götz. Evolution of dislocation structure in martensitic CrMoVsteels: the subgrain size as a sensor for creep strain and residual creep life. Steel Res, 1999, 70, 274–278. 16 S. Straub, M. Meier, J. Ostermann and W. Blum, ‘Development of microstructure and strengthening in the ferritic steel X20 CrMoV12 1 at 823 K during long-term creep tests and during annealing’, VGB Kraftwerkstechnik, 1993, 73, 646–653. 17 S. Straub, Verformungsverhalten und Mikrostruktur warmfester martensitischer 12%Chromstähle, Fortschr.-Ber. VDI Reihe 5 Nr. 405. VDI Verlag, Düsseldorf, 1995. 18 S. Straub and W. Blum, ‘Microstructural development and deformation kinetics in iron-and nickel-base alloys’, In Proceedings of the International Symposium on Hot Workability of Steels and Light Alloys-Composites, H. J. McQueen, E. V. Konopleva and N. D. Ryan, editors, Montréal, 1996, 189–203. 19 P. Polcik, Mikrostruktur und Verformungsverhalten des Stahles X22 CrMoV12 1 nach Zeitstandbeanspruchung im 105 h Bereich, PhD Thesis, Universität ErlangenNürnberg, 1993. 20 F. Abe, ‘Effect of quenching, tempering, and cold rolling on creep deformation behavior of a tempered martensitic 9Cr-1W steel’, Metall. Mater. Trans. A, 2003, 34A, 913–925. 21 B. Reppich, ‘Particle strengthening’, In Plastic Deformation and Fracture of Materials, H. Mughrabi (ed), Volume 6 of Materials Science and Technology, (ed. by Cahn, R. W. and Haasen, P. and Kramer, E. J.), VCH Verlagsge-sellschaft, Weinheim, 1993, 311–357. ∨ 22 K. Mili cka, ‘Internal stress and structure in creep of cold prestrained Fe-21Cr-32Ni alloy at 975 K. Metal Sci., 1982, 16, 419–424. 23 B. Wilshire and R. M. Willis, ‘Creep and creep fracture of prestrained type 316H stain-less steel’, In Proceedings of the 10th Joint International Conference on Creep & Fracture of Engineering Materials and Structures, Creep Resistant Metallic Materials, M. Filip, V. Foldyna, R. Gladiš, A. Jakobová, Z. Kuboň, J. Purmenský and J. Sobotka, editors, 8–11 April 2001, Vitkovice-Research and Development and TERIS 2002, Prague, Czech Republic, 2001, 6–15. 24 P. D. Portella, Monotones und zyklisches Kriechverhalten der Legierung 800H bei 800°C, PhD Thesis, University of Erlangen-Nürnberg, 1984.
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25 H. Mecking and U. F. Kocks, ‘Kinetics of flow and strain hardening,’ Acta Metall., 1986, 29, 1865–1875. 26 U.F. Kocks and H. Mecking, ‘Physics and phenomenology of strain hardening: the FCC case. Progr. Mater. Sci., 2003 48(3), 171–273. 27 W. Blum, ‘High-temperature deformation and creep of crystalline solids’, In H. Mughrabi, editor, Plastic Deformation and Fracture of Materials, Volume 6 of Materials Science and Technology, ed. by Cahn, R. W. and Haasen, P. and Kramer, E. J., VCH Verlagsgesellschaft, Weinheim, 1993, 359–405. 28 W. Blum, S. Straub and S. Vogler, ‘Creep of pure materials and alloys. In D. G. Brandon, R. Chaim and A. Rosen, editors, Proceedings 9th International Conference on the Strength of Metals and Alloys (ICSMA 9), Vol. I, Freund Publishing House, 12 London, 1991, 111–126. 29 S. U. An, H. Wolf, S. Vogler and W. Blum, ‘Verification of the effective stress model for creep of Inconel 617 at 800°C. In B. Wilshire and R.W. Evans, editors, Creep and Fracture of Engineering Materials and Structures, The Institute of Metals, London, 1990, 81–95. 30 W. Blum and F. Breutinger, ‘New method of determining stress relaxation behavior in creep machines by controlled unloading’, Z. Metallkd., 2003, 93(7), 649–653. 31 B. Backes, Ermittlung des Verformungswiderstandes eines neuen, hochwarmfesten 12% Cr-Stahls durch Spannungsrelaxation, Master’s Thesis, University of ErlangenNürnberg, 2003. 32 G. B. Gibbs, ‘Creep and stress relaxation studies with polycrystalline magnesium’, Philos. Mag. A, 1966, 317–329. 33 C. N. Ahlquist and W. D. Nix, ‘A technique for measuring mean internal stress during high temperature creep’, Scripta Metall., 1969, 679–682. 34 W. Blum and A. Finkel, ‘New technique for evaluating long range internal back stresses’, Acta Metall., 1982, 30, 1705–1715. 35 W. Blum, P. D. Portella and R. Feilhauer, ‘Zyklisches Kriechverhalte’, In B. Ilschner, editor, Festigkeit und Verformung bei hoher Temperatur, Deutsche Gesellschaft für Metallkunde, Oberursel, 1983, 41–59. 36 H. Wolf, M. Schießl and W. Blum, ‘Acceleration of creep due to cyclic loading of a CrMo-steel’, In H. J. McQueen, J. P. Bailon, J. I. Dickson, J. J. Jonas and M. G. Akben, editors, Proceedings 7th International Conference on the Strength of Metals and Alloys (ICSMA 7), Pergamon Press, Oxford, Montreal, Canada, 1985, 607–612. 37 W. Blum and J. Granacher, ‘Cyclic creep of heat resistant steels’, In D. G. Brandon, R. Chaim and A. Rosen, editors, Proceedings 9th International Conference on the Strength of Metals and Alloys (ICSMA 9), Vol. I, Freund Publishing House, London, 1991, 429–436. 38 S. Straub, P. Polcik, D. Henes and W. Blum, ‘Simulation of the long-term cyclic creep behaviour of a low alloyed ferritic chromium steel’, Mater. Sci. Eng., 1997, A234–236, 1037–1040. 39 A. Seeger, Dislocations and Mechanical Properties of Crystals, Wiley, New York, 1957. 40 E. Nes, ‘Modelling work hardening and stress saturation in FCC metals,’ Progr. Mater. Sci., 1998, 41(3), 129–193. 41 H. Oikawa and T. G. Langdon, ‘The creep characteristics of pure metals and metallic solid solution alloys’, In B. Wilshire and R. W. Evans, editors, Creep Behaviour of Crystalline Solids, Swansea, Pineridge Press, 1985, 33–82. 42 A. Cottrell and M. Jaswon, ‘Distribution of solute atoms round a slow dislocation’,
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45 46 47 48
49
50 51 52 53 54 55
56 57 58 59 60 61
62
Creep-resistant steels Proc. Roy. Soc. London, Series A, Mathematical and Physical Sci. (1934–1990), 1949, 199(1056), 104–114. W. Blum, P. Eisenlohr and F. Breutinger, ‘Understanding creep – a review’, Metall. Trans., 2002, 33A, 291–303. F. Breutinger and W. Blum, ‘Effect of dynamic strain ageing on creep of commercially pure titanium. In J.D. Parker, editor, Proceedings 9th International Conference on Creep and Fracture of Engineering Materials and Structures, The Institute of Metals, London, 2001, 39–48. W. Blum, Y. J. Li and F. Breutinger, ‘Deformation kinetics of coarse-grained and ultrafine-grained commercially pure Ti’. Mater. Sci. Eng. A, 2007, 462, 275–278. B. Schönfeld, ‘Local atomic arrangements in binary alloys’, Prog. Mater. Sci., 1999, 44, 435–543. H. Stöcker, Taschenbuch der Physik, Verlag Harri Deutsch, Thun und Frankfurt am Main, 1994. M. Meier and W. Blum, ‘Accounting for subgrain hardening in NiCr22Co12Mo9 with the composite model. In B. Wilshire and R. W. Evans, editors, Creep and Fracture of Engineering Materials and Structures, The Institute of Materials, London, 1993, 167–178. P. Polcik, Modellierung des Verformungsverhaltens der warmfesten 9–12%Chromstähle im Temperaturbereich von 550–650°C. D29, Dissertation, Universität ErlangenNürnberg, Shaker Verlag, Aachen, 1999. B. Garofalo, Fundamentals of Creep and Creep Rupture, Macmillan, New York, 1965. E. N. da C. Andrade, ‘On the viscous flow in metals and allied phenomena’, Proc. Roy. Soc. Lond., A, 1910, 84, 1–12. B. Ilschner, Hochtemperaturplastizität, Springer, Berlin, 1973. H. Wolf, Kriechen der Legierungen NiCr22Co12Mo und 10CrMo9 10 bei konstanter und zyklischer Beanspruchung, PhD Thesis, University of Erlangen-Nürnberg, 1990. F. Abe, ‘Creep rates and strengthening mechanisms in tungsten-strengthened 9cr steels’, Mater. Sci. Eng., 2001, A319–321, 770–773. F. Abe, ‘Strengthening mechanisms in steel for creep and creep rupture’, In Creep Resistant Steels, F. Abe, U. Kern and R. Viswanathan (eds), Chapter 9. Woodhead Publishing, Cambridge, 2007 (this book). W. Blum, J. Hausselt and G. König, ‘Transient creep and recovery after stress reduction during steady state creep of AlZn. Acta Metall., 1976, 24(4), 293–297. W. Blum, ‘On the evolution of the dislocation structure during work hardening and Creep’, Scripta Metall., 1984, 18(12), 1383–1388. W. Müller, M. Biberger and W. Blum, ‘Subgrain boundary migration during creep of LiF, III. Stress reduction experiments. Phil. Mag. A, 1992, 66, 717–728. M. Hättestrand, M. Schwind and H. O. Andrén, ‘Mycroanalysis of two creep resistant 9–12% chromium steels’, Mater. Sci. Eng. A, 1998, 250, 27–36. J. Hald. Long-term stability of 9–12% Cr Steels – current understanding and future perspectives. In Werkstoffe und Qualitätssicherung 2004, VGB, Dortmund, March 2004. W. Blum, ‘Creep Simulation,’ In L.-Q. Chen, F. Barlat, F. Roters and D. Raabe, editors, Continuum Scale Simulation of Engineering Materials Fundamentals, Microstructures, Process Applications. Wiley-VCH, Weinheim, 2004. J. S. Dubey, H. Chilukuru, J. K. Chakravartty and W. Blum. On effects of cyclic deformation and hold periods on subgrain structure and creep behaviour in tempered martensitic 9–12%CrMoV-steels. Mater. Sci. Eng. A, 2005, 406, 152–159.
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63 W. Blum and H. Chilukuru, ‘Tertiary creep of tempered martensite 9–12 microstructural origin of creep life limitation’, In Pressure Vessels and Piping: Materials and Properties, Proceedings International Conference on Pressure Vessels and Piping (OPE 2006– Chennai, Feb 7–9, 2006), B. Raj, B. K. Choudhary and A. Kumar, (eds), Narosa Publishing House, Chennai, ASM International, Alpha Science, 2007. 64 A. Kostka, K.-G. Tak, R. J. Hellmig, Y. Estrin and G. Eggeler, ‘On the contribution of carbides and micrograin boundaries to the creep strength of tempered martensite ferritic steels, Acta Mater., 2007, 55, 539–550. 65 A. Umantsev and G. B. Olson, ‘Ostwald ripening in multicomponent alloys, Scripta Metall. Mater., 1993, 29(8), 1135–1140. 66 J. Ågren, M. T. Clavagura-Mora, J. Golcheski, G. Inden, H. Kumar, and C. Sigli. Application of computational thermodynamic to phase transformation nucleation and coarsening. Calphad, 2000, 24(1), 41–54. 67 D. Goblirsch, Verformung von X20 CrMoV 12 1 bei konstanter Druckspannung unter besonderer Berücksichtigung der Übergangsverformung nach spannungswechsel, Master’s Thesis, University of Erlangen-Nürnberg, Erlangen, Germany, 1990. 68 R. Agamennone, W. Blum, C. Gupta and J. K. Chakravartty, ‘Evolution of microstructure and deformation resistance in creep of tempered martensitic 9–12%Cr– 2%W–5%Co steels. Acta Mater., 2006, 54, 3003–3014. 69 H. Chilukuru. On the Microstructural Basis of Creep Strength and Creep–fatigue Interaction in 9–12% Cr Steels for Application in Power Plants, PhD Thesis, University of Erlangen-Nürnberg, Erlangen, Germany, 2007. 70 P. D. Portella and W. Blum, ‘Cyclic creep of Incoloy 800 H at 1073 K’, In J. B. Bilde-Sørensen, N. Hansen, A. Horsewell, T. Leffers and H. Lilholt, editors, Deformation of Multi-phase and Particle Containing Materials, Risø National Laboratory, Roskilde, Denmark, 1983, 493–498. 71 H. Wolf and W. Blum, ‘Acceleration and deceleration of creep of a 21/4 Cr–1Mo steel by cyclic stressing at 550°C’, In B. Wilshire and R. W. Evans, editors, Proceedings 3rd International Conference on Creep and Fracture of Engineering Materials and Structures, The Institute of Metals, London, 1987, 649–662. 72 P. Eisenlohr, On the Role of Dislocation Dipoles in Unidirectional Deformation of Crystals. Dr.-Ing. Thesis, Universität Erlangen-Nürnberg, 2004. 73 P. Eisenlohr and W. Blum, ‘Bridging steady-state deformation behavior at low and high temperature by considering dislocation dipole annihilation’, Mater. Sci. Eng. A, 2005, 400–401, 175–181. 74 P. Weinert, Modellierung des Kriechens von Ferritisch/Martensitischen 9–12%Cr– Stählen auf Mikrostruktureller Basis. PhD Thesis, Technische Universität Graz, Graz, Austria, 2001. 75 D. A. Hughes, N. Hansen and D. J. Bammann. Geometrically necessary boundaries, incidental dislocation boundaries and geometrically necessary dislocations. Scripta Mater., 2003, 48, 147–153. 76 J. Rösler and E. Arzt, ‘A new model-based creep equation for dispersion strengthened materials’, Acta Metall. Mater., 1990, 38, 671–683. 77 G. Eggeler, ‘The effect of long-term creep on particle coarsening in tempered martensite ferritic steels’, Acta Metall., 1989, 37, 3225–3234. 78 E. Cerri, E. Evangelista, S. Spigarelli and P. Bianchi, ‘Evolution of the microstructure in a modified 9Cr–1Mo steel during short term creep’, Mater. Sci. Eng. A, 1998, 245(2), 285–292. 79 H. Chilukuru, K. Durst, M. Göken and W. Blum, ‘On the roles of M2X and Z-phase
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Creep-resistant steels in tempered martensitic 9–12% Cr steels’, In J. Lecomte Beckers, M. Carton, F. Schubert and P. J. Ennis, editors, Materials for Advanced Power Engineering, Proceedings of the 8th Liege Conference, volume Part III, 2006, 1181–1190. T. Barkar and J. Agren, ‘Creep simulation of 9–12% cr steels using the composite model with thermodynamically calculated input’, Mater. Sci. Eng. A, 2005, 395, 110–115. H. Mughrabi, ‘Dislocation walls and cell structures and long-range internal stresses in deformed metal crystals’, Acta Metall., 1983, 31(9), 1367–1379. W. Blum, A. Rosen, A. Cegielska and J. L. Martin, ‘Two mechanisms of dislocation motion during creep’, Acta Metall., 1989, 37, 2439–2453. S. Vogler and W. Blum, ‘Micromechanical modelling of creep in terms of the composite model’, In B. Wilshire and R. W. Evans, editors, Creep and Fracture of Engineering Materials and Structures, The Institute of Metals, London, 1990, 65–79. ∨ R. Sedlá c ek and W. Blum, ‘Microstructure-based constitutive law of plastic deformation’, Comp. Mater. Sci., 2002, 25(1–2), 200–206. D. Henes, H. Möhlig, S. Straub, J. Granacher, W. Blum and C. Berger, ‘Microstructurebased modelling of the long-term monotonic and cyclic creep of the martensitic steel X20(22) CrMoV12 1’, In Microstructure and Mechanical Properties of Metallic High-Temperature Materials, H. Mughrabi, G. Gottstein, H. Mecking, H. Riedel and J. Tobolski, (eds) Wiley-VCH, Weinheim, 1999, 179–191. P. Polcik, S. Straub and W. Blum, ‘Simulation of the deformation behaviour of martensitic 9–12% chromium steels on a microstructural basis’, In The 4th European Conference on Advanced Materials and Processes, Associatione Italiana di Metallurgia, Mailand, 1995, 313–318. U. Messerschmidt, M. Bartsch, C. Dietzsch, R. Agamennone, C. Gupta and W. Blum, unpublished results. Y. Qin, G. Götz and W. Blum, ‘Subgrain structure during annealing and creep of the cast martensitic Cr-steel G-X12CrMoWVNbN10-1-1, Mater. Sci. Eng. A, 2003, 341, 211–215. B. Holmedal, K. Marthinsen and E. Nes, ‘A unified microstructural metal plasticity model applied in testing, processing, and forming of aluminium alloys. Z. Metallkd., 2005, 96(6), 532–545. S. F. Exell and D. H. Warrington, ‘Sub-grain boundary migration in aluminum’, Phil. Mag. A, 1972, 26, 1121–1136. R. Herzog, H. Schuster, M. Weiße and W, ‘Blum. Mikrostruktur und mechanische Eigen-schaften der ODS-Eisenbasislegierung PM 2000 bei quasistationärer und instationärer Verformung’, In Werkstoffwoche ’96, Symposium 7, Materialwissenschaftliche Grundlagen, F. Aldinger and H. Mughrabi (eds), DGMInformationsgesellschaft, Frankfurt, 1997, 243–248. U. Hofmann, Kriechverhalten des austenitischen stahles X3 CrNi 18 9 bei 923 K, Master’s Thesis, Universität Erlangen-Nürnberg, Erlangen, Germany, 1991.
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13.12 Appendix: Microstructural model Mikora Soft (s) subgrains surrounded by hard (h) boundaries (thickness a) deform inelastically (inel) at equal mechanical (mech) strain so that the local stresses differ.84 ε mech = ε inel,s +
σs σ = ε inel,h + h E E
[A.1]
[A.2] fh = 2 a w where f are the volume fractions. The local rates of deformation are described by Equation [13.2]. For steels from the group of 9–12%Cr steels the local dislocation velocities were formulated as: σ = fsσs + fhσh,
fs = 1 – fh ,
Q v s = As exp – s σ *s ⋅ exp (( β σ *s ) ξ ) RT
[A.3]
b 2 s ( σ h – σ p,h ) Q v h = Ah exp – h sinh RT Mk B T
[A.4]
σ *s = σ s – α MbG ρf0.5 – C p,s σ p,s
[A.5]
Particle hardening was expressed by Polcik49 by adding the Orowan stresses in the soft region and the Zener stresses in the hard region:
σ p,s = 3.32 Gb Σ k
σ p,h =
0.5 f p,s,k dk
f p,h,k k1 Σ b k dk
[A.6] k1 = 7.7 × 10–9 N.
[A.7]
where fp,s,k and fp,h,k are the local volume fractions of the precipitate phases k existing in the soft and hard regions, respectively. The dislocation interaction parameter α evolves with strain from α = 0 at εinel = 0 as:
dα = 0.25 – α d ε inel,s ( b / M ) ρf0.5
[A.8]
The empirical factor Cp,s ≤ 1 was introduced by Polcik49 to take into account that the particle strengthening term falls below the Orowan stress σs,Or owing to climb over particles and to ensure that σ *s > 0 even for small σ. The characteristic dislocation spacings w, ρf–0.5 and a evolve with strain –0.5 = 100 according to Equation [13.10] from their starting values w 0 , ρf,0 17 nm and a0 = 0.025 w0 towards their steady state values given by Equation [13.1], δ∞ = 0.39 w∞ and a∞ = 0.025 w∞ with kw = 0.12, kδ = 0.0005 and ka
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= 0.003. s was set equal to 7.6 nm. The precipitate sizes evolve with time t according to Equation [13.9]; see Agamennone et al.68 and Chilukuru69 for values of growth constants. The volume fractions fp are assumed to remain constant except for the Laves phase which precipitates during creep. The starting values w0, fp,0 and dp,0 are the microstructural input parameters. Material parameters: b = 0.248 nm, M = 3, E = 165 GPa, G = 62 GPa. Constants: Qs = Qh = 562 kJ mol–1, As = 3.6 × 1015 m s–1 Pa–1, Ah = 4.4 × 1020 m s–1, β (MPa) and ξ are (i) 0.043 and 1.55 and (ii) 0.48 and 0.45 for σ *s below and above 62 MPa, respectively.
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14 Constitutive equations for creep curves and predicting service life S . R . H O L D S W O R T H , EMPA – Materials Science & Technology, Switzerland
14.1
Introduction
Creep strain curves are determined from the results of continuous-measurement or interrupted tests involving the application of a constant load (or stress) to a uniaxial testpiece held at constant temperature (Fig. 14.1). In continuousmeasurement tests, the creep strain, εf, is monitored without interruption by means of an extensometer attached to the gauge length of the testpiece (EN 10291, 2000). In interrupted tests, the total plastic strain, εp, is determined from optical measurements of εper at room temperature during planned interruptions (where εper = εp – εk i.e. εper = εi + εf – εk, Fig. 14.2, and εp = εper when εk ≈ 0). A list of the symbols and terms used in this chapter is given in the Nomenclature Section 14.7. The creep curve data collected in this way may then be modelled by a constitutive equation. Depending on the nature of the creep model application, the analysis will be of several εf(t) or εp(t) curves determined for a single cast or several casts of the specified material. The creep strain curves may have been determined from a matrix of t(T, σ) tests for which T and σ are: (1) relatively homogeneously distributed or (2) inhomogeneously distributed in terms of T, σ and cast of material. Case (1) is the ideal situation and generally arises within R&D projects or well co-ordinated data generation activities. Case (2) is more typical of large multi-national datasets, comprising information from many casts, gathered to produce alloy representative creep strength values for Standards (e.g. Holdsworth et al., 2005). A constitutive equation is a relation between two or more physical quantities which may be simply phenomenological or be directly derived from first principles with a physical basis. The requirement for a representative description of a material’s εf(t, T, σ) creep strain behaviour is no longer just for scientific interest and metallurgical understanding. The creep deformation behaviour of engineering components is now routinely evaluated using PC-based finiteelement analysis tools. Design engineers require the parameters for model equations to describe the long-time creep behaviour of a specified alloy type 403 WPNL2204
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Au
Strain
T Creep regimes S P tu
Time
14.1 Schematic representation of creep curve showing primary (P), secondary (S) and tertiary (T) deformation regimes.
εt εe
εp
εi
εf
σ0 t = t1
Strain
t=0
ET
T = constant
εper
εk
εe
Strain
14.2 Schematic representation of strains generated during a creep test.
(not simply the characteristics of a single cast), typically in the primary and secondary deformation regimes (i.e. the P and S regimes in Fig. 14.1). In contrast, remaining life assessment engineers are more likely to require the best model description for a single cast of material, in the secondary and tertiary creep regimes. The following sections contain a review of creep constitutive equations (Section 14.2) and an approach for assessing model-fitting effectiveness (Section 14.3). In Section 14.4, the application of constitutive equations for
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service life prediction is examined. Finally, there is a consideration of future trends (Section 14.5) and some concluding remarks (Section 14.6).
14.2
Constitutive equations
A wide range of creep model equations are in use today to represent the high temperature-time dependent deformation behaviour of engineering materials (e.g. Table 14.1). Many of these comprise components originating from a small number of classical representations of primary, secondary and/or tertiary creep deformation (e.g. Table 14.2, with reference to Fig. 14.1). Typically, the effect of temperature is acknowledged by incorporating an Arhenius, A exp (–Q/RT), function into the equation (examples are given in Table 14.1). As a generality, logarithmic creep only occurs at lower temperatures (i.e. below 0.3Tm). At higher temperatures, primary creep is more typically represented by power, exponential or sinh functions of time (Table 14.2). Similarly, secondary creep rate can be represented by power, exponential and sinh functions of stress. For secondary creep rate, a sinh formulation reduces to a power law at low stresses and an exponential law at higher stresses. Tertiary creep or creep rate is typically modelled by power or exponential representations, with or without a damage accumulation function, see for example Kachanov (1958). More recently there has been a tendency, in particular for precipitation strengthened alloys, to replace σ in certain of the models listed in Table 14.1 by (σ – σi) to acknowledge the existence of a friction stress, see for example McLean (1980). The list of constitutive equations contained in Table 14.1 is not exhaustive, but is representative of the range of creep deformation models currently employed within the power generation sector in Europe, as reviewed recently by the European Creep Collaborative Committee (ECCC, 2005b).
14.3
Constitutive equation selection
No single constitutive equation effectively represents the creep deformation characteristics of all materials over their entire temperature application range. The effectiveness of a constitutive equation to model primary, secondary and/or tertiary creep deformation for specific applications can vary with material characteristics and source data distribution. In particular, not all model equations and fitting procedures are suitable for the prediction of alloy-mean long-time creep strength behaviour (Holdsworth et al., 2005).
14.3.1 Model fitting effectiveness The ability of a constitutive equation effectively to characterise the creep deformation behaviour of a material depends not only on model characteristics
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Table 14.1 Review of creep equations used in ECCC assessment intercomparisons Model reference
Creep equation
Norton, (Norton, 1929)
ε˙ f,min = d 1 exp (–Q /RT ) σ n
Modified Norton Norton–Bailey Bartsch
ε˙ f,min = b1 exp(–QB/RT )σn + c1 exp(QC/RT )σn εf = d1σntp εf = e1 exp (–Q1/RT )σ exp(–b1σ)t p
(Bartsch, 1995) Garofalo, (Garofalo, 1965) Modified Garofalo (Granacher, et al., 2001) BJF (Jones and Bagley, 1996) Li–Akulov model (Li, 1963; Akulov, 1964) Theta (Evans and Wilshire, 1985) Modified Theta Graham–Walles (Graham and Walles 1955) Modified Graham–Walles
+ e2 exp(–Q2/RT )σ exp(b2σ)t εf = εt[1 – exp(–b1t)] + ε˙ f,min t εf = εf1[1 – exp(–g1(t/t12)u) + ε˙ f,min t + c23(t/t23)f ]
εf = n1[1 – exp(–t)]β + n2t where t = (σ/A1)n exp (–Q/RT) ε˙ i – ε˙ f,min ln 1 + (1 – exp (– kt )) + ε˙ st k ε˙ f,min + εT(exp (t/tt) – 1) εf =
ε˙ f,min
εf = θ1[1 – exp(–θ2t)] + θ3[exp(θ4t) – 1] where log(θi) = ai + biT + ci σ + di σT εf = θ1[1 – exp(–θ2t)] + θmt + θ3[exp(θ4t) – 1] where θm = Aσnexp(–Q/RT) εf = at1/3 + dt + ft 3 σ (1 + ε ) ε˙ f = e ( Q 1 /T ) 10 A 1 1+ω
σ (1 + ε ) + e ( Q 2 /T ) 10 A 2 1+ω
n1
ε m1
n2
where ω˙ = e(–QD/T)10AD(σ(1 + ε))nDεmD
h1σ n k 1σ ν ω˙ = (1 – ω ) (1 – ω ) ζ
Rabotnov–Kachanov
ε˙ =
(Kachanov, 1986) Dyson and McClean, (Dyson and McClean, 1998)
σ (1 – H ) ε˙ f = ε 0′ (1 + D d )exp(–Q /RT )sinh σ 0 (1 – D p )(1 – ω )
Baker–Cane model
εf = Atm + εp + φεs + εs(λ – φ)
(Baker and O’Donnell, 2003) Mech. E (CSWP, 1983) Characteristic strain model (Bolton, 2005a) MHG model, (Grounes, 1969) (Holmström and Auerkari, 2004) Omega, (Prager, 1995)
where l = εu/εs, εs = ε˙ mt u Ru/t/T = (a1 + b1/ε – c1ε2)Rε/t/T + d1 + e1/ε + f1/ε2 – g1ε2 εf(σ) = ε(Ru/t/T/Rε/t/T – 1)/(Ru/t/T/σ – 1)
Modified Omega (Merckling, 2002)
1– φ
t /t u – φ λ – φ l – 1–φ and φ = tp/tu
tε = exp(TF(ε, σ) + C) where the F(ε,σ) function is freely selected from multilinear combinations of σ and ε with an optimised value of C ε˙ f = ε˙ f,min /(1 – ε˙ f,min Ωt ) ε f = 1 – 1 (– ln(tu – t) + ln(tu)) Ω 2C tr + Ctr(1 – exp(mtrt))
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Table 14.2 Classical representations of primary, secondary and tertiary creep Model equation
Primary creep Logarithmic: εf = a log(1 + bt) Power: εf = at b Exponential: εf = a(1 – exp(–bt)) Sinh: εf = a sinh (bt c ) Secondary creep Power: ε˙ f,min = d σ n Exponential: ε˙ f,min = d exp (e σ ) Sinh: ε˙ f,min = d sin h (e σ ) Tertiary creep Power: εf = ft g Exponential: εf = f (exp(–gt)–1) c σk aσn where ω˙ = ε˙ f = (1 – ω )r (1 – ω )q Omega: ε˙ f = ε˙ 0 exp (Ωε )
Source reference
Phillips (1905) Graham and Walles (1955) McVetty (1933) Conway and Mullikin (1962) Norton (1929) Nadai (1938) Graham and Walles (1955) McHenry (1943) Kachanov (1958) Rabotnov (1969) Prager (1995)
but also on model fitting approach. In an ECCC creep strain assessment intercomparison activity, it became clear that model fitting effectiveness can be strongly influenced by the rigour of the analyst and the model fitting procedure applied. Three generic approaches were adopted in the ECCC εf(T, σ, t) model-fitting activity: Approach-1 • model-fitting individual experimental creep curves with the selected constitutive equation (simultaneously or consecutively for individual deformation regimes) to establish the model parameters for specific conditions of T and σ, • determination of the temperature and stress dependence of the selected model parameters to define the material mean master equation for all εf(T, σ, t). Approach-2 • determination of specific εf(T, σ, t) coordinates from individual experimental creep curves either directly (unconstrained by a formal model description) or as a result of model fitting (with a model different to that used for final fitting), • parametric model-fitting of the specific coordinates to establish parameters to define the mean master equation for the material in the form of either εf(T, σ, t) or εf(T, σ).
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Approach-3 • derivation of mean σε(T, t) relationships from specific observed tε(T, σ) coordinates from individual experimental creep curves, • model-fitting with derived σε(T, t) relationships to establish parameters for the selected model to define the mean master equation for the material. Using these approaches, ECCC investigated ways of quantifying the effectiveness of a number of constitutive equations to represent the creep deformation characteristics of large datasets of three alloys, namely 21/4CrMo low alloy ferritic steel (P22), 9CrMoVNb martensitic stainless steel (P91) and 18Cr8Ni austenitic stainless steel (TP316) (Holdsworth et al., 2005). The initial focus was on a large single-laboratory/single-cast multi-temperature t(T, σ) creep curve dataset for 21/4CrMo steel (Holdsworth and Merckling, 2003). The T, σ test conditions responsible for the creep curves in this dataset were relatively homogeneously distributed to give six rupture times in the range ~100 to ~3000 h at five temperatures between 510 and 600°C. Moreover, each creep curve comprised approximately the same number of ε(t) observations (~400). The study involved the application of 16 constitutive equations by nine analysts and resulted in the development of the Z-parameter to provide a measure of model-fitting effectiveness in specific creep strain regimes. It was concluded that as a generality, specific model equations are better suited to representing creep strain accumulation characteristics for a given material in either the primary/secondary regimes or the secondary/ tertiary regimes (e.g. Fig. 14.3), although some models can be suitable for both (e.g. Fig. 14.4). The Z-parameter provides a means of quantifying modelfitting effectiveness (Equation [14.1], Figs 14.3 and 14.4). log ( t p*ε / σ / T ) = log ( t pε / σ / T ) ± 2.5s A–RLT = log ( t pε / σ / T ) ± log ( Z ) [14.1] For a normal distribution, almost 99% of the observed times to specific strain values would be expected to lie within the boundary lines defined by Equation 14.1. A perfect prediction of tpε/σ/T by the master equation is represented by Z equal to zero. Ideally Z should be ≤2 (Holdsworth and Merckling, 2003). The multi-source/multi-cast multi-temperature datasets for the P91 and TP316 steels were much larger, each containing almost 100 creep curves with rupture durations ranging from ~100 to >50 000 h (Fig. 14.5). The temperatures and stresses for which the creep curves were generated were not uniformly distributed and the number of ε(t) observations per curve ranged between 2 and 250, the creep curves having been generated by both interrupted and continuous measurement testing. These datasets were more typical of those used to form the basis of mean alloy (rather than single
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3
Log predicted time (h)
Theta model
z=2
2
Eqn. 14.1 1
510C 540C 565C 580C 600C
0
–1 –1
0
1 Log observed time (h) (a)
2
3
3
Log predicted time (h)
Theta model
2
1
510C 540C 565C 580C 600C
0
z=2 Eqn. 14.1 –1 –1
0
1 Log observed time (h) (b)
2
3
14.3 Example comparisons of predicted and observed times to (a) 0.2% and (b) 1.0% plastic strain for a constitutive equation providing a poor fit to the experimental data in the low primary/secondary strain range regime (Z >> 2), but a good fit to the experimental data at high strains in the secondary/tertiary deformation regime (Z = 2).
cast) long time strength values for European Product Standards (Holdsworth et al., 2005). In the case of the P91 and TP316 datasets, 11 constitutive equations were applied by 11 analysts. For such large inhomogeneously distributed datasets, Approaches 2 or 3 appeared to be the most appropriate
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Log predicted time (h)
Garofalo modified model
z=2 Eqn. 14.1
2
1
510C 540C 565C 580C 600C
0
–1 –1
0
1 Log observed time (h) (a)
2
3
3
Log predicted time (h)
Garofalo modified model 2
1
510C 540C 565C 580C 600C
0 Eqn. 14.1
z=2 –1 –1
0
1 Log observed time (h) (b)
2
3
14.4 Example comparisons of predicted and observed times to (a) 0.2% and (b) 1.0% plastic strain for a constitutive equation providing good fits to the experimental data in the primary, secondary and tertiary creep regimes (Z~2).
for model fitting, both effectively averaging the multi-cast εf(T, σ, t) data in a balanced way prior to final model fitting. Even so, a criterion of Z ≤ 7 appeared to be more reasonable as a pragmatic model-fit criterion of acceptability. However, subsequently it became apparent that Z ≤ 2 is still
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attainable for strains in the secondary to tertiary creep regimes, even for such complex alloy datasets when the assessment involves a quantitative consideration of cast-to-cast variability, for example, by strength compensation (Bolton, 2005a).
102
Strain (%)
10
H1 (CT) H2 (CT) H3 (CT) H4 (CT) H5 (IT) H6 (IT)
1
170
165 150 140 164 140 180
120 115 100
90 80
1 80
72
56 50
10–1
10–2 101
1
102
103
104
105
Time (h) (a) 102
125 190
75 100 90 80 78 98 123
63125 50 98 90 62 49
165
Strain (%)
101
62
40 1
123b 123e 123m 123o
10–1
10–2 10–1
1
101
102 Time (h) (b)
103
104
105
14.5 Examples of the distribution (at one temperature only) of large multi-source/multi-cast multi-temperature datasets typical of those used for determining long time creep strength values for European Product Standards, (a) P91 at 600°C, and (b) TP316 at 700°C. (CT means continuous measurement test, IT means interrupted measurement test.)
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14.3.2 Model selection The selection of constitutive equation and model-fitting approach can depend on a number of factors including material characteristics, data distribution and practical application. Material characteristics The effectiveness of a creep equation in representing a material’s characteristic ε(t) curve shape can depend on features such as the relative proportions of primary, secondary and tertiary creep as fractions of the strain and time at rupture (e.g. Fig. 14.6(a)) and the way in which they vary over the T, σ regime of interest (Fig. 14.6(b)). The Z-parameter (Equation [14.1]) provides a useful indication of model effectiveness in these circumstances. Data distribution Model selection and the choice of model-fitting approach can be influenced by the distribution of the data to be assessed. A creep dataset typically consists of a number of ε(t, T, σ) curves (creep test records), each comprising a number of ε(t) observations, for example Fig. 14.5. Both the ε(t, T, σ) curve and ε(t) distribution characteristics are influential (see Section 14.3.1). Practical application Model selection can also depend on the purpose for which the material’s creep strain description is required. Typically, the priority of the scientist is for a creep constitutive equation to have a sound physical basis in the primary, secondary and tertiary regimes. As a generality, it is more important for design and assessment engineers for the constitutive equation to be simple to implement and effective in its description of creep deformation at long times. For design engineers, effective modelling is more important in the relatively low strain primary/secondary creep regimes whereas for remaining life assessment engineers, the priority is more likely to be an accurate knowledge of secondary/tertiary behaviour.
14.4
Predicting service life
There is no universally preferred constitutive equation for predicting service life. In practice, the selection often depends on which model best represents the high temperature deformation characteristics of the material and the preference of the analyst and/or the requirements of the available application
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1.0
Type-316 650°C
Normalised plastic strain
0.8
0.6
0.4 21/4CrMo 540°C 0.2 Steel-91 600°C 0.0 0.0
0.2
0.4 0.6 Normalised time (a)
0.8
1.0
0.4 0.6 Normalised time (b)
0.8
1.0
1.0 450C, 520MPa 450C, 430MPa 650C, 120MPa 650C, 80MPa
Normalised plastic strain
0.8
0.6
0.4
0.2
0.0 0.0
0.2
14.6 Examples of creep curve shape variations for (a) 21/4CrMo, Steel-91 and Type-316 at typical application temperatures, and (b) Steel-91 over a wide T, σ range.
tools. An approach for testing model-fitting effectiveness is described in the previous section. The life assessment of most engineering structures involves consideration of the multiaxial stress state at a critical location and it is conventionally
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assumed that multi-axial creep deformation is characterised by the vonMises effective stress, σ VM . For a given σ VM , multiaxial loading does not exert a strong influence on deformation characteristics during primary and secondary creep (Fig. 14.7), but has a strong effect on tertiary creep deformation behaviour and time to rupture (Dyson and Osgerby, 1993) and in particular the magnitude of the rupture ductility (Rice and Tracey, 1969; Cocks and Ashby, 1980; Shammas and Marchant, 1986). σ VM = 1 2
[( σ 1 – σ 2 ) 2 + ( σ 2 – σ 3 ) 2 + ( σ 3 – σ 1 ) 2 ]
[14.2]
When the end-of-life criterion is rupture and the critical location in the component is subject to multi-axial loading, the simple use of σ VM as the representative stress is not always appropriate (see Section 14.4.2).
14.4.1 End-of-life criteria Constitutive creep equations may be used to predict service life directly in a defect-free design assessment of a component subject to steady loading at high temperature. In such circumstances, an important consideration is the end-of-life criterion. For most engineering applications, this is unlikely to be rupture (or the attainment of rupture strain as determined in a uniaxial test). More usually, the end-of-life is a strain limit, for example to avoid loss of clearance or interference during service duty (Bolton, 2005b). Alternatively, it is a stress-state dependent rupture strain (Rice and Tracey, 1969; Cocks
T, σVM
Torsion
Creep strain
Uniaxial tension Multi-axial
Time
14.7 Schematic representation of influence of stress state on creep deformation behaviour.
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and Ashby, 1980; Spindler, 2003), with an applied safety factor, or a conservative estimate of uniaxial rupture strain such as that provided by the Monkman–Grant relationship ε˙ f,min t u (Monkman and Grant, 1956).
14.4.2 Multi-axial stress rupture When rupture (or a safe fraction of the rupture time) is the end-of-life criterion, one approach for the determination of service life in components subject to multiaxial loading is to use an appropriate creep constitutive equation first to determine the steady-state stress distribution. With a knowledge of the multiaxial stress rupture criterion obeyed by the material, an appropriate representative stress can then be computed for critical locations in the component (see below). Ultimately, the service life may be based on uniaxial tu(T, σ) creep-rupture relationships (e.g. ECCC, 2005a) for which the stress input is the calculated representative stress state. The following examples do not provide an exhaustive list of representative stress formulations (see also Webster et al., 2004). Equations [14.3] to [14.6] are respectively due to Sdobyrev (1958), Othman and Hayhurst (1993), Cane (1979) and Huddleston (1985): σrep = ασ1 + (1 – α) σ VM (with 0 ≤ α ≤ 1)
[14.3]
σrep = ασ1 + 3βσm + (1 – α – β) σ VM (with 0 ≤ α + β ≤ 1) [14.4] σ rep = ( σ 1 / σ VM ) λ / n σ VM
(with 0 ≤ γ ≤ n)
σ rep = σ VM exp [ C (3σ m / Ss – 1)] (with s s =
[14.5] σ 12
+
σ 22
+
σ 32
) [14.6]
where α, β and γ are parameters reflecting the multi-axial rupture criterion obeyed by the material and are determined from the results of uniaxial and multi-axial rupture tests, see for example Webster et al., 2004). In Equation [14.6], C is a material constant (e.g. 0.24 for austenitic stainless steels). Such formulations acknowledge that creep rupture may be controlled by the maximum principal stress, σ1, the von-Mises effective stress, σ VM , and/or the hydrostatic stress, σm. Creep life assessment can also be determined by analysing the stress– strain state of the structure in its entirety using detailed finite element solutions with demanding material property input data requirements or by using more simplified approaches involving some form of reference stress, e.g. skeletal point stress (Calladine, 1969), equivalent yield stress (Goodall et al., 1979), structural stress (Bolton, 2005b). The application of such reference stress approaches is usually limited to creep ductile materials for which both deformation and failure are controlled by σ VM .
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14.4.3 Defect assessment Service lives determined on the basis of a high temperature defect assessment procedure also require the implementation of appropriate creep constitutive equations, e.g. (R5, 2003). For assessments involving the application of high temperature failure assessment diagrams (HTFADs), or creep crack initiation or creep crack growth rate model equations, it is first necessary to determine the steady-state stress distribution and creep deformation response at critical locations in the component.
14.5
Future trends
The continuously increasing availability of higher performance, lower cost computer technology will lead to the wider practical application of the most effective constitutive equations for predicting service life. In principle, there will be no limit to the complexity of the constitutive equations which can be implemented. In practice, the most commonly adopted equations will be those requiring readily available and/or determinable stress and temperaturedependent material property parameters. There will continue to be research effort aimed at developing constitutive equations that require results from short duration creep tests to predict long time service lives. While this is a laudable objective, it will require more vision and advanced computing power than is currently being applied to address the difficulty of reliably extrapolating beyond the mechanism regime(s) for which there exist experimentally determined material property data.
14.6
Concluding remarks
The chapter has reviewed the formulation of a wide range of creep constitutive equations currently adopted for scientific research and predicting service life. No single constitutive equation effectively represents the creep deformation of all materials over their entire temperature application range. An approach for quantifying model-fitting effectiveness has been reviewed. Important considerations for the prediction of service life at high temperatures are the selection of representative stress to characterise deformation and rupture due to multiaxial loading in the component material, and the end-of-life criteria. A number of options have been examined.
14.7 ECCC MSRC n
Nomenclature European Creep Collaborative Committee multiaxial stress rupture criterion stress exponent
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Q R Rpε/t/T, Ru/t/T
activation energy for creep universal gas constant creep strength and rupture strength for a given time and temperature standard deviation of all residual log times sA-RLT t time observed time to rupture, maximum observed time to rupture tu, tu,max observed and predicted times to given plastic strain tpε/σ/T, t*pε/σ/T time to a specific creep strain as a function of temperature tεf(T,σ) and stress temperature, melting temperature of material T, Tm Z Parameter quantifying effectiveness of master creep equation to predict times to specific strains (see Equation [14.1]) strain, elastic strain, instantaneous plastic strain ε, εe, εi creep strain, plastic strain, anelastic strain, permanent strain εf, εp, εk, εper ˙ε , ε˙ f,min , ε˙ ave strain rate, minimum creep strain rate, average strain rate creep rupture ductility εu σ stress friction stress σi representative rupture stress σrep principal stresses (where σ1>σ2>σ3) σ1, σ2, σ3 hydrostatic stress, i.e. σm = (σ1 + σ2 + σ3)/3 σm σ VM Von-Mises effective stress stress to give a specific strain as a function of time and σε(t, T) temperature ω , ω˙ damage, rate of damage accumulation ECCC terms and terminology recommendations are given in Volume 2 (ECCC, 2005b)
14.8
References
Akulov N S (1964), ‘On dislocation kinetics’, Acta Metall., 12, 1195. Baker A J and O’Donnell M P (2003), ‘R5 high temperature structural integrity assessment of a cracked dissimilar metal weld vessel test’, in Proceedings 2nd International Conference on Integrity of High Temperature Welds, 10–12 November 2003, IOM and I. Mech. E, London. Bartsch H (1995), ‘A new creep equation for ferritic and martensitic steels’, Steel Res., 66 (9), 384–388. Bolton J L (2005a), ‘A ‘characteristic-strain’ model for creep’, in Proceedings ECCC Creep Conference on Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, Shibli I A, Holdsworth S R and Merckling G, (eds) I. Mech. E., 12–14 September 2005, London, 465–477. Bolton J L (2005b), ‘Analysis of structures based on a characteristic-strain model of
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creep’, in Proceedings ECCC Creep Conference on Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, Shibli I A, Holdsworth S R and Merckling G, (eds), I. Mech. E., 12–14/9/05, London, 1032–1045. Calladine C R (1969), ‘Time scales for redistribution of stress in creep of structures’, Proc. Roy. Soc. London., A309, 363–375. Cane B J (1979), ‘Creep cavitation and rupture in 21/4CrMo steel under uniaxial and multiaxial stresses’ in Proceedings 2nd International Conference on Mechanical Behaviour of Materials, Miller K J and Smith R F (eds) Pergamon Press, Oxford, 2, 173–182. Cocks A C F and Ashby M F (1980), ‘Intergranular fracture during power-law creep under multi-axial stress’, Met Sci., 14, 395–402. Conway J B and Mullikin M J (1962), ‘An evaluation of various first stage creep equations’, in Proceedings AIME Conference, Detroit, Michigan. Creep of Steels Working Party (CSWP) (1983), High Temperature Design Data for Ferritic Pressure Vessel Steels, Institute of Mechanical Engineers, London. Dyson B F and McClean M (1998), ‘Microstructural evolution and its effects on the creep performance of high temperature alloys’, in Strang A and McClean M (eds), Microstructural Stability of Creep Resistant Alloys for High Temperature Applications, IOM, 371–393. Dyson B F and Osgerby S (1993), Modelling and Analysis of Creep Deformation and Fracture in a 1Cr1/2Mo Ferritic Steel’, NPL Report DMM(A)116. ECCC (2005a), Data Sheets for Rupture Strength, Creep Strength and Relaxation Strength Values for Carbon–Manganese, Low Alloy Ferritic, High Alloy Ferritic and Austenitic Steels, Nickel Base Alloys and High Temperature Bolting Steels/Alloys, Robertson D G and Holdsworth S R (eds), ECCC(ETD) publishers. ECCC (2005b), Recommendations for Creep Data Validation and Assessment Procedures, Holdsworth S R, Brown T B, Buchmayr B, Bullough C K, Calvano F, et al. (eds) Vol. 1: Overview, Vol. 2: Terms and terminology, Vol. 3: Data acceptability criteria, Data generation, Vol. 4: Data exchange and collation, Vol. 5: Data assessment, Vol. 6: Characterisation of microstructure and physical damage for remaining life assessment, Vol. 7: Data assessment – creep crack initiation, Vol. 8: Data assessment – multi-axial, Vol. 9: Component assessment, ECCC(ETD) publishers. EN 10291 (2000), Metallic Materials, Uniaxial Creep Testing in Tension, Method of Test, European Norm. Evans R W and Wilshire B (1985), Creep of Metals and Alloys, Institute of Metals, London. Garofalo F (1965), Fundamentals of Creep and Creep Rupture in Metals, New York, Macmillan. Graham A and Walles K F A (1955), ‘Relations between long and short time properties of commercial alloys’, JISI, 179, 105–120. Granacher J, Möhlig H, Schwienheer M and Berger C (2001), ‘Creep equation for high temperature material’, in Proceedings 7th International Conference on Creep & Fatigue at Elevated Temperatures (Creep 7), 3–8/6/01, Tsukuba, NRIM, 609–616. Grounes M (1969), ‘A reaction rate treatment of the extrapolation methods in creep testing’, J. Basic Eng., Series D, Trans ASME. Goodall J W, Leckie F A, Ponter A R S and Townley C H A (1979), ‘The development of high temperature design methods based on reference stresses and design methods’, J. Eng. Mat. Tech., 101, 349–355. Holdsworth S R and Merckling G (2003), ‘ECCC developments in the assessment of
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creep-rupture data’, in Proceedings 6th International Charles Parsons Conference on Engineering Issues in Turbine Machinery, Power Plant & Renewables, Trinity College, Dublin, 16–18 September 2003. Holdsworth S R, Askins M, Baker A, Gariboldi E, Holmström S, Klenk A, Ringel M, Merckling G, Sandström R, Schwienheer M. and Spigarelli S (2005), ‘Factors influencing creep model equation selection’, in Proceedings. ECCC Conference on Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, 12–14 September 2005, eds. Shibli I A, Holdsworth S R and Merckling G, (eds), I. Mech. E., London. Holmström S and Auerkari P (2004), ‘Prediction of creep strain and creep strength of ferritic steels for power plant applications, in Proceedings Baltica Conference on Life Management and Maintenance for Power Plants, VTT Symposium 234, 8–10 June 2004, Espoo. Huddleston R L (1985), ‘An improved multi-axial creep-rupture strength criterion’, Trans ASME, J. Press. Vessel Technol., 107, 412–429. Jones D I G and Bagley D L (1996), ‘A renewal theory of high temperature creep and inelasticity’, in Proceedings Conference on Creep and Fracture: Design and Life Assessment at High Temperature, London, 15–17/4/96, MEP, 81–90, London. Kachanov L M (1958), ‘Time to failure under creep conditions’, Izv. Akad. Navk. SSR. Otd Teck. Nauk., 8, 26–31. Kachanov L M (1986), Introduction to Continuum Damage Mechanics, Kluwer Academic, Dordrecht. Li J C M (1963), ‘A dislocation mechanism for transient creep’, Acta Metall., 11, 1269. McHenry D (1943), ‘A new aspect of creep in concrete and its application to design’, Proc. ASTM, 43, 1069. McLean M (1980), ‘Friction stress and recovery during high-temperature creep: interpretation of creep transients following a stress reduction’, Proc. Roy. Soc. London, Series A, Mathematical Phys. Sci., 371 (1745), 279–294. McVetty P G (1933), ‘Factors affecting the choice of working stresses for high temperature service’, Trans ASME, 55, 99. Merckling G (2002), ‘Metodi di calcolo a confronto per la previsione dellulteriore esercibilità in regime di scorrimento viscoso’, in Proceedings Conference on Fitness for Service, Giornata di Studio CESI-CONCERT, Milan, 28 November 2002. Monkman F C and Grant N J (1956), ‘An empirical relationship between rupture life and minimum creep rate in creep-rupture tests’, Proc. ASTM, 56, 593–620. Nadai A (1938), ‘The influence of time upon creep, The hyperbolic sine creep law’, in Stephen Timoshenko Anniversary Volume, Macmillan, New York. Norton F N (1929), The Creep of Steel at High Temperature, McGraw-Hill. Othman A M and Hayhurst D R (1993), ‘Determination of large strain multi-axial creep rupture criterion using notched bar data’, Int. J. Damage Mech., 2, 16–52. Phillips F (1905),‘The slow stretch in india rubber, glass and metal wire when subjected to a constant pull’, Phil. Mag., 9, 513. Prager M (1995), ‘Development of the MPC Omega method for life assessment in the creep range’, ASME J. Pressure Vessel Technol., 117, May, 95–103. Rabotnov Yu N (1969), Creep Problems in Structural Members, North-Holland, Amsterdam. Rice J R and Tracey D M (1969), ‘On the ductility enlargement of voids in triaxial stress fields’, J. Mech. Phys. Solids, 17, 201–217. R5 (2003), Assessment Procedure for the High Temperature Response of Structures, Procedure R5 Issue 3, British Energy, Gloucester, UK.
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Sdobyrev V P (1958), ‘Long term strength of alloy EI-437B under complex stresses’, Izv. Akad. Nauk. SSR. Otd. Teck. Nauk., 4, 92–97. Shammas M S and Marchant K D (1986), ‘Torsion testing in an inert atmosphere’, in Techniques for Multiaxial Creep Testing, Gooch D J and How I M. (eds), Elsevier Applied Science, London. Spindler, M W (2003), ‘The multi-axial creep ductility of austenitic stainless steel’, Fatigue Fract. Eng. Mater. Struct., 27, 273–281. Webster G A, Holdsworth S R, Loveday M S, Nikbin K., Perrin I J, Purper H, Skelton R P and Spindler M W (2004), ‘A code of practice for conducting notched bar creep tests and for interpreting the data, Issue 3’, Fatigue and Fract. Eng. Mater. Struct., 27(4), 319–342.
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15 Creep strain analysis for steel B . W I L S H I R E and H . B U R T University of Wales Swansea, UK
15.1
Introduction
When a tensile stress (σ) is applied to a metal or alloy at a temperature (T), the total strain (εtot) which accumulates over a time (t) can be expressed as: εtot = ε0 + ε
[15.1]
where ε0 is the initial strain on loading and ε is the subsequent creep strain. The loading strain depends on stress and temperature as: ε0 = ƒ1(σ, T)
[15.2]
allowing the magnitude of ε0 to be discussed by reference to the values of Young’s modulus (E), the yield or 0.2% proof stress (σY) and the ultimate tensile stress (σTS) determined from high-strain-rate tests (~10–3s–1) at the creep temperature. Hence, ε0 is predominantly elastic when σ < σY (with ε0 = σ/E), whereas ε0 has elastic and significant plastic components when σ > σY. In contrast to ε0 (Equation [15.2]), ε varies with stress, temperature and time as: ε = ƒ2 (σ, T, t)
[15.3]
with the creep strain/time characteristics being dependent on the T/Tm ratio, where Tm is the absolute melting point. When T < 0.4Tm, the total creep strains are low and failure rarely occurs. Conversely, when diffusion can take place at around 0.4Tm and above, the creep strains can be large, leading to eventual failure. Thus, while creep at low temperatures is normally of little practical significance, creep and creep fracture are often the life-limiting phenomena during component service in power plant and other hightemperature applications. For design of large-scale components in power plant, a knowledge is usually required of stresses which the relevant steels can sustain at the operating temperatures without failure occurring in 100 000 h. To provide this information, major experimental programmes involving tests lasting up to 30 000 h or more are currently completed but, for many steel grades, results 421 WPNL2204
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have already been obtained for stress–temperature conditions giving failure times exceeding 100 000 h. To reduce the scale and cost of these programmes, reliable techniques must be evolved to allow accurate estimation of long-term properties by extrapolation of short-term measurements. However, confidence in any predictive methodology is improved when the analysis procedures have a sound theoretical basis. For this reason, a variety of theoretical and practical approaches to high-temperature creep and creep fracture are now summarized in relation to straightforward descriptions of the changes in creep strain/time behaviour as the stress and temperature conditions are altered (Equation [15.3]). In this way, a coherent foundation is laid for consideration of the factors affecting creep strain accumulation, as well as the times and strains to failure. The industrial implications of these concepts are then addressed by defining physically meaningful techniques for rationalization of the behaviour patterns observed for power plant steels, before introducing easily applicable extrapolation procedures for long-term design data prediction. To ensure that the approaches adopted and the results achieved are open to scrutiny, the present analyses use only the internationally respected data sets1,2 produced by the National Institute for Materials Science (NIMS), Japan. In particular, attention is focused on two 9% Cr steels, namely, Grade 91 (9 Cr–1Mo–V–Nb) and Grade 92 (9 Cr–0.5Mo–1.8W–V–Nb).
15.2
Creep-induced strain
Under a sustained tensile stress at temperatures above 0.4Tm, most metals and alloys exhibit normal creep strain/time curves. Thus, following the initial loading strain, the creep strain rate ( ε˙ = dε/dt) decreases continuously with time during the primary stage, reaching a minimum or secondary rate ( ε˙ m ) before accelerating during the tertiary stage which leads to fracture after a time (tf). The product, ε˙ m t f , is often but not always a constant (M), signifying that creep failure is strain controlled because tf increases as ε˙ m falls with decreasing stress and temperature. Numerous relationships have been proposed to quantify the variations in creep strain with time. Several of these equations seek to describe only the early stages of the creep curves, while others attempt to define the shape of the entire ε/t trajectories, but no agreement has been reached on the relationships which should be used. Despite the distinctive shape of normal curves, it has therefore become common practice to ignore the primary and tertiary stages, assuming that the secondary rate remains constant with increasing time and strain. Equation [15.3] then reduces to: ε˙ m = ƒ3 (σ, T)
[15.4]
with a further simplification giving:
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ε˙ m = ƒ4 (σ)ƒ5(T)
423
[15.5]
so the variables are treated as separate and independent. In Equation [15.5], at a fixed temperature, the stress dependence of ε˙ m can be defined as: ε˙ m ∝ ƒ4 (σ) ∝ σn
[15.6]
where n is the stress exponent. Alternatively: ε˙ m ∝ ƒ4 (σ) ∝ exp σ
[15.7]
giving an exponential dependence of ε˙ m on stress. Conversely, at a fixed stress, the temperature dependence of ε˙ m in Equation [15.5] is generally represented by an Arrhenius equation of the form: ε˙ m ∝ ƒ5 (T) ∝ exp (–Qc/RT)
[15.8] –1
where Qc is the activation energy for creep in units of J mol when the gas constant, R = 8.314 J mol–1 K–1. Combining Equations [15.5], [15.7] and [15.8] then gives: M/tf = ε˙ m = B exp [–(Qc – V σ)/RT]
[15.9]
where B and V are treated as constants. However, in most theoretical and practical studies carried out over the last half century, Equations [15.5], [15.6] and [15.8] have been combined to obtain the standard power law relationship: M/tf = ε˙ m = A σn exp (–Qc/RT)
[15.10]
but the values of the parameter, A, as well as n and Qc, vary in different stress/temperature regimes.
15.2.1 Parametric approaches to data analysis Although NIMS Creep Data Sheet No. 43 (1996) details only the stress rupture properties,1 results available from other sources3,4 allow the creep lives to be considered in relation to the creep rate characteristics of tube samples of Gr. 91 steel. Thus, using Equation [15.10], the log ε˙ m /log σ plots in Fig. 15.1 can be represented3 by a set of straight lines showing a decrease from n ≅ 16 at 848 K to n ≅ 9 at 923 K. Similarly, the stress/creep life relationships determined over extended stress ranges at 773–973 K for multiple batches of Gr. 91 tube1 reveal gradient changes corresponding to decreases from n ≅ 17 to n ≅ 4.5 with increasing temperature (Fig. 15.2). With Qc ranging from 600 to 700 kJ mol–1, these anomalously large values of n and Qc are typical of the behaviour patterns reported for power plant steels and other particle-hardened alloys.
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Minimum creep rate (s–1)
10–5
10–6
10–7
10–8 848K : n ≅ 16 873K : n ≅ 11 898K : n ≅ 10 923K : n ≅ 9
10–9
10–10 75
100
150 Stress (MPa)
200
300
15.1 Stress dependence of the minimum creep rates recorded for Gr. 91 tube steel3 at 848–923 K. 500
773K 823K 873K 923K 973K
300
Stress (MPa)
200
100
50
20 106
107 Time to fracture (s)
108
15.2 Stress dependence of the creep lives recorded for multiple batches of Gr. 91 tube steel1 at 773–973 K.
Because of the complex stress and temperature dependences of ε˙ m and tf (Figs. 15.1 and 15.2), estimation of long-term properties by extrapolation of short-term measurements involves the continued use of various parametric methods introduced in the 1950s.5–7 These empirical approaches define ‘correlation parameters’, incorporating both creep life and temperature, which can be plotted as functions of stress to superimpose multi-batch results onto a single ‘master curve’ for a given steel. Unfortunately, no one parametric
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method has proved capable of fitting the experimental data reported for the majority of power plant steels and, even when the best fitting procedure is selected, the accuracies achieved are not always satisfactory.8 One limitation inherent in parametric methods is linked to the ‘variable constants’ encountered with standard power law relationships (Equation [15.10]). For instance, Equation [15.10] can be re-written to give the parameter (POSD) proposed by Orr, et al.,7 as: POSD = log tf – (P1/T )
[15.11]
where P1 includes Qc. Hence, variations in Qc ensure that the superimposed parametric plots are non-linear. The curvatures of plots based on Equation [15.11] are not then removed by replacing Equation [15.10] with Equation [15.9]. In fact, Equation [15.9] can be rearranged to give the parameter (PLM) proposed by Larson and Miller5 as PLM = T(log tf + P2)
[15.12]
where P2 now contains Qc. Combined with the broad scatter bands generally associated with multibatch data sets, the unknown curvatures of traditional parametric plots limit extrapolation to only about three times the longest reliable test measurements available. For this reason, tests lasting up to 30 000 h and more must usually be completed to estimate 100 000 h rupture strengths.
15.2.2 Alternative procedures for data rationalization For pure metals, rearranging Equation [15.10] and using the activation energies expected for diffusion, the stress/creep rate relationships recorded at different temperatures are superimposed9 simply by plotting the dependences on (σ/ E) of the temperature-compensated creep rate, ε˙ m exp(Qc/RT). Similarly, the corresponding stress rupture properties are rationalized by plotting the temperature-compensated creep lives, tf exp (–Qc/RT), against (σ/E). However, this approach is not applicable to power plant steels, because minor variations in the thermomechanical processing conditions selected for component manufacture can alter the resulting microstructures. Elastic moduli are temperature dependent but do not vary markedly with changes in microstructure, whereas the creep and fracture properties of particlehardened alloys are both temperature and microstructure sensitive, as are σY and σTS. Hence, early rationalization procedures based on normalizing σ through E9 are up-dated by normalizing σ through σY or σTS.10–13 In this way, using the σTS values measured for each batch of Gr. 91 steel investigated,1 the multi-batch stress rupture data in Fig. 15.2 are superimposed in Fig. 15.3 using a modified power law expression:11–13 M/tf = ε˙ m = A*(σ/σTS)n exp (– Qc* / RT ) [15.13]
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where A* ≠ A and Qc* is obtained from the temperature dependence of ε˙ m at constant (σ/σTS) rather than at constant σ as in the determination of Qc in Equation [15.10]. From Fig. 15.3, Qc* = 300 kJ mol–1, a value close to that for lattice diffusion in the alloy steel matrix. With Gr. 91 steel, as with other metals and alloys,10–14 creep property sets are also rationalized effectively by normalizing σ through σY, so that (σ/σTS) can be replaced by (σ/σY) in Equation [15.13]. Irrespective of whether σY or σTS is chosen, the resulting ‘master curve’ in Fig. 15.3 is at least as impressive as that obtained using parametric methods. Moreover, with Equation [15.13], the empirical terms in parametric relationships (Equations [15.11] and [15.12] are replaced by physically meaningful properties, namely, a sensible activation energy and the measured σY and σTS values.
15.2.3 Interpretation of power law behaviour Equation [15.13] avoids the large and variable Qc values observed when data sets for Gr. 91 steel are described using Equation [15.10] (Fig. 15.2), but does not eliminate a decrease from n ≅ 20 to n ≅ 4 (Fig. 15.3), a trend generally expected to continue towards n ≅ 1 or less as the test duration and temperature increase. One early attempt to explain the anomalously large n values suggested15 that creep occurs not under the full applied stress (σ) but under a reduced stress (σ – σo), such that
0.8 0.6
σ/σTS
0.4
773K 823K 873K 923K 973K
0.2
0.1 10–15
10–14
10–13
10–12 10–11 10–10 tf exp (–Qc*/RT) (s)
10–9
10–8
15.3 Dependence of the temperature-compensated creep life on (σ/σTS), using the σT S values for each batch of Gr. 91 tube steel investigated,1 with Q *c = 300 kJ mol–1.
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ε˙ m ∝ (σ – σo)m
427
[15.14]
where m = 4, with σo now called a ‘threshold stress’. Comparing Equations [15.10] and [15.14], n ≅ m ≅ 4 when σo ≅ 0 or when σo ∝ σ, whereas n > m when σo is large. This approach has been widely applied to creep of particlehardened alloys,16 but little progress has been made because σo cannot be measured or predicted reliably. More commonly, n value variations have been interpreted on the basis that different creep mechanisms become dominant in different stress– temperature regimes. Thus, while the superimposed results in Fig. 15.3 suggest that the gradient changes continuously as (σ/σTS) decreases, such curves could be approximated by a series of straight line segments corresponding to n >> 4, n ≅ 4 and eventually n ≅ 1, with each gradient change linked to a mechanism transition. However, no agreement has been reached on the detailed processes involved and, from a practical viewpoint, current theories do not allow prediction of the creep and creep fracture properties of engineering steels. Moreover, accepting multi-mechanism concepts, analysis of results recorded in one mechanism regime would not allow prediction of properties in another regime, necessitating the completion of long-term test programmes. In this context, it therefore seems reasonable to consider whether standard power law relationships offer a valid basis for representation and interpretation of creep properties. With Equation [15.10], the fact that n and Qc are themselves functions of stress and temperature means that, in the simplification of Equation [15.4] to obtain Equation [15.5], the variables are not separate and independent. In addition, the assumption that the variations in creep strain with time, stress and temperature (Equation [15.3]) can be quantified adequately through the stress and temperature dependences of a secondary or ‘steady state’ creep rate (Equation [15.4]) is highly questionable. For these reasons, without invoking mechanism transitions, an alternative approach17 contends that the behaviour patterns displayed by power plant steels (Figs. 15.1, 15.2 and 15.3) are easily understood through Equation [15.3] by considering the systematic changes in creep curve shape which occur with increasing test duration and temperature.
15.3
Patterns of creep strain accumulation
Although conventional ε/t plots give the impression that normal creep curves exhibit clearly defined primary, secondary and tertiary stages prior to failure, this view is contradicted by the available experimental evidence. Instead, for most metals and alloys,4,10–14,17 plotting the variations of the tensile creep strain rate ( ε˙ ) with time or strain demonstrates that a minimum rate ( ε˙ m ) rather than a secondary or ‘steady state’ value is reached when the decaying
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primary rate is offset by the tertiary acceleration, as illustrated for Gr. 91 tube steel4 in Fig. 15.4. Hence, the creep rate at any instant ( ε˙ ) should be expressed as:
ε˙ = ε˙ p + ε˙ t
[15.15]
where the primary rate ( ε˙ p ) decays and the tertiary creep rate ( ε˙ t ) accelerates with time. The ε˙ m and tf data derived from sets of normal creep curves should then be interpreted in terms of the deformation mechanisms controlling strain accumulation and the damage processes causing the creep rate to accelerate during the tertiary stage. In line with Equation [15.15], inspection of the ε˙ /t trajectories in Fig. 15.4 emphasizes that much information is lost by ignoring the primary and tertiary characteristics when creep mechanism studies have traditionally focused on ‘steady state’ behaviour. Thus, most equations introduced to describe creep curve shapes feature a ‘steady-state’ term. Yet, by explicitly quantifying the decaying primary and accelerating tertiary components (Equation [15.15]), the θ Projection Concept17 provides accurate descriptions of the curve shape dependence on stress and temperature, allowing extrapolation of short-term data to provide long-term property estimates. 10,12,17 However, the comprehensive sets of high-precision constant-stress creep curves required to carry out valid θ analyses have been produced for relatively few materials. Even so, the dominant deformation and damage processes can be clarified simply by monitoring some basic quantities defining the shapes of normal creep curves. 10–5
200MPa
160MPa 140MPa
–6
10
Creep rate (s–1)
10–7 10–8 10–9 10–10 120MPa 10–11 110MPa 103
104
105 106 Time (s)
100MPa 107
108
15.4 Variations of the creep strain rate ( ε˙ ) with time (t) for tube samples of Gr. 91 steel4 over a stress range from 100 to 200 MPa at 873 K.
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15.3.1 Variations in creep curve shape With power plant steels, creep and stress rupture tests are generally performed under tensile stresses less than σY so, from Equation [15.1], the essentially elastic loading strain (εo) is small in relation to the total creep strain to failure, often termed the creep ductility (εf). In turn, εf can be regarded as the sum of the primary creep strain (εp) accumulated over the period from t = 0 to t = tm (where tm is the time to the minimum rate), plus the tertiary strain (εt), where εt = (εf – εp). Using these easily visualized quantities, several key trends become evident from the ε˙ /t trajectories in Fig. 15.4 as the stress is reduced from 200 to 100 MPa at 873 K for Gr. 91 steel.4 The εf value falls from over 0.3 towards 0.2 as the creep life increases above about 10 000 h at stresses below 140 MPa, accompanied by a gradual decrease from εp ≅ 0.03 towards εp < 0.01 with increasing test duration (Fig. 15.5). The decaying primary rate is then offset by the tertiary acceleration at a progressively earlier fraction of the creep life, so the tm/tf ratio decreases as tf increases. The creep curves therefore become increasingly tertiary dominated as the primary stage becomes less significant with decreasing applied stress. With Gr. 91 steel at 873 K (Fig. 15.5), the modest εp values fall gradually with decreasing stress, a trend which would be expected when creep occurs by movement of dislocations in the alloy steel matrix. This view is consistent with the present observation that Qc* = 300 kJ mol–1, (Fig. 15.3). In fact, for Gr. 91 and other power plant steels, detailed microstructural studies have indicated that dislocation processes are dominant under long-term test conditions4 and even under the low stress levels experienced during powerplant service.18 0.5 εf . εm · t f εp
0.4 0.3
0.04 0.1
0.02
0.0 80
100
120
140 160 180 Stress (MPa)
200
ε˙ m t f & εp
εf
0.06 0.2
0.00 220
15.5 Variations of the primary strain (εp), the creep ductility (εf) and the product, ε˙ mt f with stress at 873 K for Gr. 91 tube steel (from Fig. 15.4).
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While no major change appears to occur in the dislocation mechanisms controlling strain accumulation over extended stress–temperature ranges, several different processes can start the tertiary acceleration. These include cavity and crack development, neck formation and evolution of the microstructures of particle-hardened alloys. In addition, once the tertiary stage begins, more than one process can influence the subsequent rates of strain accumulation, while the damage mechanisms initiating tertiary creep and causing failure may also differ. The relative importance of different damage processes can then vary as the test conditions alter, causing the creep curve shape to change in the manner illustrated in Fig. 15.4.
15.3.2 Creep damage tolerance values An informative method of quantifying the tertiary characteristics is introduced because the product, ε˙ m t f , is linked to the creep ductility, εf, as: λ = ε f / ε˙ m t f
[15.16]
where λ is the creep damage tolerance parameter. Strictly, because damage development influences the tertiary not the primary stage, εf should be replaced by εt in Equation [15.16] when the primary strains are substantial.20,21 In practice, the contribution of εp to εf can usually be neglected because εf >> εp (Fig. 15.5), allowing λ to be calculated from the values of ε˙ m , tf and εf reported for power plant steels. The measured λ values then relate to the damage processes initiating tertiary creep. In many instances, the tertiary acceleration has been modelled by defining a damage parameter (ω), which is zero for undamaged material.22 The value of ω increases as the damage levels increase, so: 19
ε˙ t = ε˙ o (1 + ω )
[15.17]
with the rate of damage accumulation ( ω˙ ) increasing as ε˙ t accelerates from an initial value ( ε˙ o ) when t = 0. Adopting this approach, modelling exercises23 have predicted that λ ≅ 1.5 to 2.5 when tertiary creep and fracture are attributable to cavitation, with higher values expected when the tertiary stage begins as a consequence of necking (with λ > 2.5) or precipitate coarsening (with λ > 5). Unfortunately, experimental studies covering the tertiary behaviour of various metals and alloys20,21 have shown that the λ value alone does not allow unambiguous identification of the dominant damage mode. When λ progressively exceeds about 1.5 with increasing test duration and temperature, the tertiary acceleration is due to microstructure evolution, for example the strength loss associated with precipitate coarsening. In contrast, when λ ≅ 1.5, tertiary creep can begin by intergranular or transgranular cracking
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431
and/or necking, which must be distinguished through use of additional approaches such as microstructural studies, profiling of fractured testpieces, and so on. As evident from Fig. 15.5, although εf decreases, the accompanying fall in ε˙ m t f means that an increase from λ ≅ 5 towards λ > 15 occurs as the test duration increases from about 40 to almost 40 000 h at 873 K. This result, as well as the patterns of strain accumulation for Gr. 91 steel (Fig. 15.4), supports the outcomes of microstructural investigations4 showing that the tertiary acceleration starts through evolution of the initial tempered martensite structure. Although dislocation processes appear to govern creep strain accumulation at all stress levels, the creep curve shape changes because microstructure evolution causes the minimum rate to occur at a progressively earlier fraction of the creep life with increasing test duration and temperature (Fig. 15.4). The gradual decrease in n value (Figs. 15.1, 15.2 and 15.3) is then attributable, not to creep mechanism transitions, but to the complex dependence of ε˙ m and tf on the systematic variations in the creep curve shape as (σ/σTS) decreases. However, the curve shape variations may also be accompanied by changes in the creep ductility (εf), as indicated in Fig. 15.5. Under tensile creep conditions, the creep life can be considered conveniently as the time taken for the creep strain to reach the limiting creep ductility, that is, t → tf as ε → εf. An abrupt fall in εf during long-term exposure could then result in a reduction in the failure times expected by projection of short-term data. Hence, in seeking valid extrapolation procedures, it is necessary to clarify the manner in which εf varies over stress–temperature ranges leading to fracture in times up to 100 000 h or more. Yet, compared with the emphasis placed on measurement and interpretation of ε˙ m and tf data, comparatively little attention has been devoted to the factors affecting creep ductility (εf) and the related reduction of cross-sectional area at fracture (RoA).
15.3.3 Creep ductility With power plant steels at a fixed creep temperature, εf is often observed to decrease with increasing test duration, as illustrated in Fig. 15.5, but detailed ductility trends may be masked by the broad scatter bands encountered with multi-batch εf measurements. Hence, the εf and RoA values recorded for Gr. 91 tube samples1 are plotted, not against stress (Fig. 15.5), but against the temperature-compensated creep life, tf exp (– Qc* / RT ) , with Qc* = 300 kJ mol–1 (Fig. 15.6). Although a fall in εf is not immediately apparent because of scatter, there is an obvious decrease in the corresponding RoA values with increasing test duration and temperature (Fig. 15.6). In fact, this RoA decrease occurs with failure times of less than 100 000 h at 883 K, the upper service temperature for Gr. 91 steel.
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εf and RoA
0.8
0.6
0.4
773K 823K 873K 923K 973K
0.2 Open symbols: εf Closed symbols: RoA
0.0 10–17 10–16 10–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8 tf exp (–Qc*/RT) (s)
15.6 Changes in creep ductility (εf) and reduction of area at fracture (RoA) as a function of the temperature-compensated creep life, tf exp (– Q *c /RT ) , with Q *c = 300 kJ mol–1, using data obtained1 for multiple batches of Gr. 91 tube steel at 723–923 K.
The decrease in εf as tf increases (Fig. 15.5) has been linked to the loss of creep strength caused by progressive evolution of the martensitic microstructure, leading to the suggestion that extrapolation of short-term tf measurements overestimates long-term performance.4 Noting a related change in the stress dependence of the creep lives, it has been recommended that results obtained under the higher stress ranges covered at each test temperature should be excluded from data projection exercises,8,24 making long-term test programmes mandatory. However, this conclusion appears to be negated by several features of the ductility characteristics displayed by Gr. 91 steel: •
•
•
Under all test conditions investigated,1 the results presented in Fig. 15.6 reveal that the measured reduction of area at fracture (RoA) is significantly greater than the creep ductility (εf). Fracture is therefore preceded by neck formation, indicating that failure occurs in a relatively ductile manner, even during long-term creep exposure. As evident from Fig. 15.4, even in the test lasting almost 40 000 h at 873 K, the creep rate is accelerating rapidly late in the creep life. With most of the creep strain accumulating just prior to failure, only a major decrease in εf would cause a marked reduction in tf. The magnitude of λ (Equation [15.16]) is important in practical situations when high strains develop in regions, say, where a change in component cross-section leads to stress concentrations. With λ values of 5 or more considered to be adequate, the present estimates suggest that the localized stress concentrations typically encountered during service should not lead to premature cracking of Gr. 91 pipework, even though problems are being experienced with type IV failure of weldments.25
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These observations suggest that, despite the property trends shown in Fig. 15.6, the failure mode is essentially unchanged over extended stress– temperature ranges. Hence, with no definitive evidence for transitions in the dominant creep mechanism, there seems to be no a priori reason why analysis of short-term measurements should not allow prediction of long-term properties. Indeed, the procedures shown to rationalize multi-batch tf values (Figs. 15.2 and 15.3) introduce straightforward extrapolation methods for estimation of 100 000 h design data.12,13
15.4
Practical implications of creep strain analysis
Even without conceiving new extrapolation procedures, replacing Equation [15.10] by Equation [15.13] allows standard multi-batch tf measurements (Fig. 15.2) to be superimposed onto master curves (Fig. 15.3), suggesting that the total number of tests required to supply 100 000 h creep rupture estimates can be reduced substantially.10 Thus, tf measurements could be made for one batch of steel at the planned service temperature for stresses causing failure in times up to 100 000 h. By determining σTS at different temperatures for several batches, only a limited number of additional tests would be required to confirm that an activation energy of 300 kJ mol–1 allows effective data rationalization. Standard stress rupture plots (Fig. 15.2) can then be computed from the resulting ‘master curves’ (Fig. 15.3). However, not only the number of tests but also the maximum test durations can be reduced through alternative relationships12,13 introduced to describe the dependence of the temperature-compensated creep lives on (σ/σTS).
15.4.1 Creep life extrapolation for Gr. 91 steel With σTS representing the maximum stress which can be applied at the creep temperature, the failure times approach zero as σ → σTS, whereas infinite creep lives must be recorded when σ → 0. These criteria are met by replacing Equation [15.13] with a modified version of Equation [15.9],12,13 giving:
( σ / σ TS ) = exp {– k1 [ t f exp (– Qc* / RT )]u}
[15.18]
where k1 and u are evaluated by plotting ln [ t f exp (– Qc* / RT )] against ln[–ln(σ/σTS)], with Qc* = 300 kJ mol–1, as illustrated in Fig. 15.7. In this case, k1 and u were determined for results acquired1 only under test conditions giving creep lives of less than 5000 h, although it is evident from Fig. 15.7 that virtually identical values would be derived by including all measurements presented in Fig. 15.3. Using Equation [15.18], the ‘master curve’ in Fig. 15.8 then demonstrates that tf → 0 as (σ/σTS) → 1 and tf → ∞ as (σ/σTS) → 0, with a point of inflection occurring at about 0.5 σY, providing an impressive description of the stress rupture properties reported1 for Gr. 91 steel.
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ln [tf exp (–Qc*/RT)] (s)
–20 –22
k1 = 26.37 u = 0.134
–24 –26 –28
773K 823K 873K 923K 973K
–30 –32 –34 –1.5
–1.0
–0.5 0.0 ln [–ln(σ/σTS)]
0.5
1.0
15.7 Dependence of ln [tf exp (– Q *c /RT ) ] on ln [−ln (σ/σTS)], using the σTS values recorded for each batch of Gr. 91 tube steel investigated,1 with Q *c = 300 kJ mol–1. Test conditions where tf < 5000 h are shown as closed symbols, while results for tf > 5000 h are included as open symbols.
0.8
773K 823K 873K 873K 923K 973K
0.6
σ/σTS
434
0.4
0.2
0.0 10–16 10–15 10–14 10–13 10–12 10–11 10–10 10–9 10–8 10–7 tf exp (–Qc*/RT) (s)
15.8 Dependence of log [tf exp (– Q *c /RT ) ] on (σ/σTS), using the σTS values recorded for each batch of Gr. 91 tube steel investigated,1 with Q *c = 300 kJ mol–1. Results for tests carried out at stresses greater than 0.5 σY are presented as closed symbols and for tests at stresses less than 0.5 σY as open symbols, with the behaviour pattern predicted from Equation [15.18] shown as a solid line.
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Knowing k1 and u (Fig. 15.7), plus the σTS values for the different batches of Gr. 91 tube,1 Equation [15.18] also allows tf values to be computed over extended stress ranges at various creep temperatures. These predicted curves are included in Fig. 15.9, together with recently reported multi-batch tf values, produced through NIMS for test conditions giving creep lives up to 100 000 h for Gr. 91 plates, pipes, forgings and tubes.26 Clearly, the current projections derived using Equation [15.18] to describe tf values for tests lasting less than 5000 h fit well with the general property trends in Fig. 15.9, despite the inevitable scatter in the long-term measurements. However, to assess the predictive accuracy of Equation [15.18], noting that the present estimates were obtained for Gr. 91 tube samples only, the predictions from Fig. 15.9 are compared in Table 15.1 with recent 100 000 h rupture strength estimates from other sources.26–28 Several features of the results in Table 15.1 then merit comment:
•
In two cases,26,27 the estimates in Table 15.1 were obtained using the Larson–Miller relationship (Equation [15.12]). Both analyses highlight the fact that very different strength estimates are produced, depending on the values chosen for the empirical constant. In Table 15.1, the 100 000 h rupture strengths were estimated by Kimura26 using the latest long-term NIMS data for Gr. 91 steel (Fig. 15.9). Yet, contrary to recent suggestions,4,8,24 the property projections now derived by applying Equation [15.18] to the short-term NIMS measurements Time to fracture (h) 101 102 103
100
104
105
500 300 200
Stress (MPa)
•
100
50
20 103
773K 823K 873K 923K 973K 104
105 106 107 Time to fracture (s)
108
109
15.9 Stress/creep life behaviour predicted from Equation [15.18] for Gr. 91 tube steel (solid lines) compared with long-term multi-batch stress rupture data recorded26 for plate, pipe, forging and tube samples at 773–973 K.
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•
Creep-resistant steels
(i.e. tf < 5000 h) slightly underestimate rather than overestimate the results achieved by excluding the high-stress tf values at each test temperature from the extrapolation exercise based on the long-term properties.26 Given the spread in the 100 000 h strength values obtained using parametric methods (Table 15.1), the present predictions appear reasonable, agreeing well with the outcomes of extensive long-term test programmes28 undertaken through the European Collaborative Creep Committee (ECCC).
Having considered the creep and creep fracture properties of Gr. 91 steel, as a further check on the general applicability of Equation [15.18], the same data analysis methods are now adopted to discuss the stress rupture behaviour of Gr. 92 steel. Again, for this material, microstructure studies combined with parametric extrapolation exercises have suggested that the inclusion of results from test lasting less than about 3000 h cause overestimation of longterm performance.29,30
15.4.2 Creep life extrapolation for Grade 92 steel The data sets reported for Gr. 92 tube2 essentially replicate the behaviour patterns shown for Gr. 91 tube in Figs. 15.1 and 15.2, with Equation [15.13] again superimposing the stress/creep life relationships at different test temperatures, in line with Fig. 15.3. Hence, to evaluate the predictive capabilities of Equation [15.18], the stress–rupture properties are plotted in Fig. 15.10 to show the variation of ln [ t1 exp (– Qc* / RT )] as a function of Table 15.1 100 000 h creep rupture estimates (MPa) for Gr. 91 steel at 823, 873 and 923 K Temp (K)
823 873 923
Present estimates
154 87 43
Ref 26a
150 98 42
Ref 27b
Ref 28c
A
B
A
B
132 83 48
159 100 56
153 86 46
150 85 44
a
Estimates obtained using the Larson–Miller relationship to analyse data produced at stresses less than 180, 130 and 90 MPa at 823, 873 and 923 K, respectively whereas, by describing the full data sets from tests lasting up to almost 100 000 h, significantly larger estimates were determined for the 100 000 h strengths. b Estimates obtained using the Larson–Miller relationship with constants of 20 (Set A) and 36 (Set B). c Estimates obtained for tubes and pipes, with the longest test having a creep life exceeding 110 000 h under a stress of 150 MPa at 823 K (Set A). Set B quotes the 2005 estimates produced through the European Collaborative Creep Committee (ECCC).
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ln[–ln(σ/σTS)]. In contrast to the single-line fit achieved for Gr. 91 (Fig. 15.7), it is evident from Fig. 15.10 that the values of k1 and u in Equation [15.18] change as the temperature-compensated creep life increases with the Gr. 92 samples. Specifically, linear extrapolation of the results obtained under high stresses at low temperatures overestimates the creep rupture strengths under low stresses at high temperatures (Fig. 15.10). These changes in k1 and u appear to support the proposal26,29,30 that shortterm tf values should not be included when seeking to predict long-term stress-rupture properties. However, it is apparent from Fig. 15.10 that the long-term behaviour of Gr. 92 steel can be determined through Equation [15.18] using only the data recorded under stress–temperature conditions giving maximum creep lives of around 5000 h or so. In fact, the k1 and u values in Fig. 15.10 were calculated on this basis. Essentially, the tf values determined at 823 and 873 K represent the failure properties when (σ/σTS) is high, while the results at 973 and 1023 K describe the stress rupture characteristics when (σ/σTS) is low. In agreement with the results presented for Gr. 91 steel (Fig. 15.8), Fig. 15.11 then shows that incorporating the derived k1 and u values into Equation [15.18] provides a sensible description of the tf data for Gr. 92 steel, that is, tf → 0 as (σ/σTS) → 1, again with a point of inflection at σ ≅ 0.5 σY ensuring that tf → ∞ as (σ/σTS) → 0. Knowing the k1 and u values (Fig. 15.10), Equation [15.18] also allows prediction of the trends in the log tf/log σ relationships obtained2 from tests lasting up to 30 000 h at different temperatures for multiple batches of Gr. 92 tube and pipe (Fig. 15.12). In addition, as found for the Gr. 91 material –16
ln [tf exp (–Qc*/RT)] (s)
–18 –20 –22 –24
k1 = 37.96 u = 0.16
k1 = 8.35 u = 0.097
–26
823K 873K 923K 973K 1023K
–28 –30 –32 –1.0
–0.8 –0.6 –0.4 –0.2 0.0 0.2 ln [–ln(σ/σTS)]
0.4
0.6
0.8
15.10 Dependence of ln[tf exp (– Q *c /RT ) ] on ln[–ln(σ/σTS)], with Q *c = 300 kJ mol–1, using data reported2 for Gr. 92 tube steel. Results are presented for creep lives up to about 5000 h or so as closed symbols and for longer term data as open symbols.
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σ/σTS
0.6
0.4
823K 873K 923K 973K 1023K
0.2
Gr.91
0.0 10–15 10–14 10–13 10–12 10–11 10–10 10–9 tf exp (–Qc*/RT) (s)
10–8
10–7
15.11 Dependence of log[tf exp (– Q *c /RT ) ] on (σ/σTS)] using the σTS values recorded for each batch of Gr. 92 tube steel investigated,2 with Q *c = 300 kJ mol–1. Results for tests carried out at stresses greater than 0.5 σY are presented as closed symbols and for tests at stresses less than 0.5 σY as open symbols, with the behaviour pattern predicted from Equation [15.18] shown as a solid line. For comparison purposes, the behaviour predicted for Gr. 91 tube steel (Fig. 15.8) is included as a broken line.
Time to fracture (h) 102 103 104
101
105
500
200
Stress (MPa)
438
100 50
20 10 104
823K 873K 923K 973K 1023K 105
106 107 Time to fracture (s)
108
109
15.12 Stress/creep life behaviour predicted from Equation [15.18], using the k1 and u values derived in Fig. 15.10 (solid lines), compared with the multi-batch stress rupture data recorded2 at 823–1023 K for tube and pipe samples of Gr. 92 steel.
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(Table 15.1), the present estimates of the 100 000 h creep rupture strengths at 823, 873 and 923 K are certainly reasonable in relation to the ranges of values derived27,29 when the detailed analysis procedures differ using the Larson–Miller relationship (Table 15.2). However, the current predictions are of particular interest when compared with the strength estimates reported for Gr. 92 steel over the last decade. In 1995, a 100 000 h strength value of 132 MPa at 873 K was determined through Larson–Miller assessments31,32 of results from 284 tests at 823 to 1023 K, with the longest test duration being 43 513 h at 873 K. Using these and additional results, four different analysis methods were then applied33–36 to 704 creep life measurements, giving revised 873 K strength values of 116 to 129 MPa, with a figure of 123 MPa presented in the 1999 ECCC data sheet. In 2005, again adopting several different procedures to analyse over 800 data points recorded at 823 to 973 K (with a maximum creep life exceeding 110 000 h), the 873 K strength estimate was further reduced30 to within the range 107 to 118 MPa (Table 15.2). This progressive reduction in the 100 000 h strength values seems to confirm the importance of long-term test measurements, simultaneously justifying the omission of short-term results from the data analysis exercises. Yet, the present predictions obtained by applying Equation [15.18] only to tf values up to around 5000 h or so are in excellent agreement with the strengths recently quoted30 for creep lives of 100, 1000, 10 000 and 100 000 h at 823, 873 and 923 K in Table 15.3. On this basis, rather than focussing attention on extending the test durations and selective editing of short-term measurements, it seems that the parametric procedures still widely adopted for data analysis should be seriously reevaluated. Table 15.2 100 000 h creep rupture estimates (MPa) for Gr. 92 steel at 823, 873 and 923 K Temp (K)
823 873 923
Present estimates
182 104 53
Ref 29a
Ref 27b
Ref 30c
A
B
A
B
183 112 51
192 128 70
141 91 52
186 120 70
a
187 to 190 107 to 118 53 to 62
Using the Larson–Miller relationship to describe the stress rupture data recorded in tests giving creep lives up to 70 000 h, the estimated 100 000 h creep rupture strengths derived by excluding tf measurements less than 3000 h (Set A) are substantially lower than the predictions obtained by analysing all results available (Set B). b Estimates obtained using the Larson–Miller relationship with constants of 20 (Set A) and 36 (Set B). c Estimates obtained using several different procedures to analyse extensive data sets, with a maximum creep life of 110 450 h at 873 K. See Table 15.3.
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Table 15.3 Comparisons of the present predictions obtained using Equation [15.18] with estimates of the stresses (MPa) causing creep failure of Gr. 92 steel in 100, 1000, 10 000 and 100 000 h at 823–923 K derived by applying several different analysis procedures to long-term data30 Temperature (K)
Analysis method
823
Graphical ISO 6303 Larson–Miller Present estimate Graphical ISO 6303 Larson–Miller NIMS Present estimate Graphical ISO 6303 Larson–Miller NIMS Present estimate
873
923
Creep rupture strength (MPa) for times of 100 h
1000 h
10 000 h
100 000 h
276 284 294 272
251 254 257 245
221 223 221 215
190 188 187 182
209 213 223
183 184 186
202 147 153 157
177 122 121 122
144
121
149 150 151 150 148 90 86 89 86 86
107 111 118 109 104 53 56 62 55 53
15.4.3 Appraisal of proposed data analysis procedures With both Gr. 91 and Gr. 92 steels, microstructural studies4,26,28,29 have established that the patterns of creep strain accumulation and failure are strongly influenced as the tempered martensite structures of the as-received batches evolve with increasing time and temperature. Yet, while the overall property trends seem similar (Figs. 15.8 and 15.11), the effects of the strength loss caused by microstructure evolution appear to be more severe with Gr. 92 than with Gr. 91 steel, as reflected by the fact that the k and u values change in Fig. 15.10 but not in Fig. 15.7. These data comparisons then indicate that, using Equation [15.18], detailed differences in long-term behaviour can be diagnosed by analysis of short-term property sets. Given the scale of attention currently being directed to identification of reliable methodologies for prediction of long-term stress rupture properties,33–41 it is therefore useful to consider several key features of the approaches now proposed for data rationalization and extrapolation. •
By plotting the temperature-compensated creep lives as a function of either (σ/σTS) or (σ/σY), stress rupture measurements obtained over broad stress–temperature ranges are superimposed onto well-fitted ‘master curves’, using the activation energy for lattice diffusion in the alloy steel matrices (i.e. Qc* = 300 kJ mol–1). In this way, data rationalization can
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•
•
441
be achieved through Equation [15.13] (Fig. 15.3) or Equation [15.18] (Fig. 15.8). In either case, without assuming distinctive mechanism transitions, the results can be explained in terms of the dislocation processes controlling creep strain accumulation and the damage processes causing the tertiary acceleration and eventual failure. Avoiding the problems associated with the choice of the empirical constants in parametric relationships (Tables 15.1 and 15.2), the coefficients (k1 and u) in Equation [15.18] are easily determined (Figs. 15.7 and 15.10), allowing reasonable prediction of long-term stress rupture properties by analysis of tf measurements from tests lasting only about 5000 h (Figs. 15.9 and 15.12). Of course, with any forecasting method, the accuracy of the predictions must inevitably improve as the extent of the extrapolation decreases. Because forecasts are generally based on continuity of trends, the shorter the extent of the extrapolation, the lower is the risk of unforeseen events influencing the predictions. Even so, the use of short-term data offers advantages with steels prone to severe oxidation, that is the predictions are not significantly affected by the stress intensification associated with the loss of cross-sectional area during long-term creep exposure.10 Inspection of the data sets superimposed using Equation [15.18] reveals that a point of inflection in the sigmoidal curves occurs at around 0.5 σY for both the Gr. 91 and Gr. 92 steels (Figs. 15.8 and 15.11). This observation coincides with the proposal by Kimura8 that only data determined at stresses less than 0.5 σY should be incorporated in extrapolation exercises based on traditional parametric relationships. However, the general applicability of the 0.5 σY inflection in Figs. 15.8 and 15.11, as well as its interpretation, would require analysis of data sets for a range of other steels.
In seeking to validate the proposed rationalization and extrapolation procedures, the means of assessment are already available through the NIMS Creep Data Sheets. These comprehensive reports detail not only the stress rupture properties but also a wide range of other relevant information, including the measured σY and σTS values for each batch of steel investigated. Additionally, these documents should even allow evaluation of other data analysis concepts now suggested for future study.
15.5
Future data analysis options
Earlier investigations12,13 have demonstrated that, as with the rationalization and extrapolation of stress rupture properties through Equation [15.18], ε˙ m data recorded over extended stress-temperature ranges can also be analysed as:
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( σ / σ TS ) = exp {– k 2 [ ε˙ m exp ( Qc* / RT )] v}
[15.19]
with the coefficients (k2 and v) again determined by plotting the temperature compensated creep rate as a function of either (σ/σTS) or (σ/σY). The sigmoidal curves produced through Equation [15.19] then satisfy the essential criterion that ε˙ m → ∞ as tf → 0 when (σ/σTS) → 1, with a point of inflection again ensuring that ε˙ m → 0 as tf → ∞ when (σ/σTS) → 0. In a similar manner, the times to reach certain pre-defined creep strains (tε) can be described as: (σ/σTS) = exp {–k3 [tεexp (– Qc* / RT ) ]w}
[15.20]
with a set of k3 and w values quantifying the behaviour at different strains. Provided that full creep curves are recorded, a series of tε exp (– Qc* / RT ) against (σ/σTS) plots at various creep strains must map onto the appropriate tf exp (– Qc* / RT ) against (σ/σTS) relationship when ε = εf so tε = tf (Figs. 15.8 and 15.11). Adopting these procedures, detailed computer-efficient descriptions of the changes in creep strain (or strain rate) with time, stress and temperature (Equation [15.3]) could be made available for incorporation into the modern finite element codes developed for high-temperature engineering design. Moreover, this form of property representation would link directly with plots of εf and RoA measurements as functions of either the temperature-compensated creep lives (Fig. 15.6) or the (σ/σTS) and (σ/σY) values, clarifying long-term creep ductility trends. Overall, these new approaches should then provide an effective development avenue for future analysis, interpretation and extrapolation of creep and creep fracture data for power plant steels and other creep-resistant alloys.
15.6
References
1 NIMS Creep Data Sheet No. 43, Data sheets on the elevated-temperature properties of 9Cr-1Mo-V-Nb steel tubes for boilers and heat exchangers (ASME SA-213/SA213M Grade T91) and 9Cr–1Mo–V–Nb steel plates for boilers and pressure vessels (ASME SA-387/SA-387M Grade 91), 1996. 2 NIMS Creep Data Sheet No. 48, Data sheets on the elevated-temperature properties of 9Cr-0.5Mo-1.8W-V-Nb steel tube for power boilers (ASME SA-213/SA-213M Grade T92) and 9Cr–0.5Mo–1.8W–V–Nb steel pipe for high temperature service (ASME SA-335/SA-335M Grade P92), 2002. 3 Spigarelli S, Cerri E, Bianchi P and Evangelista E, ‘Interpretation of creep behaviour of a 9Cr–Mo–Nb–V–N (T91) steel using threshold stress concept’, Mater Sci Tech, 1999, 15, 1433–1440. 4 Kimura K, Kushima H and Abe F, ‘Heterogeneous changes in microstructure and degradation behaviour of 9Cr–1Mo–V–Nb steel during long term creep’, Key Eng Mater, 2000, 171–174, 483–490.
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5 Larson F R and Miller J, ‘A time-temperature relationship for rupture and creep stresses’, Trans ASME, 1952, 74, 765–775. 6 Manson S S and Haferd A M, ‘A linear time-temperature relation for extrapolation of creep and stress rupture data’, NASA TN 2890, 1953. 7 Orr R L, Sherby O D and Dorn J E, ‘Correlations of rupture data for metals at elevated temperature’, Trans ASM, 1954, 46, 113–128. 8 Kimura K, ‘Present status and future prospect of NIMS creep data sheet’, Creep Deformation and Fracture, Design and Life Extension, R S Mishra, J C Earthman, S V Raj and R Viswanathan (eds), MS&T, Pittsburgh, 2005, 97–106. 9 Barrett C R, Ardell A J and Sherby O D, ‘Influence of modulus on the temperature dependence of the activation energy for creep at high temperatures’, Trans AIME, 1964, 230, 200–204. 10 Wilshire B and Burt H, ‘Yield stress rationalization of creep and creep fracture properties’, Scripta Mater, 2005, 53, 909–914. 11 Wilshire B and Burt H, ‘A unified theoretical and practical approach to creep and creep fracture’, Creep Deformation and Fracture, Design and Life Extension, R S Mishra, J C Earthman, S V Raj and R Viswanathan (eds), MS&T, Pittsburgh, 2005, 3–12. 12 Burt H and Wilshire B, ‘Theoretical and practical implications of creep curve shape analyses for 7010 and 7075’, Metall Mater Trans A, 2006, 37A, 1005–1015. 13 Wilshire B and Burt H, ‘Creep data rationalization for power plant steels’, Material Science Forum, 2007, 539–543, 254–261. 14 Wilshire B and Battenbough A J, ‘Creep and creep fracture of polycrystalline copper’, Mater Sci Eng A, 2007, 443, 156–166. 15 Williams K R and Wilshire B, ‘On the stress and temperature dependence of creep of Nimonic 80A’, Metal Sci J, 1973, 7, 176–179. 16 Arzt E, ‘Creep of dispersion strengthened materials: A critical assessment’, Res Mech, 1991, 31, 399–453. 17 Evans R W and Wilshire B, Creep of Metals and Alloys, Institute of Metals, London, 1985. 18 Williams K R and Wilshire B, ‘Effects of microstructural instability on the creep and fracture behaviour of ferritic steels’, Mater Sci Eng, 1977, 28, 289–296. 19 Leckie F A and Hayhurst D R, ‘Constitutive equation for creep rupture’, Acta Metall, 1977, 25, 1059–1070. 20 Wilshire B and Burt H, ‘Tertiary creep of metals and alloys’, Z Metallkd, 2005, 96, 552–557. 21 Wilshire B and Burt H, ‘Damage evolution during creep of steels’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech Publ, London, 2005, 191–200. 22 Kachanov L M, ‘On the time to failure under creep conditions’, Izv Acad Nauk USSR, Otd TeKd Nauk, 1957, 8, 26–38. 23 Ashby M F and Dyson B F, ‘Creep damage mechanics and micromechanisms’, Advances in Fracture Research, S R Valluri (ed.), Pergamon Press, Oxford, 1993, Volume 1, 3–30. 24 Foldyna V, Kuboň Z, Jakobová A and Vodarek V, ‘Development of advanced high chromium ferritic steels; microstructural development and stability’, High Chromium Ferritic Power Plant Steels, A Strang and D J Gooch (eds), Institute of Materials, London, 1997, 73–92. 25 Brett S J, Oates D L and Johnston C, ‘In-service type IV cracking in a modified 9Cr
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26
27
28
29
30
31
32 33 34 35 36
37
38
39
40
Creep-resistant steels (Grade 91) header’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech London, 2005, 563–572. Kimura K, ‘Review of allowable stress and new guideline of long-term creep strength assessment for high Cr ferritic creep resistant steels’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech London, 2005, 1009–1022. Masuyama F, ‘Creep rupture life and design factors for high strength ferritic steels’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech London, 2005, 983–996. Di Gianfrancesco A, Cipolla L and Cirilli F, ‘Microstructural stability and creep data assessment of Tenaris Grades 91 and 911’, 1st International Conference, Super-High Strength Steels, Rome, Italy, (AIM), 2005 CD-Rom. Ennis P, ‘The significance of microstructural changes and steam oxidation for the service life of chromium steel components’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech, London, 2005, Volume 2, 79–87. Bendick W and Gabrel J, ‘Assessment of creep rupture strength for the new martensitic 9% chromium steels E911 and T/P92’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DEStech London, 2005, 406–418. Naoi H, Mimura H, Ohgami M, Morimoto H, Tanaka T, Yazoki Y and Fujita T, ‘NF616 pipe production and properties and welding consumable development’, EPRI/ National Power Conference, London, 1995, 8–29. Masuyama F, ‘ASME code approval for NF616 and HCM 12A’, EPRI/National Power Conference, London, 1995, 98–113. Granacher J and Monsees M, ECCC DESA CRDA Procedure Document, ECCC Recommendations, 2001, Vol 5, Appendix D2. Holdsworth S R, Bullough C K and Orr J, BS PD6605 Creep Rupture Assessment Procedure, ECCC Recommendations, 2001, Vol. 5, Appendix D3. Orr J, ECCC ISO CRDA Procedure Document, ECCC Recommendations, 2001, Vol 5, Appendix D1a. Granacher J and Schwienheer M, ECCC Procedure Document for Graphical MultiHeat Averaging and Cross Plotting Method, ECCC Recommendations, 2001, Vol 5, Appendix D4. Merckling G, ‘Long term creep rupture strength assessment: the development of the European Collaborative Creep Committee post assessment tests’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 3–19. Yagi K, ‘Acquisitions of long term creep data and knowledge for new applications’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 31–45. Swindeman R W and Swindeman M J, ‘A comparison of creep models for nickel base alloys for advanced energy systems’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 361–371. Holdsworth S R, Askins M, Baker A, Gariboldi E, Holmstrom S, Klenk A, Ringel M,
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Merckling G, Sandstrom R, Schwienheer M and Spigarelli S, ‘Factors influencing creep model equation selection’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 380–393. 41 Abe F, ‘Stress to produce minimum creep rate of 10–5 %/h and stress to cause rupture at 105h for ferritic and austenitic steels and superalloys’, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, I A Shibli, S R Holdsworth and G Merckling (eds), DESTech, London, 2005, 997–1008.
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16 Creep fatigue behaviour and crack growth of steels C . B E R G E R, A . S C H O L Z, F . M U E L L E R and M . S C H W I E N H E E R, Darmstadt University of Technology, Germany
16.1
Introduction
The lifetime of components of power plants depends in most cases on variable loading conditions. This concerns fatigue and creep–fatigue as well as crack initiation and crack propagation. The normal variable service conditions which encompass phases of start-up, full load, partial load and shutdown cause variable stress–strain distributions and temperature transients and lead to a large variety of combined static (primary load) and variable loading (secondary load) situations. A wide range of loading parameters have been introduced in order to simulate complex loading of components. Conventional testing of temperature-induced loading deals with fatigue experiments representing the cyclic loading at the heated surface of components, while creep experiments represent the quasi-static loading phases of components. A key problem deals with the identification and interpretation of different physical damage mechanisms. Therefore, phenomenological solutions were traditionally introduced and put into practical use. Fatigue cracks occur at the surface of a component (Fig. 16.1), but creep damage is initiated by creep cavities and micro cracks preferably at grain boundaries. Their interaction was intended for conventional heat-resistant steels, but their consideration in life estimation methods has not been realised satisfactorily. A clear influence on endurances caused by the superposition of fatigue and creep has been observed.1 Increasing tensile hold time decreases the fatigue endurance limit. Thus, the failure mode changes from a fatigue dominant to a creep dominant condition. In the high strain regime, damage is characterized as fatigue dominated. With decreasing strain, creep damage predominates. It is known from the literature2 that hold times at the tension and compression stage can have an individual material influence on endurance. Additionally, oxidation effects can contribute significantly to the reduction of endurance limits.2 Summarizing, either fatigue damage or creep damage prevails, or they may interact under variations of strain range, tension and/or compression hold time, frequency, temperature and ductility of the material. 446 WPNL2204
(a)
(b)
Total strain range
Creep fatigue behaviour and crack growth of steels
447
Pur e fa tigu Cre ea ep– inte nd fati com rac g tion ue pre (b) ssi g ve in ll s dw Cr ells ea we Cre ee r p-d ep– (a) nc le d I i n f i om atig ter s act n fai u ina e e ion t lur ted (b) e( c)
Number of cycles to failure Nf
(c)
16.1 Failure modes at fatigue, creep fatigue and creep-dominated material behaviour of heat-resistant steels, schematically.2
The initial stage of fatigue failure is characterized by dislocation processes which lead to surface defects. This is followed by growth processes by bulk deformation and is completed by tearing up a small remaining ligament. Creep damage owing to nucleation is much more difficult to define. It is assumed that microstructure damage is nucleated during the creep life and can be interpreted as the onset of tertiary creep.
16.2
Creep–fatigue experiments
Generally, conventional mechanical experiments are needed to measure the material properties in order to provide a consistent basis for quality control purposes and for design data. In addition, complex experiments may represent service conditions and contribute to verification in life prediction methods. The creep–fatigue behaviour at the heated surface of heavy components, such as turbine rotors, normally is investigated by conventional low-cycle fatigue (LCF) experiments with standard (LCF) cycles (Fig. 16.2(a)) and creep–fatigue cycles with dwell periods at maximum and minimum strain (Fig. 16.2(b)). In contrast, a single-stage service-type strain cycle (Fig 16.2(c) to (e)) was developed,1,3 which is characterized by a compressive strain hold phase 1 simulating the start-up condition, a zero strain hold phase 2, with approximate temperature equilibration at constant loading, a tensile strain hold phase 3, simulating shut-down conditions and an additional zero strain hold phase 4, which characterizes a zero loading condition. Anisothermal experiments (type ‘an’) approximate the service conditions more closely than isothermal experiments. They were carried out up to failure times of 8000 h and gave only insignificantly smaller numbers of cycles to
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0
t
(b) T= const.
0
tht
1
ε˙ r
T
ε
Hold phase j
0
ε˙ r
ε3 = ∆ε/3
ε1 = –2∆ε/3
(c) ε
σ
(d)
ε˙ r = 6%/min ε3
2 Total strain range ∆ε
ε1
3
σeff σx
4 σi
ε˙ r
ε +
Cycle period tp
Tmax Temperature range ∆T
Tmin (e)
t
σ
σ
3 0
σmax
t 1
ε˙ r thc
Stress range ∆σ
σm
0
t t1 = 0.075 tp t2 = 0.700 tp t3 = 0.150 tp t4 = 0.075 tp
σm
+
ε ∆σeff
∆εeff
16.2 Different strain cycles simulating the conditions at the heated surface of heavy components. Standard cycle without (a) and with (b) hold times as well as service-type cycle (c), the stress–strain path (d) and (e) show mean stress σm, stress σmax, internal stress σi, effective stress σeff and corresponding effective values.
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(a) T= const. ε
Creep fatigue behaviour and crack growth of steels
449
failure than comparable isothermal tests (type ‘iso’). Considering the design life of power plants of up to 200 000 h or more, long-term strain cycling is of interest. This could be realised by an isothermal package-type testing procedure (‘pa’) (Fig. 16.3). It is composed of packages of strain cycles with short hold times which are periodically inserted into much longer creep packages. Maximum test durations of 70 000 h have been achieved.
16.3
Stress–strain behaviour
In order to develop a creep–fatigue life estimation procedure, the stress– strain path of the service-type experiments were analysed. Deformation analyses led to the determination of an effective stress concept, σeff = σ – σi (Fig. 16.4). The internal stress σi for any time of the measured hysteresis loop, for example point A, stress σx, is defined as the centre of the hypothetical elastic–plastic flank curve loop which is inserted in the flank curve loop enveloping the whole measured loop. The flank curves are derived from a cyclic or quasi-static yield curve by multiplying the latter by a factor of two. The cyclic yield curve can be experimentally determined by a strain cycle without hold times which is inserted into the service-type strain cycling. Owing to the different relaxation phases during the hold times, the mean stress σm varies from zero (Fig. 16.4) and has to be considered for the lifetime estimation. For this purpose a mean stress factor νσ which includes the Smith–Watson–Topper parameter4 can be used.
16.4
Creep–fatigue interaction, life estimation
For life estimation under creep fatigue loading the life fraction rule5,6 is widely used. Failure is determined by the summation of fatigue damage Lf as a cycle fraction and creep damage Lc as a time fraction up to a critical creep fatigue value L, whereby L depends on the material. In this paper L is used in the sense of a damage fraction. The damage mechanisms during creep–fatigue conditions are mainly influenced by microstructural coarsening7 and cavity growth. Typical mechanisms of cavity growth have been identified in 1CrMoNiV steels during 3
ε
∆ε
1
ε
4
2
tp
t
σ
3
t ∆ε 1 Strain cycle package
Service type cycle
2
t +
Creep package
n – times (n > 10)
16.3 Simulation of long-term isothermal service-type testing by isothermal package type tests.
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450
A
PSWT Measured loop
∆σeff σmax
σeff
Flank curve loop (twice the yield curve)
Nf σi
σm
ε
Inserted loop defining the internal stress σi of point A
PSWT =
σ max ⋅ ∆ε ⋅ E /2
N f ( σ max , ∆ε eff ) νσ = N f ( ∆σ eff /2,∆ε eff )
Course of internal stress σi ∆εelf
16.4 Presentation of a hysteresis loop of an isothermal service-type strain cycle according to Fig. 16.2(c) and definition of internal stress σi, mean stress σm and the mean stress factor νσ.
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σ A′
Creep fatigue behaviour and crack growth of steels
451
stress relaxation.8 Grain boundaries may be cavitated in tensile hold time, but cavitation is usually not found in a balanced cycle containing hold time of equal duration. In service-type strain cycling, (Fig. 16.2(c)) cavitation dominates. As indicated by Scholz and Berger,9 a high deformation rate can be associated with transgranular damage, while a low deformation rate leads to intergranular damage (Fig. 16.5). Cavities were found at the transition of transgranular damage to mixed-mode damage on a 1CrMoNiV steel at 525°C. This assumption was applied to an interaction concept developed for life analysis using the damage accumulation rule, at first for 1CrMoNiV steel (see e.g. Granacher et al.)10 (Fig. 16.6). The identification of a transition time ttr is addressed in the separation of dominant fatigue damage at region 1 and dominant creep damage at region 2. A popular rule for creep–fatigue life analysis in the long-term region is the generalized damage accumulation rule:
Σ Σ ( ∆t j / t uj ) + Σ ( N k / N fok ) = L k
j
[16.1]
k
which combines the Miner rule for fatigue damage and the life fraction rule for creep damage.1,3,8,9 The damage summation over all cycles k including damage at hold times j = 1 to 4 leads to creep–fatigue damage L. To calculate creep damage, the ratio of time increments tj and rupture time tuj is considered. The structure of this relationship is relatively simple and therefore in practical use for life monitoring systems.11 Some features of this rule were modified empirically in order to cover physical aspects of the material, for example cyclic stress–strain–time behaviour and mechanisms given in
Stress
Region 1 transgranular
Region 2 intergranular
Deformation cavity growth
ttr
Constrained diffusion cavity growth Time
16.5 Various mechanisms of cavity growth during tensile dwell in a strain controlled cycle, 1CrMoNiV steel,2 and determination of a transition time tr for application to damage accumulation.
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gran
ular
Mixed mode
Stress
Tran s
Intergranular
Time (a) ε ∆ε
Transgranular Stress
σ
ttr
Intergranular
ttr
t
tp
(b)
ttr
t
Time
16.6 Association between failure mode at creep rupture testing (a) and failure mechanisms at stress relaxation (b), steel of type 1CrMoNiV.
(Figs 16.4–16.7). The reference value of the number of cycles to failure Nfo: Nfo = Nf (∆ε, 2tHsta) · νσ
[16.2]
is taken from standard strain-cycling tests, with tension and compression hold times tHsta (Fig. 16.2(b)). In the case of a viscoelastic stress–strain path, at low strain amplitudes where plastic deformation disappears, creep damage dominates owing to intergranular damage (region 2, Figure 16.5) and fatigue damage was calculated by a fatigue life curve (tp = 0 h). In the full elastic– plastic regime at high values of total strain range, creep–fatigue damage dominates owing to transgranular damage (region 1, Fig. 16.5). Here fatigue damage is calculated on the basis of a failure life curve based on symmetrical creep–fatigue experiments with symmetric hold times (Figure 16.2b). The hold time tH1 = tHsta at compression strain is fully considered as creep– fatigue damage while creep damage is derived from hold phase 2 and 4 and the remaining time of hold phase 3 (tH3 – tHsta).
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The reference value of rupture time tu (eq. (1)) is taken from a creep rupture curve for time dependent stresses σeff(t) and is also affected by the hold times tHsta (Fig. 16.2(b)); Fig. 16.7). Thus, cyclic softening effects are taken into account. In addition, preloading effects can have an influence on the reference values of Nfo and tuo.1,7 Furthermore, Nfo depends on a mean stress factor νσ (Fig. 16.4): νσ = Nf (∆σmax, ∆εeff)/Nf (∆σeff/2, ∆εeff)
[16.3]
This factor νσ is derived from the Smith–Watson–Topper parameter (PSWT)4 and considers a mean stress σm ≠ 0 in the cycle. Herein the values σmax, ∆εeff and ∆σeff are taken from the flank curve loop (Fig. 16.2(e) and Fig. 16.4). The main result of a creep-fatigue life analysis as described above is a creep–fatigue damage mean value L . Mean values of L = 0.54 at 500°C and L = 0.52 at 525°C were identified for the conventional European ferritic 1CrMoNiV rotor steel.1,3 Higher mean values were obtained for the martensitic 12CrMoV steel, where L = 0.75 at 550°C and L = 0.93 at 600°C are reported in1,3. Finally, for the modern 600°C steel of type 10CrMoWVNbN and its cast version, mean values L = 0.71 and L = 0.65 were found at 600°C.3 As a ε
tH3
tH2
0
tH4
σ
Pure elastic
t
∆ε
tH1
Transition regime
Elastic plastic
0
t Creep damage
∆ε
∆ε<0.36%
ttr
Creep fatigue damage
0.36<∆ε<0.44%
0.0
∆ε>0.44%
0.0 3/3
0/0 3/3
tH1/tH1
Nf Fatigue
Creep fatigue 3/3 min
σ
σ
σ ε
Creep fatigue tH1/tH1
ε
ε
tp = 2 · tHsta = 2 · ttr
16.7 Association of elastic–plastic deformation and failure mechanisms at stress relaxation for the calculation of creep damage and fatigue damage.
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Creep-resistant steels
result, all features introduced within this creep–fatigue interaction concept yielded the smallest scatter band in the creep fatigue life diagram (Fig. 16.8). Three-stage service-type strain cycling (Fig. 16.9) demonstrates typical service loading conditions as cold start, warm start and hot start. Such loading 1
Mean value
∑∆t/tuo
L
0.1
10CrMoWVNbN - forged steel 10CrMoWVNbN – cast steel 9CrMoVNb - pipe steel 0.01 0.01
100 000
0.1 ∑N /Nfo (a)
1
T = 600°C Single-stage cycling single heat No 1A grade
Nf cycle
10 000
Three-stage cycling single heat No 1A grade
Extrapolated
ε˙ r = 0.06%/min 1000
100 100
Material: heat No 1A: X12CrMoWVNbN10-1-1 grade: 10%CrMo(W)VNbN 1000
N ** f (b)
10 000
100 000
16.8 Results of creep–fatigue life assessment of service-type strain cycling experiments, (a) for 10CrMoWVNbN forged and cast steels and 9CrMoVNb pipe steel and (b) measured number of cycles to failure Nf of isothermal service-type strain cycling tests versus predicted number of cycles to failure N *f* ; single heat and material grade of 10%CrMo(W)VNbN.3,12
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H
C
T
C
∆TH = 50°C
∆TW = 100°C
Tm
Tmin Frequency: 1 Coldstart, 3 Warmstarts, 16 Hotstarts
Collective (1 + 3 + 16) · tc H3
H3
ε
H4
C1
C2
tC
W3
C3
H3
H4
W4
H2
∆εH
H1 W1
C example for elementary cycles CH HH HH
HH
t
W2
H2 H1
C4
HW
WH …
∆εH
∆εW H2.4 W2 H
W Rainflow – counting Rangemean – counting
16.9 Three-stage service-type strain cycling and principle of cycle counting.1,10
W3
∆εC
…
C HC
Creep fatigue behaviour and crack growth of steels
∆TC = 300°C
455
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Creep-resistant steels
sequences are of specific design interest. The frequency is typical for a medium loaded power plant. Three-stage cycling tests have reached longest times to failure up to 1.5 × 104 h for 1CrMoNiV steel and 104 h for 10CrMoWVNbN steel. The creep–fatigue damage assessment of the three-stage service-type experiments leads to values L within the scatter band of the single-stage service-type tests (Fig. 16.8). This important result confirms the ability of the damage accumulation rule (Equation [16.1]) for multi-stage service-type loading. In order to transfer knowledge of deformation behaviour and damage assessment obtained in several research programmes, a software tool SARA was developed for lifetime studies and industrial application as a life-cycle counter in power plants.3,10 The numerical simulation by SARA envelopes the input of cycle and material data, the synthesis of stress–strain hysteresis loops according to cycle counting methods, the individual assessment of fatigue damage, creep–fatigue damage and finally an output of the predicted number of cycles-to-failure N f** , failure time t f** as well as cycle specific results. Life estimation on the basis of single heat data and material grade data leads to an acceptable result (Fig. 16.8).
16.5
Multiaxial behaviour
Life estimation concepts either of a conventional type or of a advanced type require suitable multiaxial experiments for verification purposes. In addition to tension/torsion or internal pressure experiments developed in the past, experiments with cruciform test pieces (Fig. 16.10) are of great interest. Investigations by Ohnami and co-workers13 have demonstrated the large influence of loading ratio on number of cycles to crack initiation. Pure fatigue tests on a 1%Cr-steel show a factor of 10 in life between the biaxial strain ratio Φε = –1 and Φε = +1 (Fig. 16.11(a), solid lines). Long-term service-type creep-fatigue experiments on a 1%CrMoNiV rotor steel with four hold times (cycle period tp) are performed in the cruciform testing system. Experiments run under strain-controlled mode with long hold times. The strain ratio is given as Φε = 0.5 and Φε = 1. A total of five biaxial service-type creep–fatigue experiments were performed on the 1%CrMoNiV rotor steel with relevant test durations of up to about 2000 h. As a first result at strain ratio Φε = 1, a clear influence of superimposed creep at hold times of a factor of two can be observed. Secondly, a strain ratio of Φε = 0.5 leads to an increase of the number of cycles to crack initiation Ni of a factor 1.5 compared to Φε = 1. This result confirms the pure fatigue biaxial experiment results.13 Further, biaxial strain-controlled experiments lead to a reduction of number of cycles to crack initiation Ni up to a factor 3 compared to uniaxial service-type creep fatigue experiments (Fig. 16.11(b)). On the one hand, an increase of total strain range leads to a larger difference in Ni
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PEEQ (Ave. Crit.: 75%) +2.00e–03 +1.50e–03 +1.00e–03 +5.00e–04 +0.00e+00
B TE
ε A, ε B A TE
y x
16.10 Scheme of the cruciform specimen (a), and elastic–plastic finite element calculation showing equivalent plastic strain εp distribution (b); maximum deformation in the test zone (TE thermocouple).
Creep fatigue behaviour and crack growth of steels
(a)
457
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Total strain range ∆εx(%)
0.6 Φε –1.0 –0.5 0.0 0.5 1.0
0.4
0.2
0.1 102
Ohnami Φε = 1.0, Φε = 1.0, Φε = 0.5, Φε = 1.0, Φε = 1.0,
tp tp tp tp tp
= = = = =
1.33 1.23 1.33 3.53 3.43
h, h, h, h, h,
t1 t1 t1 t1 t1
≈ ≈ ≈ ≈ ≈
530 h, ∆εx = 0.60% 2055 h, ∆εx = 0.42% 830 h, ∆εx = 0.60% 1016 h, ∆εx = 0.60% 2230 h, ∆εx = 0.42%
103 104 105 Number of cycles to crack initiation Ni (cycles) (a)
1 0.8 0.6
Total strain range ∆εx(%)
458
0.4
0.2
0.1 102
Uniaxial service-type, tp = 1.0 h Uniaxial service-type, tp = 3.2 h Φε = 1.0, tp = 1.33 h, t1 ≈ 530 h, ∆εx = 0.60% Φε = 1.0, tp = 1.23 h, t1 ≈ 2055 h, ∆εx = 0.42% Φε = 1.0, tp = 3.53 h, t1 ≈ 1016 h, ∆εx = 0.60% Φε = 1.0, tp = 3.43 h, t1 ≈ 2230 h, ∆εx = 0.42% 103 104 105 Number of cycles to crack initiation Ni (cycles) (b)
16.11 Influence of biaxial strain ratio Φε on the number of cycles to crack initiation Ni in fatigue testing, 1%CrMoV, T = 550°C13 and first results of long-term service-type creep–fatigue experiments (Fig. 16.2(c)) (Φε = 1.0 and Φε = 0.5), dε/dt = 0.06% min–1 (a), comparison with uniaxial experiments (b), tp periodic time, ti time to crack initiation corresponding to a crack depth of about 0.2 mm; 1%CrMoNiV, T = 525°C.
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values, but on the other hand, increasing hold time has a more significant influence on number of cycles to crack initiation Ni. Results of experiments with cruciform specimens can be employed in combination with finite element (FE) simulations in order to verify various material models. These simulations were performed with ABAQUS whereby a constitutive material model is implemented in a user-defined subroutine UMAT. A constitutive material model describes elastic–viscoplastic behaviour for small deformations and was introduced by Tsakmakis and Reckwerth.14 A key feature is the combination of effective stress with a generalized energy equivalence principle. An undamaged fictitious material is described by means of effective variables which are the basis of the constitutive material model. The structure of this model can be attributed to Lemaitre and Chaboche.15 A damage variable D is defined by an approach proposed by Lemaitre additionally for another set of variables for the damaged real material. The known behaviour of the undamaged fictitious material is then mapped to the unknown behaviour of the real material with damage. This step is done by substitution of the effective variables using relations which implicate the damage variable D. Within current research work on creep–fatigue, the material parameters of the constitutive model were determined by a two-step approach using a combination of the neural networks method and the optimization method by Nelder–Mead.29 The neural networks method is already established in similar non-linear problems and can deliver a ‘global’ solution. The method by Nelder–Mead is a direct search method without the need for numerical or analytical gradients and leads only to a ‘local’ solution. This method is commonly referred to as unconstrained non-linear optimization. In the first step, the neural network identifies a parameter vector close to the global solution within a parameter interval. This result is used subsequently in the second step as an initial parameter vector in the Nelder–Mead method for further improvement of the solution. In order to identify the parameters of the model under examination for 1%CrMoNiV steel by means of neural networks, one-dimensional (1D) calculations depending on parameter variations were performed. The example in Fig. 16.12 demonstrates the applicability of the material model introduced.
16.6
Creep and creep–fatigue crack behaviour
In the case of unavoidable notches, creep–fatigue cracks may be initiated or propagated by static and/or cyclic high temperature loading. A quantitative description is required in order to establish crack initiation and propagation methodologies. For a description of crack behaviour based on experiments on standard compact–tension specimens16 and/or non-standard test specimens17 under creep–fatigue conditions, the parameter C* and the stress intensity factor KI
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Φε = 1.0, tp = 1.33 h ∆εx = ∆εx = 0.60%
Strain (%)
0.3
0.0
–0.3 Experiment Simulation
–0.6 0.00
0.25
0.50
0.75 1.00 Time (h)
1.25
1.50
16.12 Finite element simulation of biaxial service-type creep–fatigue experiments (see Fig. 16.11(a)), comparison of strain versus time curves from experiment to those of FE model calculated by the constitutive material model, and the material parameters determined, Φε = 1.0, 1%CrMoNiV steel, T = 525°C, dε/dt = 0.06% min–1.
can be applied to heat-resistant steels. The creep fracture mechanics parameter C* is valid for stationary creep in the crack tip environment.16 The stress intensity factor KI is valid only for linear elastic behaviour,18 but it can be used as an approximation if the plastic zone, near the crack tip, is limited.19 A model for creep crack growth was proposed which assumed that crack advance occurs when the creep ductility is exhausted at the crack tip.20 The creep crack growth rate under steady state conditions may be written as: da/dtNSW = [(n + 1)/ ( Au* ) ] · (A/rc)[1/(n+1)](C*/In)[n/(n+1)]
[16.4]
where rc is the size of the creep process zone (usually related to the grain size of the material), In represents a non-dimensional function and Au* is the appropriate (multiaxial) crack tip creep ductility. This model is known as the Nikbin–Smith–Webster model (NSW model). For ductile steels it has been found that most experimental data approach the Nikbin-Smith-Webster plane stress prediction.21 The creep crack growth rate is most sensitive to the multiaxial creep ductility, Au* . Therefore, the steady state creep crack growth rate, da/dtNSW-A, may be approximated as:22 da/dtNSW-A = 3·(C*)0.85/ Au*
[16.5] –1
–1
where da/dt and C* have the units of mm h and MPa m h , respectively and Au* is taken as the uniaxial failure strain, Au, for plane stress conditions and Au/30 for plane strain.23 This model has been validated for a range of materials.21 If ∆a is the minimum crack extension that can be measured reliably, then the initiation time, ti, may be estimated by:
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ti = ∆a/(da/dt)
461
[16.6]
If the approximate Nikbin–Smith–Webster Equation [16.5] is used, then Equation [16.6] becomes: t i = ∆a ⋅ Au* /[3 ⋅ ( C *) 0.85 ]
[16.7]
Predictions from Equation [16.7] will vary under conditions of plane stress and plane strain. In the following, the crack initiation is defined by a constant technical crack initiation length ∆ai = 0.5 mm which is independent of specimen geometry and size. Other definitions are possible, for example a crack initiation length, which is dependent on geometry and size. The parameter C* is shown in Fig. 16.13 against the creep crack initiation time ti for the 10CrMoWVNbN cast steel at 550°C and 600°C. In a shortterm regime there is a difference between small scale and large scale specimens. In a long-term regime the influence of specimen geometry and size on crack initiation time becomes smaller. The plane strain and plane stress initiation time predictions, given by Equation [16.7] are included in Fig. 16.13. The plane strain Nikbin–Smith–Webster A model provides a very conservative estimate of the time to creep crack initiation. Applying the formula for plane stress, a better agreement between creep crack initiation data and the approximate Nikbin–Smith–Webster model can be observed. The creep crack initiation behaviour of the 10CrMoWVNbN cast steel in terms of KI is shown in Fig. 16.14. In this figure the Larson–Miller parameter PLM,24 known from the description of creep data, is used as a time–temperature
C* (N mm–1 h–1)
101 Specimen/a0/W/T/environment Cs25/0.55/550°C/air Cs25/0.55/660°C/air Cs25/0.55/600°C/shielded gas Ds60/0.20/550°C/air Ds60/0.20/600°C/air
10–1
ti = (∆a· Au* )/(3·(C*/1000)0.85) Au* = Au (NSW-A-stress) Au* = Au/30 (NSW-A-strain
10–2
10–3
10–4 10–5 101
10CrMoWVNbN-cast steel T = 550 and 600°C ∆ai = 0.5 mm 102
103
104
105
t i (h)
16.13 Creep crack initiation time ti versus parameter C*; 10CrMoWVNbN cast steel, 550°C and 600°C.
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106
462
Creep-resistant steels
75
Specimen/a0 /W/T/environment
KIi (MPa m1/2)
50
Ds60/0.20/550°C/air Ds60/0.20/600°C/air
10CrMoWVNbN-cast steel T = 550 and 600 °C ∆ai = 0.5 mm
25
10 7.5
5 11
Specimen/a0 /W/T/environment Cs25/0.55/550°C/air Cs25/0.55/600°C/air Cs25/0.55/600°C/shielded gas 12
13
14 15 PLM = T · [13.0 + log(t)]
16
17
16.14 Stress intensity factor KIi versus Larson–Miller–Parameter PLM; 10CrMoWVNbN cast steel, 550°C and 600°C.
parameter in order to get a temperature-independent representation. There is a significant difference between small scale and large scale specimens. The creep crack initiation on large scale specimens occurs later than on small scale specimens. It is recommended that side-grooved C(T) specimens (B = 25 mm) be used for creep crack tests.16 Results from these specimens lead to a conservative estimation of creep crack initiation with regard to large specimens or components. In the two-criteria-diagram19 the nominal stress σnpl is related to the stress situation in the ligament, that is in the far field of the creep crack and the fictitious elastic parameter KI at time zero characterizes the crack tip situation. These loading parameters are normalized in a two-criteria-diagram by the respective time- and temperature-dependent values, which indicate material resistance against crack initiation. The normalized parameters are the stress ratio Rσ = σnpl/Ru/t/T for the far field and the stress intensity ratio RK = KI/KIi for the crack tip, with creep rupture strength Ru/t/T and parameter KIi which characterize the creep crack initiation of the material. This parameter has to be determined from specimens with high ratio KI/σnpl, preferably side grooved C(T)25-specimens. The two-criteria-diagram (Fig. 16.15) distinguishes three fields of damage mode separated by lines with a constant ratio Rσ/RK. Above the line Rσ/RK = 2.0 ligament damage, and below the line Rσ/RK = 0.5 crack tip damage is expected. Between these lines a mixed damage mode is observed. Crack initiation is only expected above a boundary line. In order to transfer
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Cr
Techn. crack initiation *) With
1%
1%
CT100, a/W = 0.55 CT50, a/W = 0.4 – 0.55 DENT, B = 60 mm, a/W = 0.4 DENT, B = 60 mm, a/W = 0.2 DENT, B = 60 mm, a/W = 0.1
1%
Small specimens, CT25, a/W = 0.55 DENT, B = 9-15 mm, a/W = 0.2 - 0.4
M
ix
Rσ = σnpl/Ru / t /T
ed
da
m
ag
e
Ligament damage
Burst 1
Castings with manufacturing defects Small manufacturing defects in smooth specimens
1%
Small specimens, CT25, a/W = 0.55 DENT, B = 9-15 mm, a/W = 0.2 - 0.4
=2
*0.1 - 0.5 mm crack length No crack
Rσ /
Rk
0.5
R
/R k σ
.5 =0
Leakage Crack tip damage
0 0
0.5
1
1.5 RK = Kl/Kli
2
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16.15 Two-criteria-diagram for creep crack initiation; 1 Cr and 12 Cr steels, 550°C. DENT represents the double edge notched tension specimen. CT represents the compact tension specimen.
Creep fatigue behaviour and crack growth of steels
1%
Without
464
Creep-resistant steels
creep crack initiation data from specimens with different sizes to large components with a similar far field and crack tip situation, the two-criteriadiagram has been developed. The applicability of the two-criteria-diagram was validated by results of more than 100 small and large scale specimens with artificial and natural defects. Creep–fatigue experiments, with increased hold time tH up to 3.0 h, show a significant decrease of crack initiation time (Fig. 16.16). With increasing hold times the data points approach pure creep crack behaviour. Short hold times lead to shorter crack initiation times, which decrease due to the influence of fatigue. To predict creep-fatigue crack initiation, a modified two-criteria-diagram has been introduced and validated on 1CrMoNiV and 10CrMoWVNbN steels.25–27 For practical applications, a reduction of the crack initiation time from tic to ticf = 0.6 tic is proposed.25 This characterizes the materials resistance in the modified two-criteria-diagram by a new time-dependent parameter KIicf(ticf). Besides this proposed rule, the structure of the two-criteria-diagram for creep crack initiation remains unchanged. The validity of the boundary line is confirmed by the results of creep–fatigue crack tests. The creep crack growth rate da/dt is plotted against parameter C* in Fig. 16.17 for 1CrMoNiV steel and in Fig. 16.18 for 10CrMoWVNbN steel in cast and forged conditions. An almost linear correlation between da/dt and parameter C* on the log–log scale can be observed. Furthermore, the data are compared with the Nikbin–Smith–Webster A model. All data fall more or less close to the plane stress line. The Nikbin–Smith–Webster A plane 75 10CrMoWVNbN cast steel T = 600 °C ∆ai = 0.5 mm R = 0.1
KIi (MPa m1/2)
50
25 Specimen/a0 /W/tH
10 7.5 5 101
Cs25/0.55/0.3 h Cs25/0.55/3.0 h Cs25/0.55/∞ Cs50/0.55/0.3 h Cs50/0.55/3.0 h Ds60/0.20/0.3 h Ds60/0.20/3.0 h
Pure creep 1/t H
102
103 t i (h)
104
105
16.16 Stress intensity factor KIi versus creep-fatigue crack initiation time ti; 10CrMoWVNbN cast steel, T = 600°C.
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Creep fatigue behaviour and crack growth of steels
465
100 1CrMoNiV steel T = 550°C 0.5 mm < ∆a < 3.0 mm
–2
da/dt (mm h–1)
10
Specimen/a0 /W Cs25/0.55 Cs50/0.40–0.55 CT100/0.50
10–3
10–4
da/dt = (3·(C*/1000)0.85)/ Au* Au* = Au (NSW-A-stress) Au* = Au/30 (NSW-A-strain)
10–5
10–6 10–5
10–4
10–3
10–2 C* (N mm–1 h–1)
Specimen/a0 /W D15/0.4–0.6 D30/0.2–0.4 D60/0.1–0.4
10–1
101
16.17 Creep crack growth rate da/dt versus parameter C*; 1CrMoNiV steel, T = 550°C.
100 10CrMoWVNbN steel T = 500 and 600°C 0.5 mm < ∆a < 3.0 mm
da/dt (mm h–1)
10–2
10–3
Specimen/a0/W/steel Cs25/0.55/forged Cs25/0.55/cast
10–4
Ds60/0.20/forged Ds60/0.20/cast da/dt = (3·(C*/1000)0.85)/ Au* Au* = Au (NSW-A-stress) Au* = Au/30 (NSW-A-strain)
10–5
10–6 10–5
10–4
10–3
10–2 C* (N mm–1 h–1)
10–1
16.18 Creep crack growth rate da/dt versus parameter C*; 10CrMoWVNbN steels, T = 550°C and 600°C.
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101
466
Creep-resistant steels
strain line provides a conservative prediction for the data and thereby for components. As an example, the fatigue crack growth results for pipe steel 9CrMoVNb are presented in Fig. 16.19. The results of this constant amplitude tests show a common scatter band. The crack growth under creep–fatigue conditions can be described by depicting the crack propagation per cycle da/dN against the range of effective stress intensity ∆KIeff, which relates to the stress ratio R (Fig. 16.20). With increasing hold time there is an increase of the creep crack propagation per cycle observed. Results for creep–fatigue crack tests with tH ≤ 0.3 h are comparable to results of tests under pure fatigue conditions. As mentioned above, fatigue processes dominate the crack growth at short hold times and creep processes at long hold times. For intermediate loading conditions an accumulative crack growth is assumed, which can be determined by the summation of increments of creep crack growth and fatigue crack growth.21 The creep fatigue crack propagation per cycle is then given by the accumulation rule: da/dNcf = da/dNf + tHda/dtc
[16.8]
Beginning from the initial crack length a0, creep crack growth increments were considered when the accumulated time increments exceeded the creep crack initiation time tic. Results of such a calculation are represented in Fig. 16.21. The predicted creep–fatigue crack length ∆a cf′ is composed of the fatigue portion ∆a f′ ( t ) and the creep portion ∆a f′ ( t ) . The predicted 100
da/dt (mm h–1)
9CrMoVNb steel T = 600°C R = 0.1
Specimen/a0 /W Cs25/0.55 D15/0.20 D60/0.20
10–2
10–3
10–4
10–5 10–6 5
7.5
10
25
75
∆Kl(MPa m1/2)
16.19 Crack propagation per cycle da/dN versus range of stress intensity ∆KI; 9CrMoVNb steel, T = 600°C, range of effective stress intensity 䊐KIeff.
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Creep fatigue behaviour and crack growth of steels
467
100
10–2
tH 10–3
Specimen/tH (h)/ R Cs25/0.3/0.1 Cs25/3.0/0.1 Cs50/0.3/0.1 Cs50/3.0/0.1 DS60/0.3/0.1 Ds60/3.0/0.1 Cs25/0.3/0.6
10–4
10–5
10–6 5
7.5
10
25 ∆Kleff(MPa m1/2)
75
16.20 Creep–fatigue crack propagation da/dN per cycle versus range of effective stress intensity ∆KIeff; stress ratio R, 10CrMoWVNbN cast steel, T = 600°C.
4 1CrMoNiV steel T = 550°C Cs25-specimen σn0 = 173 MPa tH = 1.0 h/R = 0.1
3
∆acf (mm)
da/dN (mm cycle–1)
10CrMoWVNbN cast steel T = 600°C
Measured ∆acf (t) Calculated ∆acf′ (t)
∆a c′ (t ) 2
1
∆a f′ (t )
0 0
2000
4000
6000
8000
10 000
t (h)
16.21 Experimental value ∆acf and predicted value ∆a cf ′ of creep– fatigue crack length as a function of time for an individual experiment; 1CrMoNiV steel, T = 550°C.
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Creep-resistant steels
values meet the measured values in an acceptable scatter. Predicted and measured values of the creep-fatigue length of different specimen types and sizes are compared in Fig. 16.22. The crack length is slightly overestimated by the prediction. The basic crack properties of materials are responsible for the scatter.28
16.7
Concluding remarks
The multi-stage creep–fatigue behaviour of conventional and modern heatresistant steels was investigated by service-type experiments and a numerical simulation. Knowledge of cyclic deformation and creep–fatigue damage assessment has been obtained. The industrial benefit can be summarized as follows:
• •
The development of a creep–fatigue interaction concept covers physical interpretations of deformation and failure mechanisms. Stress–strain path and creep–fatigue life can be predicted by user program SARA on the basis of rules for deformation, relaxation and cyclic stress– strain behaviour including internal stress and mean stress. A creep–fatigue life estimation procedure was developed for multi-stage loading, which covers cycle counting methods. The procedure was 100 1 CrMoNiV steel T = 550°C R = 0.1
∆a cf′ (mm)
•
1
tH(h) Cs25 Cs50 0.3 1.0 3.0 10.0
0.1
0.01 0.01
0.1
1 ∆acf (mm)
D60
100
16.22 Comparison between the predicted value ∆a cf ′ and experimental value ∆acf of creep–fatigue crack length; 1CrMoNiV steel, T = 550°C.
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Creep fatigue behaviour and crack growth of steels
•
• • •
469
established in power plant applications. An extension to automotive and aero engine applications is of future interest. For verification purposes long-term service-type creep–fatigue data up to 70 000 h were generated. In order to characterize creep and creep fatigue crack behaviour, different methods exist, like the two-criteria-diagram for creep crack initiation. The methods consider the linear elastic-parameter K1 as well as the creep fracture mechanics parameter C*. These methods are validated in long term regime on numerous materials, mainly on type 1CrMoNiV steels and in recent years on type 10CrMoWVNbN steels. Long-term experiments, both uniaxial as well as multiaxial, are necessary in order to verify advanced life prediction methods for components. Modelling creep–fatigue behaviour in multiaxial and multi-stage loading with constitutive models is a current challenge. Advanced lifing methods and knowledge of materials as well as methodologies enables a reduction design efforts, an increase in component loading and an increase in design quality linked with a reduction in technical risk and an increase in efficiency and economic benefits.
16.8
Acknowledgements
Thanks are due to the Forschungsvereinigung der Arbeitsgemeinschaften der Eisen und Metall verarbeitenden Industrie e.V. (AVIF), Project No. A166, the FKM Forschungskuratorium Maschinenbau e.V., Project No. 052510, the Arbeitsgemeinschaft industrieller Forschungsvereinigungen (AiF), the VDEh-Gesellschaft zur Förderung der Eisenforschung mbH, Project No. 11200 N and to the Deutsche Forschungsgemeinschaft (DFG), Projects No. BE1890,13-1 and BE1890, 16-1 for financial support, and the working groups of power plant industry for their interest and technical support.
16.9
References
1 J. Granacher, A. Scholz and C. Berger, ‘Creep fatigue behaviour of heat resistant turbine rotor steels under service-type strain cycling’, Proceedings of the Fourth International Charles Parsons Turbine Conference, Newcastle upon Tyne, A. Strang, W. M. Banks, R. D. Conroy and M. J. Goulette (eds), The Institute of Materials, London, 1997, 592–602. 2 D. A. Miller and R. H. Priest, Materials Response to Thermal–Mechanical Strain Cycling, High Temperature Fatigue: Properties and Prediction, R. P. Skelton (ed.), Elsevier Applied Science, London and New York, 1987, 113–176. 3 A. Scholz, H. Haase and C. Berger, Simulation of Multi-Stage Creep–Fatigue Behaviour, Fatigue 2002, Proceedings of the Eighth International Fatigue Congress, A. F. Blom (ed.), Volume 5/5, Stockholm, June 2002, 3133–3140. 4 K. N. Smith, P. Watson and T. H. Topper, ‘A stress–strain function for the fatigue of metals’, J. Mater., 1970, 4, 767–778.
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5 E. L. Robinson, ‘Effect of temperature variation on the creep strength of steels’, Trans ASME, 1938, 60, 253–259. 6 S. Taira, ‘Lifetime of structures subjected to varying load and temperature’, in Creep in Structures, Hoff, N. J. (ed.), Academic Press, New York, 1962, 96–124. 7 J. S. Dubey, H. Chilukuru, J. K. Chakravartty, M. Schwienheer, A. Scholz and W. Blum, ‘Effects of cyclic deformation on subgrain evolution and creep in 9–12% Cr Steels, Mater. Sci. Engng. A, 406 (2005), S. 152–159. 8 K. Yagi, O. Kanemaru, K. Kubo and C. Tanaka, Life prediction of 316 stainless steel under creep-fatigue loading, Fatigue Fracture Engng. Mater. & Struct., 1987, 9 (6), 395–408. 9 A. Scholz and C. Berger, ‘Deformation and life assessment of high temperature materials under creep fatigue loading’, First Symposium on Structural Durability in Darmstadt, June 9–10, 2005, Orangerie Darmstadt, Proceedings, C. M. Sonsino (ed.), Fraunhofer IRB Verlag, Stuttgart, 2005, S. 311–328. 10 J. Granacher, A. Scholz, H. Möhlig and C. Berger, ‘Heat resistant power plant steels under variable long-term conditions’, 5th International Charles Parsons Turbine Conference, Cambridge, Strang, A., Banks, W.M., Conroy, R. D., McColvin, G. M., Neal, J. C. and Simpson, S. (eds), IOM Communications, London, 2000. 11 B. J. Cane, Life Management of Ageing Steam Turbine Assets, Proceedings, Institute of Materials, London, 1997, 554–574. 12 R. Znajda, Betriebsähnliches Langzeitdehnwechselverhalten wichtiger Stahlsorten, Doctoral Thesis, D17, Darmstadt University of Technology, Shaker Verlag, Aachen, 2007. 13 M. M. Itoh, M. Sakane and M. Ohnami, High temperature multiaxial low cycle fatigue of cruciform specimen, Trans. ASME, JEMT, 116, (1), 1994, 90–98. 14 C. Tsakmakis and D. Reckwerth, ‘The principle of generalized energy equivalence in continuum damage mechanics’, Deformation and Failure of Metallic Continua, Hutter K., Kirchner N. and Baaser H. (eds), Springer Series – Lecture Notes in Mechanics, 2003, 381–406. 15 J. Lemaitre and J.-L. Chaboche, Mechanics of Solid Materials, Cambridge University Press, Cambridge, 1990. 16 ASTME 1457-00, Standard Test Method for Measurement of Creep Crack Growth Rates in Metals, Annual Book of ASTM Standards, 2001, 3 (1), 936–950. 17 C. M. Davis, F. Mueller, K. M. Nikbin, K. M. O’Dowd and G. A. Webster, Analysis of creep crack initiation and growth in different geometries for 316H and carbon manganese steels, J. ASTM Inte., 3, (2), 2006. Paper JAI 13220 available on-line at www.astm.org. 18 G. R. Irwin, ‘Analysis of stresses and strains near the end of a crack traversing a plate’, Trans. ASME, J. Appl. Mech., 1957, 24, 361–364. 19 J. Ewald and K. H. Keienburg, A two-criteria-diagram for creep crack initiation, International Conference on Creep, JSME, IMechE, ASME, ASTM, Tokyo, 1986, 173–78. 20 K. M. Nikbin, D. J. Smith and G. A. Webster, ‘Influence of creep ductility and state of stress on creep crack growth’, Advances in Life Prediction Methods at Elevated Temperatures, Woodford, D. A. and Whitehead, J. R. (eds), ASME, 1983, 249–258. 21 G. A. Webster and R. A. Ainsworth, High Temperature Component Life Assessment, 1st edition, Chapman and Hall, London, 1994.
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22 K. M. Nikbin, D. J. Smith and G. A. Webster, An Engineering Approach to the Prediction of Creep-Crack Growth, Journal of Engineering Materials and Technology, 1986, 186–191. 23 M. Tan, N. J. C. Célard, K. M. Nikbin and G. A. Webster, Comparison of creep crack initiation and growth in four steels tested in HIDA, Int. J. Pressure Vessels and Piping, 2001, 78(12), 737–747. 24 F. R. Larson and J. A. Miller, Time–temperature relationship for rupture and creep stresses’, Trans. ASME, 1952, 74, 765–775. 25 J. Ewald, S. Sheng, A. Klenk and G. Schellenberg, Engineering guide to assessment of creep crack initiation on components by Two-Criteria-Diagram, Int. J. Pressure Vessels and Piping, 2001, 78, (11–12), 937–949. 26 J. Granacher, A. Klenk, M. Tramer, G. Schellenberg, F. Mueller, and J. Ewald, Creep-fatigue crack behaviour of two power plant steels, Int. J. Pressure Vessels and Piping, 2001, 78, (11–12), 909–920. 27. F. Mueller, A. Scholz, C. Berger, A. Klenk, K. Maile and E. Roos, ‘Crack behaviour of 10Cr-steels under creep and creep–fatigue conditions’, ECCC Creep Conference, 12–14 September 2005, London. 28 B. Dogan, U. Ceyhan, J. Korous, F. Mueller and R. A. Ainsworth, ‘Sources of scatter in creep/fatigue crack growth testing and their impact of plant assessment’, Proceedings of FITNET 2006, International Conference on Fitness-for-Service, 17–19 May 2006, Amsterdam. 29 J. A. Nelder and R. Mead, ‘A simplex method for function minimization’, Comput. J., 1965, 7, 308–313.
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17 Creep strength of welded joints of ferritic steels H . C E R J A K and P . M A Y R, Graz University of Technology, Austria
17.1
Introduction
Ferritic steel grades are highly valued for fabrication of components in the thermal power generation industry. As well as low chromium ferritic/bainitic steels for boiler components like waterwalls or for sub-critical steam piping and headers, steels with a chromium content in the range of 9–12% for ultrasuper critical (USC) power plants are of great interest. Welding in all its varieties is still the major joining and repair technology for power plant components. Either by repair welds in casting defects, fabrication welds, joints of similar and dissimilar steel grades, connections with small crosssections for example tube to tube welds, or with large cross-sections for example pipe to pipe or pipe to casting welds, the microstructure of the joined materials is strongly influenced by the welding process and thereby their mechanical properties are altered. As for all components exposed to high temperatures during service, the 100 000 h creep rupture strength of base material (BM), weld metal (WM) and cross-welds is still the major design criterion. Table 17.1 gives an overview of the chemical composition of commonly used ferritic creep-resistant steels and their 100 000 h creep rupture strength. Long-term experience of creep exposed welded structures has shown that the heat affected zone (HAZ), a narrow zone of base material adjacent to the weld fusion line altered by the weld thermal cycle, in respect to the creep strength, is often regarded as the weakest link in welded constructions. Considering just plain steam pressure (p) loading in a pipe, the HAZ can be stressed in series and in parallel.1 In the absence of additional mechanical or thermal loading, the hoop stress (σH) is about twice as the longitudinal stress (σL) and series loading of strong and weak zones (WM-HAZ-BM) is more significant than parallel loading. Longitudinal seam welds show the highest risk of failure (Fig. 17.1). If additional loading arises, that is insufficient dead weight support or sagging, the total axial stress (σL) can be greater than the hoop stress (σH) and the HAZ in girth welds exhibits at least similar risk of damage as in seam welds. 472 WPNL2204
Table 17.1 Chemical composition (wt%) and creep rupture strength of widely used creep-resistant ferritic steels for application in thermal power plants Low Cr steels
C
Si
Mn
Cr
Mo
Cu
W
Ni
V
Nb
B
N
Ti
T W1 (°C)
R m2 ×10 5 /T
W
(MPa3 ) min max min max min max min max
0.08 0.18 0.08 0.14 0.04 0.10 0.05 0.10
9% Cr steels X11CrMo9-1 (T9) X10CrMoVNb9-1 (T/P91) X11CrMoWV Nb9-1-1 (E911) X10CrWMoVNb9-2 (T/P92)
min max min max min max min max
0.08 0.15 0.08 0.12 0.09 0.13 0.07 0.13
min max min max
0.17 0.23 0.07 0.14
12% Cr steels X20CrMoV12-1 T/P122 HCM12A
0.40 1.00 0.40 0.80 0.10 0.60 0.30 0.70
0.70 1.15 2.00 2.50 1.90 2.60 2.20 2.60
0.40 0.60 0.90 1.10 0.05 0.30 0.90 1.10
0.50
0.30 0.60 0.30 0.60 0.30 0.60 0.30 0.60
8.00 10.00 8.00 9.50 8.50 9.50 8.50 9.50
0.90 1.10 0.85 1.05 0.90 1.10 0.30 0.60
0.50
1.00
0.50
0.70
10.00 12.50 10.00 12.50
0.80 1.20 0.25 0.60
0.35 0.50 0.50 0.15 0.45 0.25 1.00 0.50 0.10 0.50
0.30
0.012
0.30
0.012 1.45 1.75
0.20 0.30 0.20 0.30
0.02 0.08
0.0005 0.0060 0.0015 0.0070
550
49
550
68
550
130
550°C
147
0.030 0.010
0.05 0.10
550
92
600
94
600
98
600
113
600
49
600
103
0.30 0.30 0.90 1.10 1.50 2.00
0.30 0.10 0.40 0.40 0.30 0.80
0.30 1.70
1.50 2.50
0.50
0.18 0.25 0.18 0.25 0.15 0.25 0.25 0.35 0.15 0.30
0.06 0.10 0.06 0.10 0.04 0.09
0.04 0.10
0.0005 0.0050 0.0010 0.0060
0.0005 0.0050
0.030 0.070 0.050 0.090 0.030 0.070
0.040 0.100
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1service temperature; 2 100 000 h creep rupture strength at service temperature; 3 Software Stahlschlüssel 2004, Verlag Stahlschlüssel Wegst GmbH. Version 4.00.0005; www.stahlschluessel.de; 4 The T23/T24 Book – New Grades for Waterwalls and Superheaters, Vallourec & Mannesmann Tubes, 2000; 5 ECCC Data Sheets, 2005, compiled and published by D G Robertson and S R Holdsworth, ETD Ltd; 6 METI (Ministry of Economy, Trade and Industry) Thermal Power Standard Code, Japan (single phase with no delta ferrite).
Creep strength of welded joints of ferritic steels
13CrMo4-4 (T/P12) 10CrMo9-10 (T/P22) 7CrWVMoNb9-6 T/P23 7CrMoVTiB1010 (T/P24)
474
Creep-resistant steels σH =
p×D 2×t
σL =
p×D 4×t
σH : σL = 2 : 1
p
σH σL
17.1 Schematic of different loading conditions acting on the HAZ of weldments in a pressurised pipe with mean diameter D and wall thickness t.1
Awareness of the importance of knowledge of creep behaviour in welded structures has increased continuously over the last few decades. In the 1990s several failures of welded steam piping systems, some of them in a catastrophic manner (see Fig. 17.2), have backed the need for investigation into the creep behaviour of welded structures.2 At that time, these failures were not seen as anything else than an anomaly resulting from improper fabrication or installation or inappropriate service conditions. With increasing hours in service, various other cracks, leaks and even pipe ruptures occurred and the industry realised the symptomatic manner of some problems caused by the utilisation of welded ferritic steel components. These failures acted as a driving force for increased research activities on failure characterisation, non-destructive testing methods, remaining life prediction methods, repair technology as well as the development of new improved steel grades and welding procedures. The reason that many power stations are operating beyond their design life has again boosted the research efforts of manufacturers, operators as well as academics. Nowadays, an almost similar effort is put into the determination of the creep properties of weld metals and weldments as of base materials. In the following, the state of the art of science and technology of creep-exposed ferritic weldments covering topics like influence of welding procedures on the microstructure, HAZ simulation, weld metal development, creep behaviour of welded joints, selected damage mechanism in creepexposed welded structures, implications for the industry using ferritic creepresistant steels and finally future trends in this field will be discussed.
17.2
Influence of weld thermal cycles on the microstructure of ferritic heat-resistant steels
Fusion welding, as the most important joining process in power plant construction works, strongly affects the joint properties.3 Not only is a new
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Mohave
17.2 Catastrophic failure of seam welded steam pipe in the USA in 1985 causing six deaths and 10 injuries. (Picture courtesy of C. D. Lundin, University of Tennessee).2
type of material, the weld metal, deposited between the connected parts but also the base material is altered by local and very inhomogeneous heat treatment as a result of the weld thermal cycle. Good overviews on welding metallurgy are given by Easterling,4 Granjon5 and Schulze.6 This chapter focuses on phase transformations taking place in the base material of ferritic heat-resistant steels within the HAZ, strongly influencing the resultant mechanical properties of the weldment. The optimised base materials microstructure and properties, set through accurate melting techniques, an exact production process control and proper heat treatment by the base metal manufacturer, are changed completely within
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the HAZ by the applied weld thermal cycle. In addition, internal stresses generated during rapid cooling, which are in the order of the yield strength of the material, sum up with system stresses later on in service. In Fig. 17.3, the basic influences of the welding process on the metallurgy in the HAZ are shown and compared to the calculated equilibrium phase diagram of X10CrMoVNb9-1 (P91) steel. Depending on the selected welding process, the base material microstructure within the HAZ will be changed. The resulting microstructure is governed by the heating rate of the weld thermal cycle, the experienced peak temperature (TP), the dwell time, the cooling rate, the effects of multilayer welding and finally by the adjusted postweld heat treatment (PWHT) parameters. The heating rate in arc welding processes can be as high as 200–300 K s–1. As a result, transformation temperatures are shifted to significantly higher temperatures than predicted in the equilibrium phase diagram (T0). For example, ferrite (α) to austenite (γ) phase transformation can occur about 100 K higher at a heating rate of 100 K s–1 than the calculated T0, resulting in considerable superheating of the ferrite before transformation. Other parameters affected by the heating rate are: recrystallisation temperature, coarsening rate of carbides and nitrides, solution temperature of carbides and nitrides and the main proportion of grain growth. Precipitation strengthening is one of the most effective mechanisms active in ferritic creep-resistant steels. Therefore, precipitate stability is a key factor both in the base material and also in the HAZ of weldments. At higher heating rates, not only phase transformations but also solution temperatures of precipitates are shifted to higher temperatures. In most cases equilibrium is not reached at high heating rates and short dwell times and superheating occurs. It is of great importance to estimate how this may affect the microstructural evolution, especially the grain growth behaviour, which is strongly influenced by grain boundary pinning. When welding precipitation strengthened creep-resistant steels, three scenarios should be considered: • • •
Peak temperature is too low to have any noticeable effect on the precipitates; Particles only partially dissolve during weld thermal cycle, but coarsening of favoured particles occurs; Particles dissolve completely during the thermal cycle and omission of grain boundary pinning causes excessive grain growth.
Besides precipitation strengthening, grain size is of major importance, being a key factor for good mechanical properties, such as tensile strength, toughness, creep rupture strength or with respect to determining the susceptibility of the alloy to several damage mechanisms like cold cracking, reheat cracking or type IV cracking. In regions where most of the precipitates have been dissolved by the weld
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Temperature (°C)
Solid–liquid transition zone Coarse prior austenite grains + fine prior delta ferrite grains
δ
1400
L+δ
L+γ+δ
L L+γ
δ+γ
CGHAZ 1200
Grain growth zone
γ 1000
Grain refined zone
FGHAZ Intercritical zone Over-tempered region
800
α+γ
Unaffected BM α 600
Heat affected zone
0
0.2
0.4 0.6 Carbon (wt %)
0.8
17.3 Schematic of the sub-zones of the HAZ corresponding to the calculated equilibrium phase diagram of X10CrMoVNb9-1 (P91-type) steel.1
1
Creep strength of welded joints of ferritic steels
Peak temperature Tp
Solidified weld
477
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thermal cycle, excessive grain growth may take place. After completion of α/γ transformation, the stability of newly formed grains is far from equilibrium. With increasing peak temperature, crystallographic favoured grains begin to grow at the expense of smaller grains. The resulting coarse-grained microstructure generally shows low toughness characteristics and in the case of some low alloyed ferritic/bainitic Cr-steels significant susceptibility to reheat cracking. In 9–12% Cr steels, according to the equilibrium phase diagram, austenite starts to transform to delta ferrite (δ) at the highest peak temperatures. The nucleation of delta ferrite grains at austenite grain boundaries again causes the overall grain size to decrease. While the lower solubility of carbon and other austenite stabilising elements in ferrite results in the diffusion of these elements out of the delta ferrite to the remaining austenite regions, ferrite formers, like chromium, are enriched in the ferritic regions. Therefore, segregated regions differ locally in chemical composition and austenitic transformation on cooling may be incomplete, resulting in retained delta ferrite. Second, high cooling rates, predominant for low heat input welding processes, like electron beam- (EB), laser- and gas tungsten arc welding (GTAW), and thick walled components can also result in an incomplete reverse-transformation of delta ferrite back to austenite. Small amounts of delta ferrite may be present in the microstructure even at room temperature. Residual amounts of delta ferrite in a martensitic matrix have been shown to have a negative influence on impact values as well as on creep rupture strength and therefore are undesirable in these types of steels.7 Multilayer welding techniques offer more time for diffusion to compensate segregational processes and result in a complete and homogeneous retransformation of delta ferrite to austenite.8 In the following paragraphs the HAZ of creep-resistant chromium steels and their sub-regions are described in more detail.
17.2.1 Heat affected zone (HAZ) As pointed out above, the welding process strongly influences the microstructure and properties of the base material. As a result of the severe thermal cycle caused by the welding process, the original microstructure is altered and a so-called heat affected zone (HAZ) is formed (see also Fig. 17.3). The HAZ can be divided into a number of sub-zones. No distinct borderline between the different regions is recognisable; it is more a continuous gradient from the fusion line between the deposited weld metal to the unaffected base material. Each sub-zone is represented by its characteristic microstructure and properties.
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479
Grain growth zone (Tp >> Ac3) This zone adjacent to the fusion line experiences temperatures well above the Ac3 transformation temperature. Any precipitates that obstruct growth of austenite grains at lower temperatures dissolve, resulting in coarse grains of austenite. In 9–12%Cr steels delta ferrite grains may nucleate at the highest peak temperatures (TP > 1250°C) causing overall grain size to decrease. On cooling, lower chromium steels form a bainitic/martensitic microstructure and 9–12%Cr steels form a martensitic microstructure. The coarse-grained zone (CGHAZ) features the highest hardness of the HAZ and generally low toughness values are expected. The coarse-grained zone may be vulnerable to reheat cracking during creep loaded service. Grain refined zone (Tp > Ac3) Lower peak temperatures of about 1100°C, just above Ac3, result in an improper development of austenite, following the α/γ transformation during heating, producing small austenitic grains (FGHAZ). In addition, peak temperature may not be high enough to dissolve precipitates completely, limiting the grain growth by pinning the austenite grain boundaries. On cooling, either a fine grained bainitic microstructure for lower chromium steels or a martensitic microstructure for higher chromium steels is formed. The fine grained region of the HAZ is regarded as the weakest link in weldments during creep loaded service. At longer service times and lower stress levels most of weldments of creep-resistant ferritic steels fail within this region by the so-called type IV mechanism. Partially transformed zone – intercritical HAZ (Ac1 < Tp < Ac3) Peak temperatures lying between Ac1 and Ac3 transformation temperatures result in a partial transformation of α into γ on heating. While new austenite grains nucleate at favoured positions, like prior austenite grain boundaries or martensite lath boundaries, the untransformed bainitic or tempered martensitic microstructure is simply tempered for a second time by this weld thermal cycle. Partial dissolution of precipitates may be experienced in this part of the HAZ and coarsening of undissolved precipitates can occur especially during subsequent PWHT. After cooling, a twofold microstructure consisting of newly formed bainite, for low chromium steels, and virgin martensite, for higher chromium steels, and the tempered original microstructure coexist. The intercritical HAZ shows a small grain size and exhibits the lowest hardness values in weldments. This sub-zone of the HAZ shows similar susceptibility to type IV cracking as the grain refined zone.
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Over-tempered region With peak temperatures experienced below Ac1 the microstructure does not undergo any phase transformations but the original microstructure is locally tempered at higher temperatures compared to that of the annealed base material. As a result, coarsening of precipitates may be enhanced by a higher coefficient of diffusion at this temperature. Some alloys show the lowest hardness values in this region. Zone of unchanged base material The zone of unchanged base materials concerns temperatures up to ca. 700°C, in which changes in morphology of constituents do not appear to occur. Nevertheless, in this region over-tempering effects can occur which may weaken the creep strength of welded low alloyed quenched and tempered (Q+T) steels, that is 1%CrMoV steels.9
17.2.2 Multilayer welding Special attention has to be paid to the advantage of multilayer welding techniques. Multilayer welds are formed by subsequent deposition of weld beads on solidified formerly deposited runs. Significant changes in the microstructure of multilayer welds compared to single layer welds are the result of additional heat input into the weld metal but also the HAZ of the base material. Reed and Bhadeshia10 as well as Cerjak et al.11 characterised the HAZ of multilayer welds by mathematical modelling. This ‘multilayer’ effect is not only representative of the HAZ of base material and regions with different microstructure can be formed within the weld metal. In Fig. 17.4, intercritically heated areas within a cross-weld of E911 base material welded with matching filler have been marked in black, while the fusion lines are marked by grey lines.12 While a continuous region of intercritically heated material forms within the HAZ of the base material, the HAZ within the weld metal is discontinuously related to the loading direction. If subsequent weld beads are symmetrically deposited to the weld centreline, a continuous layer of intercritically heated material is also formed within the weld deposit (see Fig. 17.4 right).13 These heat affected regions within the weld metal can exhibit the same susceptibility to certain cracking mechanisms as is prominent in the HAZ of the base material. Therefore, in multilayer welds special attention has to be paid to which welding parameters are utilised, especially heat input, joint geometry and weld layer structure.
17.2.3 HAZ simulation The problem of conducting basic investigations of HAZ structures in actual welds is the presence of extremely small and inhomogeneous sub-zones (see
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481
(a)
(b)
17.4 (a) Multilayer weld of E911 base material with matching filler material. Areas exposed to peak temperatures between A1 and A3 transformation temperatures (intercritical HAZ) are marked in black. The loading direction is indicated by arrows.12 (b) creep exposed 21/4Cr–1Mo SAW weld with symmetrically deposited weld beads showing cracking in the refined weld metal at the weld centreline (Picture courtesy of C. D. Lundin, University of Tennessee).13
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Fig. 17.3). For the HAZ simulation the time–temperature profile acting as input data can either be measured from real welds or be derived by analytical solutions of the heat conduction equation or by using more sophisticated numerical thermal heat source models.14–16 HAZ simulation allows the generation of larger volumes of material with uniform microstructure and properties which represent one specific point within the HAZ. This generated homogeneous microstructure can be used for all kinds of metallographic investigations and tested by applying common standardised mechanical testing procedures such as tensile tests, creep and fatigue tests and impact tests.17 Different techniques are currently applied for the HAZ simulation, namely: • • • •
heating in a hot salt bath until the peak temperature is reached, immediately followed by cooling in a moderated tempered salt bath; induction heating with subsequent quenching in an oil bath; heating in a furnace to peak temperature, followed by cooling in air or an oil bath; controlled resistance heating in the specimen using a weld simulator (GLEEBLETM, SmitweldTM)
Each physical simulation process has its own characteristic advantages and disadvantages. A good overview is given by Buchmayr.18
17.3
Weld metal development for creep-resistant steels
Within the last few decades great efforts were put worldwide into the development of filler materials for newly developed creep-resistant steel grades. It is of great importance that matching filler metals are developed simultaneously to base materials.19 The term ‘matching’ is not directly related to the chemical composition but more specific to the design-based tensile properties, in this case to the creep rupture properties of the weld metal, compared to that of the base metal although it is known that similar chemical composition results in similar creep rupture properties.20 This is correct for a great number of filler metals, but special deoxidation practices, additions of micro alloying elements, improvement of weldability and handling remain the detailed knowledge of the individual filler material manufacturer. Baune et al. summarised the objectives for the development of new weld metals for 2.25%Cr steels (T/P23, T/P24) as follows:21 • •
adjust alloying additions to the electrodes so that mechanical properties of welds are properly balanced, i.e. impact toughness, tensile properties, hardness and creep rupture strength; minimise susceptibility to temper embrittlement. Therefore, the amount of residual elements such as P, Sb, Sn, and As should be minimised through careful control of raw materials;
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• • • •
483
maximise toughness in the as-welded condition, whereas impact toughness values higher than 27 J after PWHT are often required by fabricators; set tensile requirements of the weld metal at room temperature according to the specifications given for the base material; ensure hardness values of the weld metal, obtained in as-welded condition are below 350HV10 and below 248HV10 after PWHT; evaluate, finally, the influence of different PWHT conditions on the properties of weldments.
Matching weld metal for creep-resistant chromium steels solidifies primary ferritic. During solidification, delta ferrite is enriched by chromium and molybdenum and the residual melt is enriched by nickel. The chemical composition of segregated regions is strongly influenced by diffusion and changes continuously during cooling. The rapid cooling process inhibits a complete δ/γ-phase transformation and residual δ-ferrite is often present after cooling to room temperature. The detrimental effects of δ-ferrite on impact toughness and creep rupture properties of ferritic steels are well known.7 Therefore, weld metal development especially for 9–12% chromium steels, aims to eliminate retained δ-ferrite by modification of the weld metal chemistry. Additions of Ni, as an austenite-stabilising element, showed beneficial effects on the impact toughness properties of NF616 (Grade 92) weld metal by limiting the retention of δ-ferrite.22 While nickel addition lowers the A1 transformation temperature significantly, the addition of cobalt has almost no influence on the transformation temperatures but can also reduce retained δ-ferrite content effectively.12 In the new generation of ferritic chromium steels, tungsten has been added to improve creep rupture strength. However, tungsten is a strong ferrite stabilising element and promotes the retention of delta ferrite in the weld metal. Additions of 1%Ni or 2%Co in combination with 1%W have been shown to be sufficient in eliminating retained δ-ferrite and improving impact toughness in a modified 9Cr–1Mo steel.23
17.4
Creep behaviour of welded joints
Weldments of ferritic steels exposed to high temperature creep loading show very similar behaviour in tendency, irrespective of the chemical composition and other parameters like welding procedure, groove preparation and so on. Numerous creep tests of cross-weld samples showed that at lower temperatures there is no big difference between the base metal and the cross-weld creep strength.24 This difference becomes more prominent as the temperature increases and as the applied stress level is lowered. The time of deviation of cross-weld creep strength from base metal mean creep strength varies with creep testing conditions, material grades and welding prerequisites, for example stress level, temperature, welding procedure, PWHT parameters, and so on.
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Figure 17.5 compares the creep rupture strength of cross-welds with the mean line of base metal creep rupture strength of 9% Cr steel grade E911. At higher stress levels and lower testing temperatures the location of failure is rather randomly distributed between BM, WM and HAZ. At lower stress levels and higher testing temperatures the HAZ of creep-resistant steels appears to be the weakest link, diminishing the creep strength of weldments by up to 50%. Generally, as derived from numerous long running creep tests, the creep rupture time is longer in the order of weld metal, base metal, crosswelds and fine grained simulated HAZ material.25
17.5
Selected damage mechanism in creep-exposed welded joints
In 1974, Schüller et al.26 categorised the types of cracking observed in weldments of heat-resistant steels by a simple scheme (see Fig. 17.6). Cracks were classified depending on their location and orientation within the weldments. Cracks in the deposited weld metal correspond to type I and type II. They form in the weld metal and develop either in the longitudinal or transversal direction but rarely may also be completely unoriented. While type I cracks are arrested in the weld metal, type II cracks may propagate into the HAZ or even into the base material. The other types of cracks develop within the HAZ of the weldments. Type III cracks form in the coarse-grained part of the HAZ close to the fusion line and can prolong in 500
Stress (MPa)
300
600°C
100 80
625°C
60 40
650°C
E911 base material Crosswelds 600°C Crosswelds 625°C Crosswelds 650°C 10 101
102
103 Time to rupture (h)
104
105
17.5 Comparison of cross-weld creep rupture strength with mean base material creep rupture strength of E911 steel at different temperatures.12
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Creep strength of welded joints of ferritic steels II
III
485
IIIa
I III IV IV
HAZ
I IIIa II
HAZ
17.6 Modified schematic of cracking modes in weldments of heatresistant steels (original by Schüller et al.).26
this zone as well as into the base material, whereas cracks of type IV are restricted to the fine or intercritical zone adjacent to the unaffected base material. Brett27 added a term for type IIIa cracking to this scheme for a failure mechanism taking place close to the fusion line (type III) but in a fully refined HAZ structure with higher fracture ductility.
17.5.1 Type I, type II cracking Defects in the weld metal are commonly associated with the welding process itself or stress relief during PWHT; more rarely weld metal cracks can be related to creep damage during high temperature service.28,29 Cracks in the weld metal are mostly in the transversal direction although longitudinal failures are also reported. Previous reported cracks were almost intercrystalline and appeared as hot cracks formed during solidification of the weld metal. Since weld metal development has improved and the cleanliness of the weld deposit has increased during the last few decades, the significance of solidification cracks in ferritic steels has diminished. They are still of great concern in austenitic and nickel base weld metals. As described above, the weld metal deposit produced by a multilayer welding technique followed by sub-critical tempering does not exhibit a uniform microstructure. In multilayer welding, the solidified weld beads underneath experience a significant thermal influence from subsequent weld thermal cycles. Therefore, similar zones, compared to those of the HAZ of base metal, develop in a multilayer weld metal. These zones within the weld metal can be definitely susceptible to different forms of creep damage. While coarse-grained areas in the weld deposit might show a susceptibility to reheat cracking, fine grained regions might fail by type IV cracking with the same mechanisms active as in the analogous part of the HAZ of the base metal. Both failure mechanisms will be described in more detail in the following paragraphs.
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17.5.2 Type III cracking – reheat cracking Reheat or stress relief cracking is defined as intergranular cracking in the HAZ or the weld metal that occurs during the exposure of welded structures to elevated temperatures produced by PWHT or high temperature service up to about 20 000 h. This type of cracking in welded structures of creepresistant, precipitation-strengthened alloys has received considerable attention since the mid-1950s.30 Initial work was concerned primarily with cracking problems in austenitic steels (18Cr–12Ni–1Nb) used for power generating equipment, especially high temperature steam piping. Reheat cracking was also reported in some Ni base superalloys.31 In the early 1960s similar cracking problems during stress relief heat treatments and high temperature service were again observed in power generating constructions where the ferritic creep resisting 2CrMo and CrMoV weldments in steam pipework and valve assemblies were found to exhibit occasional cracking. While high alloyed 9–12% chromium steels show almost no susceptibility to reheat cracking, newly developed low alloyed 21/4% chromium steels can exhibit significant susceptibility depending on their alloying concept. The mechanism of reheat cracking is generally understood although details of the controlling parameters and their mechanistic interpretation still remain a subject for discussion. In general terms, cracking results when the relaxation strains, occurring with stress relief of residual stresses during PWHT or high temperature service, exceed the local ductility of the material. Factors affecting the susceptibility of welded structures to reheat cracking include the chemical composition, the microstructure resulting from the applied welding process and the stress state. The significance of segregation of alloying elements like Al, B, Mn and impurity elements like P, S, As, Sb, Sn, and so on is under discussion in literature. Some authors report a strong influence of elements segregated at the prior austenite grain boundaries (PAGB) in lowering the cohesive strength of the boundaries whilst others find no interaction between segregation and stress relief cracking failures.30–32 Generally, it is evident that all mechanisms weakening or embrittling grain boundaries enhance the susceptibility to reheat cracking. In the CGHAZ almost all precipitates are dissolved during the applied weld thermal cycle. During PWHT and high temperature service reprecipitation takes place at the grain boundaries as well as inside the grains. The grain interior gets strengthened by precipitation of finely dispersed, mainly coherent carbonitrides while PAGB are energetically favourable for precipitation of incoherent carbides like Fe3C, M23C6, M6C. Highly diffusive grain boundaries enhance the coarsening process of the latter and result in the depletion of solid solution strengthening elements and the dissolution of MX particles in the vicinity of PAGB.
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Creep strength of welded joints of ferritic steels
487
As a result of these microstructural changes, precipitation-strengthened grain interiors are surrounded by weak grain boundary areas. Hence, reduction of welding residual stresses mainly proceeds by deformation concentrated at the grain boundaries. Therefore, stress relief cracking appears to be macroscopic as intergranular cracks along PAGB with low rupture ductility. Examination has shown that damage caused by stress relief cracking forms at the boundaries of large prior austenite grains by a mechanism of creep cavitation.33 Cavities nucleate primarily on PAGB as incoherent precipitates, acting as stress concentrators. Susceptibility to reheat cracking is a function of precipitation strengthening in the grain interior, strengthening and weakening the behaviour of grain boundaries and the relaxation of residual stresses. In recent research work, the susceptibility of up-to-date 21/4% chromium steels for implementation as membrane waterwalls for USC boilers or superheater and reheater components of conventional boilers and heat recovery steam generators was investigated.33–36 The Japanese ferritic/bainitic steel HCM2S (7CrWVMoNb9-6, ASTM A213 T/P23) and the German 7CrMoVTiB10-10 (ASTM A335 T/P24) showed different susceptibilities to reheat cracking. Isothermal slow strain rate tensile tests on weld simulated specimens were used for qualification. After fracture, the reduction of area is measured as an indicator of ductility. The reduction in area is directly related to the cracking susceptibility and classified as shown in Table 17.2. The results, summarised in Fig. 17.7, show that care must be taken when welding T23 and P23. The pipe material P23 is highly susceptible in a broad PWHT temperature range from 600°C to 750°C while tube material shows only slight susceptibility to reheat cracking at PWHT temperatures higher than 675°C. In contrast neither T24 nor P24 material shows any susceptibility to reheat cracking. Figure 17.8 shows characteristic fracture surfaces of P23 and P24 pipe material after isothermal slow strain rate testing in an inert atmosphere. While the fracture surface of CGHAZ simulated P23 specimen can be categorised as almost 100% intergranular brittle fracture, P24 pipe material shows a tendency for intergranular fracture but with a predominant ductile fracture surface. Multi-pass welding, producing a fine grained microstructure Table 17.2 Classification of reheat cracking susceptibility by reduction in area measurements on a CGHAZ simulated specimen after isothermal slow strain rate tensile testing30 Susceptibility to reheat cracking
% reduction in area
Extremely susceptible Highly susceptible Slightly susceptible Not susceptible
<5 5–10 10–20 > 20
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100 90
P22 T23
P23 T24
P24
Reduction of area Z (%)
80 70 60 50 40 30 20 10 0 525
550
575
600
625 650 675 700 Testing temperature (°C)
725
750
775
17.7 Reduction in area as a function of testing temperature for isothermal slow strain rate tensile tests for different CGHAZ simulated 21/4Cr steels. P23 pipe material shows the highest susceptibility to reheat cracking according the classification in Table 17.2.
by the normalising effect on prior deposited layers, resulted in an increase in rupture ductility and eliminated reheat cracking susceptibility.
17.5.3 Cracking in dissimilar welds For the sake of completeness, cracking in dissimilar welds is mentioned in this chapter. There are a multitude of different types of dissimilar welds depending on the steel grades and welding consumables used. In thermal power stations, dissimilar welded joints between ferritic and austenitic steels and high alloyed martensitic and low alloyed ferritic/bainitic steels are commonly in use. Weldments, joining different alloys, are characterised by a sometimes very sharp transition in microstructure, physical properties, chemical potential and as a result, in mechanical properties. Owing to the diversity of different types of joints, this chapter numerates only basic mechanisms relevant for creep damage in heat-resistant weldments. These mechanisms are: • •
mismatch of physical properties decarburisation/carburisation (gradient in chemical composition and potential).37
First, additional stresses may arise from a mismatch of physical properties. Different coefficients of thermal expansion and heat conductivity introduce
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Creep strength of welded joints of ferritic steels
489
(a)
(b)
17.8 Fracture surfaces of weld simulated coarse grained material after slow strain rate isothermal tensile testing (a) P23, (b) P24.34
thermal stresses adding to system stresses. This results in an additional localised loading caused by the combination of creep and fatigue close to the fusion line leading to premature failures. The more potent mechanisms influencing the time to rupture and failure location in creep-exposed dissimilar welds are
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variations within the microstructure owing to gradients in chemical composition. Formation of localised zones with inferior creep properties can be the result of diffusion of alloying elements across the interface either initially during PWHT or during service at elevated temperature.38 Investigating welds between a 1%CrMoV cast and a 12%CrMoV forged pipe (X20CrMoV12-1) utilising a 12%CrMoV filler material, Witwer40 showed the formation of a carbide seam in the fusion line area of the high alloyed weld material. In the low alloyed material, as a result, a carbon depleted soft region is formed. The width of carbide seam and carbon depleted region is strongly influenced by PWHT and service parameters.39 Regarding creep rupture properties, this narrow zone weak in creep is surrounded by zones of higher creep strength and leads to localised premature damage in the decarburised HAZ region of the low alloyed steel (Fig. 17.9).40
200 µm
17.9 Cross-weld creep sample of 1%CrMoV cast welded with 12%CrMoV filler fractured after 806 h at 180 MPa at 550°C in the decarburised CGHAZ of 1%CrMoV steel.40 The fracture surface on the left shows decarburised (bright appearance) areas.
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Creep strength of welded joints of ferritic steels
491
17.5.4 Type IV cracking Type IV cracking is defined as the formation and propagation of failures in the fine grained HAZ and the intercritically heated region of the HAZ. A strict differentiation between ICHAZ and FGHAZ is generally difficult because of very similar microstructural features in both regions. At this time, type IV cracking is considered as the major ‘end of life’ failure mechanism for ferritic creep-resistant steel weldments in the power generating industry. Therefore, this failure mechanism is of great interest and many researchers have investigated this life limiting phenomenon in welded components. Type IV cracking has been reported in low alloy ferritic/bainitic steels (1/2Cr1/2 Mo1/2V, 1CrMo, 1CrMoV, 11/4 Cr1/2Mo, 2CrMo, T/P22, T/P23, T/P24), as well as in ferritic/martensitic 9–12%Cr steels (P91, X20CrMoV121, P92, P122, E911).41–52 Time to rupture for failures in the type IV region is sensitive to the stress state and the applied loading direction. Cracking was observed in seam welds as well as in girth welds. While girth welds, loaded by system axial stresses, tend to fail by a ‘leak before break’ mechanism, failures in seam welds, loaded by the hoop stress, can be catastrophic.53 Review papers on type IV cracking have been written by Middleton and Metcalfe (1990),46 Ellis and Viswanathan (1998)53 and Francis et al. (2006).54 In contrast to creep failures in ferritic weldments at high stress levels which take place randomly either in BM, WM or the HAZ, the fracture location of weldments exposed to lower stress levels is shifted into the very narrow FGHAZ or intercritical HAZ region. Generally, type IV cracking can be seen as the result of a microstructural zone of material weak in creep strength surrounded by materials that are stronger in creep. This mismatch of creep properties leads to highly complex material behaviour.41 Most fractures have a macroscopic appearance typical of a low ductility failure. Figure 17.10 shows a cross-weld creep sample of E911 base material welded with matching filler and creep tested for 18 000 h at 600°C. Calculation of the failure strain using the overall gauge length results in very low strain values, typically below 10%. This leads to the assumption that the fracture mode is relatively brittle. However, significant creep deformation can be measured in a narrow part of the HAZ close to the unaffected base material, as investigated by Parker and Stratford.48 This fact indicates a strong strain localisation. Parker estimated the longitudinal strain of uniaxial specimens after creep exposure from measurements of the grains dimensions present in the fractured region. Changes in grain shape suggest a localised strain of about 20–30% and therefore the failures are in fact locally of a highly ductile nature. The failure mechanism of type IV cracking is governed by creep cavitation. Creep voids generally initiate at the sub-surface and grow by a diffusive
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4 mm
17.10 Cross-weld sample prepared from E911 pipe welded with matching filler and creep tested at 600°C for 18 000 h. The macroscopic low deformation fracture is located in the FGHAZ close to the unaffected base material.1
mechanism. Figure 17.11 shows the formation of a very narrow band of voids in the outer region of the HAZ in a cross-weld sample prepared from an E911 pipe welded with matching filler and creep tested for 14 000 h at 600°C. The combination of voids forms micro-cracks which in turn coalesce to form macro-cracks finally leading to premature failure. Preferred nucleation sites for voids are particle/matrix interfaces associated with inclusions or second phase particles.55 In the IC/FGHAZ region of weldments, carbides are only partially dissolved by the applied weld thermal cycle. Precipitation on retained large particles, such as M23C6, is favoured instead of fine reprecipitation on grain boundaries, in order to decrease the interfacial energy of the microstructure.56 Therefore, retained carbides coarsen more rapidly than those in the base material or weld metal during PWHT and are preferred nucleation sites for creep voids. Letofsky57 compared the evolution of precipitates at different stages of creep testing in the ICHAZ of G-X12CrMoWVNbN10.1.1 steel with that of the unaffected base material and matching weld metal using energy filtering transmission electron microscopy (EFTEM).57 His work is in agreement with other researchers showing that kinetics of microstructural changes, for example coarsening of precipitates and formation of the Laves- and Z-phase, are significantly faster in the ICHAZ than in all other regions of weldments (Fig. 17.12).25–62 General agreement also prevails concerning other microstructural features of the IC-/FGHAZ. Both regions susceptible to type IV cracking consist of fine equiaxed subgrains. A lath-type structure that is generally observed in all other parts of weldments is missing here.56 During creep exposure the subgrain microstructure tends to coarsen significantly.63,64 TEM investigations also revealed a significantly lower dislocation density in the vicinity of the
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17.11 Cross-weld sample prepared from E911 pipe welded with matching filler and creep tested at 600°C for 14 000 h. Localised formation of voids and their coalescence to macro-cracks at the outer region of the HAZ, observed by SEM, led to final fracture.1
G-X12 CrMoWVNbN 10.1.1, HAZ (soft zone)
As received
3.919 h 150 MPa
8.065 h 130 MPa
12.271 h 120 MPa
26.094 h 90 MPa
17.12 Evolution of precipitates in the ICHAZ of creep-exposed crossweld samples of G-X12 steel.62
FGHAZ compared to the other areas in welded joints. Further recovery of excess dislocations takes place during high temperature service.61 All these microstructural changes take place at higher velocity in the IC/FGHAZ than in other regions of weldments and contribute to the continuous weakening of this specific region, finally leading to the breakdown of creep rupture strength. The formation of a ‘soft zone’ with low hardness is not always directly related to low creep rupture stress. While in some investigations the location of final fracture corresponds with that of the soft zone,12,53,60,65 other researchers clearly distinguish between softened regions and regions with low creep rupture strength.25,56,61 The role of constraint The role of constraint in the creep weak IC/FGHAZ region from the adjacent stronger CGHAZ and BM is a key factor in clarification of type IV cracking mechanism. Simplifying the load condition to plain tensile loading of the dissimilar regions of the HAZ in series, the weakest part tries to deform transversely under high strains. If this region is sufficiently thin, it is constrained from doing so by the adjacent stronger material. As a result, a triaxial stress state predominates and prevents the weaker region from yielding. Under the
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creep regime, the situation is more diverse and is discussed intensively in the literature.66 Li et al. simulated creep behaviour of P122 cross-welds using a FEM model, based on Norton’s law representing four regions with different creep properties. They concluded that the FGHAZ has the highest equivalent strain, high first principal stress and hydrostatic pressure. The constraining effects of BM and CGHAZ prevent creep deformation in the FGHAZ. Therefore, a narrow FGHAZ is calculated to decrease the equivalent strain in the FGHAZ and is considered to be beneficial in decreasing the occurrence of creep voids.67 Experimental validation is provided by Japanese researchers reducing the HAZ width of P122 steel by electron beam welding (EBW). The creep rupture strength of EB weldments was improved by a factor of two compared to weldments produced by a GTAW process, although EBW specimens still failed by a type IV mechanism.68 Crack initiation was at lower creep strain in the EB joint because of larger local multi-axiality compared to the GTAW joint. Abe et al.25 simulated type IV creep crack growth behaviour in P122 weldments using a three-dimensional FE model taking the diffusive growth of creep voids into account. As a result, creep cracks grow faster in the centre of the specimen thickness than in specimens’ surface areas. Higher multiaxiality in the centre of the specimen thickness leads to a higher vacancy concentration which is consistent with experimental observation of initial creep voids in the centre region. This is supported by the acceleration of vacancy diffusion, formation and the growth of creep voids under multiaxial conditions for a welded joint specimen and as a result, the enhancement of damage accumulation by multiaxial stress states.69 Moreover, no mechanical constraint throughout the simulated HAZ creep specimen resulted in no void formation.25 An alternative approach to modelling type IV damage was introduced by Kimmins and Smith who suggest that constraint is relaxed by grain boundary sliding.70 Their experimental results suggested that additional grain boundary sliding results in a greater number of cavities. Once sliding is accommodated, the failure time for both, cross-weld and simulated type IV samples, was similar. Therefore, they concluded that material weak in creep deforms independently of adjacent stronger material and rather than using conventional continuum damage models in FE analysis, alternative models involving the mechanism of grain boundary sliding require development.41 Detectability As mentioned earlier, type IV damage initiates as creep cavitation subsurface at about half the time of the expected life of the weldments.63 However, cracks may form relatively late in life and crack growth once initiated may
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be very rapid.71 The remaining life for propagation throughout the wall can be less than 10 000 h.46 Surface bearing cracks do not appear until late in life. Therefore, surface examination of creep-exposed weldments only by replication techniques, penetration testing or eddy current testing can be misleading and severe damage in the subsurface regions can be overlooked.72 A sound residual life investigation on weldments can only be performed by advanced ultrasonic (UT) inspection or by highly sophisticated methods like time of flight diffraction (TOFD) which can detect creep voids and microcracks even at a life ratio (t/tr) of 0.5. In any case, life assessment methods for weldments vulnerable to type IV cracking have to include a qualitative damage classification scheme and a cavity density-based model.53,73,74
17.6
Implications for industries using welded creepresistant steels
Above it was shown that creep rupture strength of base metals and weldments can differ significantly. Therefore, it is of great importance to consider a weld strength factor (WSF) or a weld strength reduction factor (SRF) during the design stage of new components and the residual life evaluation of existing structures.75,76 In design codes for nuclear power plants, weld strength reduction factors have already been implemented. The opposite value of SRF, the WSF, is defined as the ratio of creep rupture stress at a certain time and temperature between weldments and base material. The European Creep Collaborative Committee (ECCC) defined WSF and SRF in their ECCC Recommendations as follows:77
WSF ( t , T ) =
Ru ( w )/ t / T Ru / t / T
[17.1]
SRF ( t , T ) =
Ru / t / T – Ru ( w )/ t / T Ru / t / T
[17.2]
where Ru/t/T is the creep rupture strength of base material samples at a time t and temperature T and Ru(w)/t/T is the creep rupture strength of cross-weld samples at a time t and temperature T. Up until now, a time and temperature independent weld strength reduction factor has been defined. As an example, in the German AD2000-Merkblatt B0, a constant WSF of 0.8 is given for components designed using creep rupture strength values. Investigations by many researchers have proved that this factor cannot be assumed to be constant over different temperature and stress levels, but rather depends on several factors like material type, stress level, temperature and time and can be either higher but also significantly lower.78,79 Lack of long-term creep rupture data makes the determination of accurate WSF very difficult. Extrapolation of WSF from short-term creep rupture tests holds the risk of
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Creep-resistant steels
overestimation of long-term creep properties and is not recommended. This emphasises the necessity for long-term creep rupture data of base material and welded joints. In Japan a ‘Strength of High Chromium Steel’ (SHC) committee since 2004 has systematically collected and analysed creep rupture data for base metal and welded joints. Takahashi and Tabuchi80 developed creep reduction factors for welded joints of HCM12A (P122) steel. The serious influence of temperature can be seen by a strong decrease in WSF from a value of 1 at 550°C down to 0.51 at 650°C.80 Tabuchi and Takahashi made similar investigations for 9Cr–1Mo (P91) steel.81 At 550°C a weld strength factor of 1 predicts no difference in creep rupture strength of welded components and plain base material, while at 650°C a WSF of 0.7 accounts a 30% decrease in creep rupture strength for weldments. In Fig. 17.13, WSF for several steel grades at different temperatures are shown.82 A reduction of up to 50% in creep rupture strength of weldments compared to base material creep rupture strength emphasises the importance of further research in this field and the correct consideration of WSF during the design stage of new structures.
17.7
Future trends
Up to the present, inferior long-term creep properties of weldments of heatresistant steels can only be taken into account during the design stage of thermal power plant components. The awareness of designers, engineers and operators of the risk of extrapolating results from short-term creep tests to longer times has already contributed to an increase in safety. General acceptance of the necessity of long-term creep testing data for cross-welds, weld metal and base metal for a reliable material selection is inevitable.83 1.0 Carbon steels
Weld strength factor
0.9
1-2%CrMo
0.8
G20Mo5
0.7 P91 0.6 G17CrMoV5–10 0.5 12CrMoV
P122
0.4 400
450
500 550 600 Temperature (°C)
650
700
17.13 Weld strength factors for the 100 000 h creep rupture strength of different steel grades.82
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Within the last few decades suppliers of welding consumables have improved their filler materials in terms of cleanliness, weldability and mechanical properties. As consequence, weld metal cracking has almost disappeared. In most cases the HAZ was identified as the weakest link in welded structures of creep-resistant steels. Shifting of welds to lower creep-stressed regions may improve the short time performance of welded components, but failures cannot be avoided over long times. The use of the most advanced nondestructive testing methods is necessary to detect damage as soon as possible in order to be able to react with appropriate countermeasures. The HAZ, as part of the base metal, already has to be taken into consideration in the design of new steel alloys. Weldability studies using HAZ simulation techniques can provide useful information on possible bottlenecks. Recently, in Japan at the National Institute for Materials Science (NIMS) a 9Cr–3W– 3Co steel with a reduced nitrogen level and controlled addition of boron was developed. In contrast to all creep-resistant steels used up until now, this steel does not show the formation of a fine grained region within the HAZ. Figure 17.14 shows the results of an electron backscatter diffraction pattern (EBSP) analysis of the grain size, as a function of distance from the fusion line, for a conventional P92 steel versus the new 9Cr–3W–3Co material.84 111
EBSP analysis (inverse pole figure map[001]) weld microstructure (GTAW width of HAS; 2.5 mm)
101
1.5mm from fusion line
101
0.5mm from fusion line
90 ppm B
Base metal
100µm
100µm
P92
100µm
100µm
100µm
100µm
17.14 Electron backscatter diffraction pattern (EBSP) analysis results for 9Cr–3W–3CoNB steel and P92 steel HAZ microstructures.84
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By the elimination of fine grains in the HAZ, the formation of creep damage by a type IV mechanism, which is strictly limited to fine grained regions, should be avoided. Up to a duration of 10 000 h, creep tests at 650°C are very promising and no difference in creep strength between cross-weld specimen and base material is shown. Although the mechanisms active in this steel are not fully understood, this might be a possible approach for the prevention of type IV cracking in ferritic creep-resistant steels.
17.8
References
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33 Nevasmaa P, Salonen J, Holmström S and Caminada S, ‘Heat affected zone toughness behaviour and reheat cracking susceptibility of thermally simulated microstructures in new P23 (7CrWVMoNb9-6) steel’, IIW Doc. IX-A-08-06, International Institute of Welding, 2006. 34 Dhooge A and Vekeman J, ‘New generation 21/4 Cr steels T/P23 and T/P24 – Weldability and high temperature properties’, IIW Doc. XI-810-04, International Institute of Welding, 2004. 35 Nawrocki J G, DuPont J N, Robino C V and Marder A R, ‘The stress-relief cracking susceptibility of a new ferritic steel – Part 1: Single-pass heat affected zone simulations’, Welding J, 2000 355–362. 36 Nawrocki J G, DuPont J N, Robino C V and Marder A R, ‘The stress-relief cracking susceptibility of a new ferritic steel – Part 2: Multiple-pass heat affected zone simulations’, Welding J, 2001 80 (1), 18–24. 37 Buchmayr B, Cerjak H, Witwer M, Maile K, Theofel H and Eckert W, ‘Experimental and numerical investigations of the creep behaviour of the dissimilar weldment GS17CrMoV5-11 and X20CrMoV12-1’, Steel Res, 1990 61 (6), 268–273. 38 Helander T Z, Andersson H C M and Oskarsson M, ‘Structural changes in 122.25%Cr weldments – an experimental and theoretical approach’, Materials at High Temperature, 2000 17 (3), 389–96. 39 Buchmayr B, Cerjak H, Kirkaldy J S and Witwer M, ‘Carbon diffusion and microstructure in dissimilar Cr-Mo-V-welds and their influence on the mechanical properties’, in 2nd International Conference Trends in Welding Research, ASM International, Gatlinburg, 1989. 40 Witwer M, Untersuchungen an Mischschweissverbindungen warmfester CrMoVStähle, Doctoral Thesis, Graz University of Technology, 1989. 41 Brett S J, ‘Cracking experience in steam pipework welds in National Power’, in Conf VGB Conference – Materials and Welding Technology in Power Plants 1994, VGB, Essen, 1994. 42 Gooch D J and Kimmins S T, ‘Type IV cracking in 1/2Cr1/2Mo1/4V/2 1/4Cr1Mo weldments’, in 3rd International Conference Creep and Fracture of Engineering Materials and Structures, Maney Publishers, Swansea, 1987. 43 Smith D J, Walker N S and Kimmins S T, ‘Type IV creep cavity accumulation and failure in steel welds’, International J Pressure Vessels & Piping, 2003, 80, 617– 627. 44 Brear J M, Fairman A, Middleton C J and Polding L, ‘Predicting the creep life and failure location of weldments’, Key Eng Mater, 2000, 171–174, 35–42. 45 Fujibayashi S and Endo T, ‘Creep behaviour of a low alloy ferritic steel weldment’, in 9th International Conference Creep and Fracture of Engineering Materials and Structures, Cambridge, Maney Publishing, 2001. 46 Middleton C J and Metcalfe E, ‘A review of laboratory Type IV cracking data in high chromium ferritic steels’, in International Conference Steam Plants for the 1990’s, IMechE, London, 1990. 47 Brühl F, Cerjak H, Schwaab P and Weber H, ‘Metallurgical investigation on the base material and weldments of the 9% chromium X10CrMoVNb91’, Steel Res, 1991 62 (2), 75–82. 48 Parker J D and Stratford G C, ‘Strain localisation in creep testing of samples with heterogeneous microstructures’, International J Pressure Vessels & Piping, 1996 68, 135–143. 49 Tezuka H and Sakurai T, ‘A trigger of Type IV damage and a new heat treatment
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procedure to suppress it. Microstructural investigations of long-term ex-service CrMo steel pipe elbows’, International J Pressure Vessels & Piping, 2005, 82, 165– 174. Shibli I A and Le Mat Hamata N, ‘Creep and fatigue crack growth in P91 weldments’, in 9th International Conference Creep and Fracture of Engineering Materials and Structures, Woodhead Publishing, Cambridge, 2001. Ennis P J, ‘The mechanical properties and microstructures of 9% chromium steel P92 weldments’, OMMI, 2002, 1 (2), 1–23. Takemasa F, Nonaka I, Ito T, Saitou K, Miyachi Y and Kagiya Y, ‘Type IV creep damage analysis for full size component test on welded P91 boiler hot reheat piping’, in International Conference Elevated Temperature Design and Analysis, Nonlinear Analysis and Plastic Components, ASME, San Diego, 2004. Ellis F V and Viswanathan R, ‘Review of Type IV cracking in piping welds’, in 1st International Conference Integrity of High Temperature Welds, IOM, London, 1998. Francis J A, Mazur W and Bhadeshia H K D H, ‘Type IV cracking in ferritic power plant steels’, Mater Sci Techn, 2006, 22 (12), 1387–1395. Shinozaki K, Li D, Kuroki H, Harada H and Ohishi K, ‘Analysis of degradation of creep strength in heat affected zone of weldment of high Cr heat-resisting steels based on void observation’, ISIJ International, 2002, 42 (12), 1578–1584. Hasegawa Y, Muraki T and Ohgami M, ‘Metallurgical investigation of a Type IV damage at the heat affected zone of weld for tungsten containing martensitic heat resistant steels’, in International Conference Experience with Creep-Strength Enhanced Ferritic Steels and New and Emerging Computational Methods, ASME, San Diego, 2004. Letofsky E, ‘Microstructure aspects of creep resistant welded joints’, IIW Doc. IX2055-03, International Institute of Welding, 2003. Albert S K, Matsui M, Watanabe T, Hongo H, Kubo K and Tabuchi M, ‘Microstructural investigations on Type IV cracking in a high Cr steel’, ISIJ International, 2002, 42 (12), 1497–1504. Kondo M, Tabuchi M, Tsukamoto S, Yin F and Abe F, ‘Suppression of Type IV failure in high-B low-N 9Cr-3W-3Co-NbV steel welded joint’, in 4th International Conference Advances in Materials Technology for Fossil Power Plants, ASM International, Ohio, 2004. Watanabe T, Tabuchi M, Yamazaki M, Hongo H and Tanabe T, ‘Creep damage evaluation of 9Cr-1Mo-V-Nb steel welded joints showing Type IV fracture’, International J Pressure Vessels & Piping, 2006, 83, 63–71. Tabuchi M, Watanabe T, Kubo K, Matsui M, Kinugawa J and Abe F, ‘Creep crack growth behaviour in the HAZ of weldments of W containing high Cr steel’, International J Pressure Vessels & Piping, 2001, 78, 779–784. Letofsky E and Cerjak H, ‘New quantitative microstructural characterisation on creep tested welded joints’, in 6th International Conference Trends in Welding Research, ASM, Pine Mountain, 2002. Abe F and Tabucshi M, ‘Microstructure and creep strength of welds in advanced ferritic power plant steels, Science Technol. Welding Joining, 2004, 9 (1), 22–30. Eggeler G, Ramteke A, Coleman M, Chew B, Peter G, Burblies A, Hald J, Jefferey C, Rantala J, Dewitte M and Mohrmann R, ‘Analysis of creep in a welded P91 pressure vessel’, International J Pressure Vessels & Piping, 1994, 60, 237–257. Masuyama F, Komai N and Sasada A, ‘Creep failure experience in welds of advanced steel boiler components’, IIW Doc. XI-795-04, International Institute of Welding, 2004.
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66 Kimmins S T, Coleman M C and Smith D J, ‘An overview of creep failure associated with heat affected zones of ferritic weldments’, in 5th International Conference Creep and Fracture of Engineering Materials and Structures, IOM, London, 1993. 67 Li D, Shinozaki K and Kuroki H, ‘Stress-strain analysis of creep deterioration in heat affected weld zone in high Cr ferritic heat resistant steel’, Mater Sci Technol, 2003, 19, 1253–1260. 68 Tabuchi M, Albert S K, Kondo M, Watanabe T and Abe F, ‘Study on the creep fracture of advanced high Cr steel weldment’, Ultra-Steel Conference: Requirements from New Design of Constructions, Proceedings of 7th Workshop on Ultra Steel, 24– 25 June 2003, Tsukuba, Japan, NIMS, Tsukuba, 2003. 69 Perrin I J and Hayhurst D R, ‘Continuum damage mechanics analyses of type IV creep failure in ferritic steel crossweld specimens’, International J Pressure Vessels & Piping, 1999, 76 (9), 599–617. 70 Kimmins S T and Smith D J, ‘On the relaxation of interface stresses during creep of ferritic steel weldments’, J. Strain Anal, 1998, 33 (3), 195–206. 71 Ito T, Nonaka I and Takemasa F, ‘Full size internal pressure creep test and life evaluation for welded Gr.91 hot reheat piping’, IIW Doc. XI-796-04, International Institute of Welding, 2004. 72 Lundin C D and Prager M, ‘A new approach to investigation into Type IV cracking susceptibility’, in ASME/JSME Pressure Vessels and Piping Conference, 26–30 July, San Diego, CA, ASME, 1998. 73 Parker J and Bisbee L, ‘Girth weld cracking in high temperature headers’, in International Conference Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, DEStech Publications, London, 2005. 74 Brett S J, ‘The management of weld cracking in high-temperature CrMoV pipework systems’, in International Conference Assuring it’s safe: Integrating Structural Integrity, Inspection and Monitoring into Safety and Risk Assessment, IMEchE, 1998. 75 Allen D J, Harvey B and Brett S J, ‘Fourcrack – An investigation of the creep performance of advanced high alloy steel welds’, in International Conference Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, DEStech Publications, London, 2005. 76 Tu S T, Segle P and Gong J M, ‘Strength design and life assessment of welded structures subjected to high temperature creep’, International J Pressure Vessels & Piping, 1996, 66, 171–186. 77 Morris P F, ‘Terms and Terminology for weld creep testing’, ECCC Recommendations, 2001, 2(IIb). 78 Masuyama F, ‘Creep rupture life and design factors for high strength ferritic steels’, in International Conference Creep & Fracture in High Temperature Components – Design & Life Assessment Issues, DEStech Publications, London, 2005. 79 Sandstrom R and Tu S T, ‘The effect of multiaxiality on the evaluation of weldment strength reduction factors in high-temperature creep’, Trans ASME, 1994, 116, 76– 80. 80 Takahashi Y and Tabuchi M, ‘Evaluation of creep strength reduction factors for welded joints of HCM12A (P122)’, in 2006 ASME Pressure Vessels and Piping Division Conference, ASME, Vancouver, 2006. 81 Tabuchi M and Takahashi Y, ‘Evaluation of creep strength reduction factors for welded joints of modified 9Cr-1Mo steel (P91)’, 2006 ASME Pressure Vessels and Piping Division Conference, ASME, Vancouver, 2006. 82 Schubert J, Klenk A and Maile K, ‘Determination of weld strength factors for the
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creep rupture strength of welded joints’, in International Conference Creep and Fracture in High Temperature Components – Design & Life Assessment Issues, DEStech Publications, London, 2005. 83 Klenk A, ‘Creep testing of weldments: practices and investigations on the effects of sampling and size on creep test results for weldments’, ECCC Recommendations, 2005, 3(II) Appendix 1. 84 Kondo M, Tabuchi M, Tsukamoto S, Yin F and Abe F, ‘Suppressing type IV failure via modification of heat affected zone microstructures using high boron content in 9Cr heat resistant steel welded joints’, Sci Technol Welding and Joining, 2006 11 (2), 216–223.
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18 Fracture mechanics: understanding in microdimensions M . T A B U C H I, National Institute for Materials Science (NIMS), Japan
18.1
Introduction
For economic reasons, many operators plan to use power and chemical plants beyond their originally predicted operating life. At the same time, operating temperatures and loading conditions for high temperature components are tending to become more severe in attempts to improve efficiency, save energy and reduce carbon dioxide emissions. Under these circumstances, accurately predicting the lifespan and remaining life of high temperature components becomes much more important. Creep voids can form in thick components operating under stress and temperature gradients and these voids grow into micro and macro cracks, eventually causing fracture. It is therefore important to be able to predict crack initiation, growth and time to fracture under high temperature creep conditions. The application of fracture mechanics to high temperature creep crack growth first began in the 1970s. A large amount of research concerning fracture mechanics parameters, which control the stress–strain field in front of the creep crack, has been conducted using several types of specimens, including CT, CCT, SEN and DEN. The C* parameter,1,2 based on the J integral,3 is generally used to characterize the stress-strain field in front of the crack tip and to correlate crack growth rate under extensive creep conditions. C(t),4,5 Ct6 and Q*7 can also be used to evaluate creep crack growth properties. The standard test method for measuring and evaluating high temperature creep crack growth for creep ductile materials, ASTM E1457, was established in 19928 and revised to include creep brittle materials in 2000,9 based on the results of VAMAS TWA11 and 19 projects. Recently, the application of non-linear fracture mechanics to cracked components has become an important subject.10
18.2
Non-linear fracture mechanics
Just as J characterizes the crack tip fields in an elastic or elastic–plastic material, the stress and strain rate distribution in front of a crack tip under 504 WPNL2204
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steady-state creep conditions can be characterized by a path-independent energy rate line integral, C* integral,1,2 which is defined by analogy to J.3 The definition of C* is the following line integral on a contour Γ surrounding a crack tip shown in Fig. 18.1: C* =
W˙ =
∫
∫
∂u˙ Γ W˙ dy – Ti i ds ∂x
ε˙ ij
[18.1]
σ ij d ε˙ ij
[18.2]
0
where W˙ is strain energy rate density, σij and ε˙ ij are the stress and strain rate tensors respectively, Ti is the outward traction vector, u˙ i is the displacement rate vector and ds is the arc length along contour path Γ. C* can be determined from a load–displacement rate diagram for a creeping material. C* is defined from the shaded area of Fig. 18.2: y
ds r θ
x
Γ
18.1 Integration contour around the crack tip for the definition of non-linear fracture mechanics parameter C*. P d U*
Load
a
a + da
Creep displacement rate
18.2 Load–creep displacement rate responses for different crack sizes of non-linear material.
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C* = – 1 dU * B da
U* =
∫
δ˙
[18.3]
P dδ˙
[18.4]
0
where U* is the potential energy rate, a is the crack length, δ˙ is load line displacement rate and B is the thickness. For power-law creep, the C* parameter can be described as:
C* =
Pδ˙ F B (W – a)
[18.5]
where W is the width of the plate and F is a non-dimensional factor which depends on geometry and creep exponent n of the power-law creep. For CT specimens, the following calculation method for F is recommended in ASTM E1457:8,9 F=
n 2 + 0.522 W – a W n + 1
[18.6]
Usually, the relationship between creep crack growth rate (da/dt) and C* is given as:
da = DC* φ dt
[18.7]
where D and φ are constants. Figure 18.3 shows the relationship between creep crack growth rates and C* obtained by Japanese research groups for 1Cr–Mo–V turbine rotor steel using several sizes of CT specimen.11 C* correlates to creep crack growth rate better than net section stress (σnet) and stress intensity factor (K). In the relationship between da/dt and C* in Fig. 18.3, transient behaviour (the tail part) is observed in the early stages of crack growth, where the values of C* and da/dt decrease. Only after the steady-state creep crack growth stage can da/dt be expressed as Equation [18.7]. Transient crack growth behaviour is considered to occur through a combination of stress redistribution and primary creep. ASTM E1457 recommends evaluating only the data for which time exceeds transition time, tT, by C*. The transition time is estimated as: tT =
K 2 (1 – ν 2 ) E ( n + 1) C*
[18.8]
where E is Young’s modulus, ν is Poisson’s ratio and K is the stress intensity factor. The data where time is lower than tT are correlated using Ct.6
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da/dt (mm h–1)
100
10–1
10-2
10–3 10–1
100 101 C* (kJ m–2 h–1)
102
18.3 Relationship between creep crack growth rate versus C* parameter for a 1Cr–Mo–V turbine rotor steel.
18.3
Effect of mechanical constraint
The effects of crack tip constraint arising from variations in specimen size, geometry and material ductility can influence the relationship between da/dt and C* in Equation [18.7].9 Figure 18.4 shows the effect of specimen size on da/dt versus C* for 1Cr–Mo–V steel using large CT (W = 254 mm) and standard CT (W = 50.8 mm) specimens.12 Note that an increase in specimen thickness increases crack growth rate. The da/dt of the thickest specimen (B = 63.5 mm) with side grooves (SG) is about six times higher than that of the thinnest specimen (B = 6.35 mm) without SG. Figure 18.5 shows the fracture surface of large CT specimens, 63.5 mm in thickness. Comparing non-SG specimens with SG specimens, while the crack length at the thickest part was nearly the same, the crack length at the surface was one-fifth that in SG specimens. The reason why side grooves and thickness accelerate creep crack growth rate is considered to be due to mechanical constraint ahead of the crack tip. The stress and strain rate field under power-law creep, ε˙ = Aσ n , is given as follows (HRR field):13,14
σ ij = C* I n Ar
1/ n+1
ε˙ ij = A C* I n Ar
σ˜ ij ( θ, n ) n / n+1
ε˜˙ ij ( θ, n )
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[18.10]
Creep-resistant steels 1
da/dt (mm h–1)
508
W 254.0, 254.0, 254.0,
B 63.5, S.G. 63.5, Non S.G. 12.7, S.G.
0.1
0.01
Pl
an
e
st
ra
in
Pl
0.001 0.01
0.1
a
ne
s
e tr
ss
W 50.8, 50.8, 50.8, 50.8,
B 25.4, 12.7, 6.35, 6.35,
1 C* (kJ m–2 h–1)
10
S.G. S.G. S.G. Non S.G. 100
18.4 Effect of specimen thickness, width and side groove on the creep crack growth rate of 1Cr–Mo–V steel.
W = 254 mm, B = 63.5 mm without SG
W = 254 mm, B = 63.5 mm with SG
18.5 Fracture surface of large CT specimens with 63.5 mm in thickness after a creep crack growth test at 811 K for 1Cr–Mo–V steel.
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~ where σ˜ ij and ε˙ ij are non-dimensional functions of θ and n, and In is a nondimensional function of n. From Equation [18.10], the following equation to characterize creep crack growth rate is derived:15,16 n / n+1 da n + 1 1/ n+1 C* = ( Ar ) c In dt ε *f
[18.11]
where ε *f is creep ductility and rc is the size of the creep process zone. Under plane stress conditions, ε *f is the creep ductility obtained from smooth round bar creep specimens. The creep crack growth rate calculated using Equation [18.11] is shown in Fig. 18.4 as the line that corresponds well with the experimental data for the thinnest non-SG specimen (B = 6.35 mm). When the specimen thickness increases and side grooves are introduced, ε *f ahead of the crack tip decreases and the crack growth rate increases. The crack growth rate under plane strain conditions in Fig. 18.4 is obtained by assuming that creep ductility under such conditions is 1/50 of uniaxial ductility.15 The ε *f under a multiaxial stress state can be calculated as follows:17,18 ε *f n – 1/2 n – 1/2 σ m = sinh 2 sinh 2 εf 3 n + 1/2 n + 1/2 σ e
[18.12]
where ε *f and εf are creep ductility under multiaxial and uniaxial condition, and σm and σe are hydrostatic stress and equivalent von Mises stress, respectively. It is therefore necessary to keep the component dimensions in mind when applying the experimental data to large-scale structural components.9
18.4
Effect of microscopic fracture mechanisms
Creep crack growth properties are affected by microscopic fracture mechanisms, which change according to temperature and loading conditions. Over long operating times, creep crack growth by grain boundary cavitation is often observed. In most high-temperature, thick-section components operating under multiaxial stress conditions, creep cracks grow accompanied by micro-cavities, voids and micro-cracks ahead of the crack tip. It is therefore important to take microscopic aspects into account when evaluating creep crack growth properties with the aim of developing more accurate predictions of operating life. Figure 18.6 shows the microscopic features of creep cracks observed in CT specimens of Alloy 800H.19 Three types of creep crack growth, brittle intergranular fractures caused by wedge-type cracks (W-type) at lower temperature, ductile transgranular fractures (T-type) and void-type intergranular (V-type) fractures at higher temperature, were observed, depending on the testing temperature and loading conditions. Large creep damage zones ahead
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Creep-resistant steels (a)
Wedge-type creep crack (b)
Transgranular creep crack (c)
Void-type creep crack
18.6 Microscopic features of creep cracks of Alloy 800H tested at (a) 873K, (b) 973K and (c) 1073K.
of the crack tip were observed for V-type crack growth. The relationship between da/dt and C* for Alloy 800H at various temperatures is shown in Fig. 18.7. These relationships are dependent on the microscopic crack growth mechanisms; creep crack growth rate for W-type and V-type fracture was higher than that for T-type fracture. For 316 stainless steel, the same tendency was reported.20 Figure 18.8 shows the da/dt at a constant C* value (1kJ/m2h) plotted against creep ductility in uniaxial creep tests for Alloy 800H and 316 stainless
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Creep crack growth rate, da/dt (mm h–1)
1
0.1 Temp. (K)
Load (kN)
Fracture mode
873
13.69
Wedge-type
0.01
973
11.47 Wedge-type 11.71 Transgranular 8.98 Transgranular 7.62 Transgranular
1073
4.36 3.11
Transgranular Void-type
2.61
Void-type
923 0.001
Alloy 800H 0.0001 0.01
0.1
1 C* (kJ m–2 h–1)
10
100
18.7 Relationship between da/dt versus C* parameter for Alloy 800H at various temperatures.
da/dt at C* = 1 (kJ m–2 h–1) (mm h–1)
0.08 Wedge-type Transgranular Void-type
W
0.06
Alloy800H 316 stainless
V
0.04
W
V
W
V
0.02
V T
T
0 0
20
T
40 60 80 Reduction in area (%)
100
18.8 Relationship between da/dt at constant C* value 1 kJ m–2h–1 versus creep ductility for Alloy 800H and 316 stainless steel.
steel,21 which is dependent on creep fracture mechanisms under these conditions. For W-type and T-type fractures, da/dt is inversely proportional to creep ductility, as shown by the solid curve in Fig. 18.8. Therefore, the creep crack growth rate for W-type and T-type fractures can be predicted according to Equation [18.11]. On the other hand, for V-type fractures, da/ dt is faster than that predicted from creep ductility. When creep damage is only just in front of the crack tip, as in W-type and T-type fractures, da/dt can be characterized by the creep ductility. Where there is a large creep damage
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zone formed ahead of the crack tip, as in voids and micro-cracks, da/dt accelerates faster than predicted from creep ductility. Analytical models for creep crack growth taking into account the effect of creep damage formed ahead of the crack tip have been proposed.5,16 According to these analytical models, increase in creep damage (void density) in front of the crack tip accelerates the creep crack growth rate. It can be considered that, under creep conditions, the diffusion of vacancies towards cracks and voids would contribute to crack growth. The vacancy diffusion equation under stress gradient is given as follows:22,23
(
∂C = D ∇ ∇ C + C ∇v RT ∂t
)
[18.13]
ν = –σp∆V
[18.14]
where C is the vacancy concentration, D is the diffusion coefficient, R is the gas constant, T is the absolute temperature, σp is the hydrostatic stress and ∆V is the change to molar volume caused by vacancy diffusion. In this equation, vacancy diffusion is controlled by the hydrostatic stress gradient. The modified equation to calculate realistic vacancy diffusion and concentration using weight coefficients α1 and α2 is as follows:24,25
∂C = α D ∇ 2 C + α D ∇ C∇ ν 1 2 RT ∂t
[18.15]
The vacancy diffusion equation is solved for three dimensional CT specimen models using the hydrostatic stress gradient ∇σp computed by the finite element method (FEM).21 Figure 18.9 shows the computed examples for changes in vacancy concentration ahead of the crack tip during creep. The vacancy concentration increases faster and to higher levels in the centre of
Vacancy concentration C/C0
3.0
D = 1.5 × 10–9 m2s–1 D = 1.5 × 10–10 m2 s–1
2.5
2.0
Centre
Surface
1.5
Centre
1.0
Surface Alloy800H
0.5 0
100
200 Time (h)
300
18.9 Computed results for changes in vacancy concentration ahead of the crack tip during creep.
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specimen than at the surface. This is consistent with experimental results where creep voids were often observed in the thickest part of the specimen, where stress multiaxiality is large. When D is low, at lower temperatures, cracks propagate before vacancies are accumulated, which corresponds to transgranular-type crack growth. Creep crack growth properties can be characterized by creep ductility. Creep damage formed in front of the crack tip accelerates creep crack growth. Vacancy diffusion, void formation, and crack initiation and propagation are accelerated under multiaxial stress fields at high temperatures.
18.5
Type IV creep crack growth in welded joints
In an attempt to reduce carbon dioxide emissions and save energy in thermal power plants, the steam pressure and temperature conditions under which boiler components are expected to operate is increasing. High Cr ferritic steels (9–12%Cr steels), such as P91 (9Cr–1Mo–VNb), P92 (9Cr–0.5Mo– 1.8W–VNb) and P122 (11Cr–0.4Mo–2W–CuVNb) steels, which have a tempered martensite structure, are now used for boiler components in ultrasupercritical (USC) power plants operating at around 873 K because of their high creep strength. The creep strength of weldment in these steels, however, decreases at above 873 K owing to type IV creep damage formed in the heataffected zone (HAZ).26–30 It is important to predict initiation and growth of type IV creep damage in the HAZ of the weldment. The creep properties of HAZ are usually investigated using simulated HAZ specimens, which are produced by rapid heating to peak temperatures around the AC1 to AC3 transformation temperature and cooling. Figure 18.10 shows the relationship between creep rupture times and peak temperatures during simulated HAZ heat treatment for P122 steel at 923 K.29 The creep rupture time shows a minimum value for the specimens heated to around the AC3 transformation temperature. The microstructure of HAZ heated to AC3 is characterized by a fine-grained structure without lath martensite, with a grain size of about 5 µm. The creep rupture time for the simulated HAZ specimen heated to AC3 was about one-fifth that of base metal. Therefore, lower creep strengths in fine-grained HAZ structures are considered to be the primary cause of type IV failure. The width of the HAZ is very narrow in the welded joint. Therefore the creep deformation of fine-grained HAZ (with low creep strength) is mechanically constrained by weld metal and base metal, which have higher creep strengths. Figure 18.11 shows the creep crack propagation profile in a welded joint of P92 steel. This test was conducted using a CT specimen with a notch tip located in the HAZ and fatigue pre-cracked. A very sharp creep crack, which propagates in fine-grained HAZ, type IV creep crack, is observed. Creep voids were observed around the main cracks in welded joint specimens.
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Ac1
Ac3
10 000
Time to rupture (h)
80 MPa
1000
100 MPa
120 MPa
P122 steel 923 K
100
1100
1200 1300 1400 Peak temperature (K)
1500
18.10 Changes in creep rupture time as a function of peak temperature of simulated HAZ heat treatment for P122 steel at 923 K.
Base metal Creep crack HAZ Pre-crack
Weld metal
18.11 Creep crack propagation profile in welded joint of P92 steel using CT specimen.
The C* line integral of Equation [18.1] can be calculated by FEM using power-law creep obtained from uniaxial creep tests for simulated HAZ, base metal and weld metal. Figure 18.12 shows the computed C* value for base metal, welded joints and fine-grained HAZ for P122 steel plotted against loading time.31 Here, the simulated HAZ specimen has a fine-grained structure
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1 P122 923 K
C* (kJ m–2 h–1)
0.1
Simulated HAZ
Welded joint 0.01 Base metal
0.001 1
10
100 Time (h)
1000
18.12 Computed C* value by FEM for base metal, welded joints and simulated HAZ for P122 steel.
and the width of fine-grained HAZ for welded joints is 1.2 mm. The computed C* value for the simulated HAZ is about one order higher than that of base metal for the same loading conditions. The C* value of welded joints decreases as the width of HAZ decreases. Figure 18.13 shows the creep crack growth behaviour of welded joints, base metal and simulated fine-grained HAZ for P122 steel tested using CT specimen at 923 K under the same loading conditions.31 For the welded joints, the early crack growth rate is nearly the same as that in base metal, but the crack grows rapidly in the accelerating stage after the incubation period. Consequently, the crack growth life of welded joints is about onethird that of base metals. The relationship between creep crack growth rate da/dt and C* in welded joints and base metal for P122 steel is shown in Fig. 18.14. The creep crack growth rates of the welded joints are higher than those for base metal in higher C* regions, but the differences in da/dt between welded joints and base metal are small in the lower C* region. From this result, the prediction of crack initiation time (incubation period), and when crack growth accelerates, is clearly important for the evaluation of fracture life of the welded joints. The crack growth behaviour of welded joints mentioned above can be correlated with stress multiaxiality in HAZ. Creep deformation for finegrained HAZ with low creep strength is mechanically constrained by the weld metal and base metal, which have higher creep strength. Figure 18.15 shows the computed stress triaxial factor TF in front of the crack tip for welded joints and base metal, P92 steel, using FEM for three-dimensional CT specimens with side grooves.32 TF is calculated as follows:
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Creep crack length (mm)
Simulated HAZ 6.0
Welded joint
4.0
Base metal 2.0
0.0 0
500
1000 Time (h)
1500
2000
18.13 Creep crack growth behavior of welded joints, base metal and simulated fine-grained HAZ for P122 steel at 923 K for the same loading condition. 1 P122 steel 923 K
da/dt (mm h–1)
0.1
Base metal Welded joint Simulated HAZ
0.01
0.001
0.0001 0.001
0.01
0.1 C* (kJ m–2 h–1)
1
10
18.14 Relationship between creep crack growth rate da/dt versus C* parameter of welded joints and base metal for P122 steel.
TF =
σ1 + σ 2 + σ 3 σ eq
[18.16]
where σ1, σ2, and σ3 are principal stress and σeq is equivalent von Mises stress. In Fig. 18.15, the value of TF for welded joints is larger than that for base metal. The creep deformation in the HAZ of welds is small in the initial
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12 3D elastic–plastic creep FEM analysis Model : CT specimen with SG Temp : 650°C Load : 10 000 N (Kin = 15 MPa m1/2)
10
Base metal Heat affected zone
Triaxial factor TF
8
6
Increase in creep time up to 500 h
4 Increase in creep time up to 500 h 2
0 0
5 10 Distance from the crack tip (mm)
15
18.15 Stress triaxial factor TF ahead of the crack tip of threedimensional CT specimen with S.G. for welded joint and base metal of P92 steel.
stages owing to mechanical constraint under multiaxial stress conditions. However, void formation, growth and crack initiation are accelerated under multiaxial stress conditions for welded joint specimens. Predicting the lifespan of welds therefore requires evaluating the initiation of creep voids and cracks under multiaxial stress conditions.
18.6
References
1 K. Ohji, K. Ogura and S. Kubo, Proceedings 1974 Symposium on Mechanical Behavior of Materials 1, The Society of Materials Science, Japan, 1974, 455–466. 2 J. D. Landes and J. A. Begley, ASTM STP 590, ASTM, 1976, 128–148. 3 J. R. Rice, Trans. ASME, J. Appl. Mech., 1968, 35, 379–385. 4 R. Ehlers and H. Riedel, Proceedings of ICF5, Volume 2, Pergamon Press, 1981, 691–698. 5 H. Riedel, Fracture at High Temperatures, Springer-Verlag, Berlin, 1987. 6 A. Saxena, ASTM STP 905, ASTM, 1986, 185–201. 7 A. T. Yokobori, Jr. and T. Yokobori, Proceedings International Conference on Creep, JSME, 1986, 135–140. 8 ASTM E1457-92: Standard Test Method for Measurement of Creep Crack Growth Rate in Metals, ASTM, 1992 1031–1043. 9 ASTM E1457-00: Standard Test Method for Measurement of Creep Crack Growth Rate in Metals, ASTM, 2000, 936–950.
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10 Code of Practice for Creep/Fatigue Testing of Cracked Components, ISO/TTA E2006 (in press). 11 T. Yokobori, C. Tanaka, K. Yagi, M. Kikagawa, A. Fuji, A. T. Yokobori, Jr. and M. Tabuchi, Mater. High Temp., 1992, 10, 97–107. 12 M. Tabuchi, K. Kubo and K. Yagi, Eng. Fract. Mech., 1991, 40, 311–321. 13 J. W. Huchinson, J. Mech. Phys. Solids, 1968, 16, 13–31. 14 J. R. Rice and G. F. Rosengren, J. Mech. Phys. Solids, 1968, 16, 1–12. 15 K. M. Nikbin, D. J. Smith and G. A. Webster, Trans. ASME J. Engng. Mater. Tech., 1986, 108, 186–191. 16 G. A. Webster and R. A. Ainsworth, High Temperature Component Life Assessment, Chapman & Hall, London, 1994. 17 A. C. F. Cocks and M. F. Ashby, Met. Sci., 1980, 14, 395–402. 18 M. Yatomi, K.M. Nikbin and N.P. O’Dowd, Int. J. Pressure Vessels Piping, 2003, 80, 573–583. 19 M. Tabuchi, K. Kubo and K. Yagi, Tetsu-to-Hagane, 1993, 79, 732–738 in Japanese. 20 M. Tabuchi, K. Yagi and T. Ohba, ISIJ International, 1990, 30, 847–853. 21 M. Tabuchi, J. Ha, H. Hongo, T. Watanabe and A. T. Yokobori Jr., Metall. Mater. Trans. A, 2004, 35A, 1757–1764. 22 H. P. Leeuwen, Eng. Fract. Mech., 1974, 6, 141–161. 23 H. P. Leeuwen, Eng. Fract. Mech., 1977, 9, 951–974. 24 A. T. Yokobori, Jr., T. Nemoto, K. Sato and T. Yamada, Eng. Fract. Mech., 1996, 55, 47–60. 25 A. T. Yokobori, Jr., Y. Chinda, T. Nemoto, K. Sato and T. Yamada, Corrosion Sci., 2002, 44, 407–424. 26 G. Eggeler, A. Ramteke, M. Coleman, B. Chew, G. Peter, A. Burblies, J. Hald, C. Jefferey, J. Rantala, M. deWitte and R. Mohrmann, Int. J. Pressure Vessels Piping, 1994, 60, 237–257. 27 F. Masuyama, M. Matsui and N. Komai, Key Engng Mater., 171–174, 2000, 99–107. 28 T. H. Hyde, W. Sun and A. A. Becker, Int. J. Pressure Vessels Piping, 2001, 78, 765– 77. 29 M. Tabuchi, T. Watanabe, K. Kubo, M. Matsui, J. Kinugawa and F. Abe, Int. J. Pressure Vessels Piping, 2001, 78, 779–784. 30 Y. Hasegawa, T. Muraki and M. Ohgami, Tetsu-to-Hagane, 2004, 90, 609–617 in Japanese. 31 M. Tabuchi, H. Hongo, T. Watanabe and A. T. Yokobori, Jr., J. ASTM Inte, 2006, 3 (5), online. 32 R. Sugiura, A. T. Yokobori, Jr., M. Tabuchi and T. Yokobori, Eng. Fract. Mech., 2007, 74, 868–881.
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19 Mechanisms of oxidation and the influence of steam oxidation on service life of steam power plant components P. J. E N N I S and W. J. Q U A D A K K E R S Forschungszentrum Juelich GmbH, Germany
19.1
Introduction
In the development of commercial 9% chromium steels, the principal aim has been a significant increase in the creep rupture strength, to enable application temperatures of 600°C and even higher. Creep strength targets were achieved, so that with the currently available commercial steels P91/ T91, E911 and P92/T92, metal temperatures of up to 625°C can be considered, based on the creep rupture strength. Creep rupture testing to test durations of many thousands of hours have been carried out on all these steels and it was found that a chromium content of 9mass% promoted the formation of highly protective oxide scales that provided excellent protection against oxidation in air. However, the service environments of steam power plant components are either steam or combustion gases, which also contain significant levels of water vapour. Investigations into the effect of such service environments on oxide scale formation showed that in the presence of water vapour, the oxidation rates of steels containing around 9mass% of chromium were considerably accelerated. In contrast to thin protective scales based on haematite (Fe2O3) found after exposure in air, thick and highly defective external and internal scales of mainly magnetite (Fe3O4) were formed in steam-containing atmospheres. Figure 19.1 compares the oxide scales formed on the P92 steel in air and in an argon-50vol% steam atmosphere after 10 000 h exposure at 650°C; the associated mass gains were 0.1 and 45 mg cm–2, respectively. The results shown in Fig. 19.1 are from specimens exposed in an argon50vol% water vapour atmosphere not in pure steam. In laboratory testing this is a convenient means of assessing the steam oxidation resistance of steels. Instead of a steam generator, argon is passed through a water bath at a constant temperature. By adjusting the temperature of the water bath, the concentration of water vapour in the argon can be controlled. After passing through the water bath, the gas is led through heated pipes to the furnace retort, to prevent the water condensing out. Comparison tests have shown 519 WPNL2204
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Ni layer to preserve scale during preparation
20 µm
200 µm
(a)
(b)
19.1 Micrographs for P92 steel of the oxide scale formed after 10 000 h at 600°C (a) in air and (b) in Ar-50 vol % steam (note difference in magnification). (a) P92 steel, 10 000h/650°C in air, mass gain 0.1 mg cm–2; (b) P92 steel, 10 000 h/650°C in Ar-50%H2O, mass gain 45 mg cm–2.
that there is no significant difference between testing in pure steam and in argon-50 vol% H2O.
19.2
Mechanisms of enhanced steam oxidation
19.2.1 Stability of oxides Figure 19.2 shows the thermodynamic stability of the oxides of iron and indicates which oxide, if any, is stable at a given temperature and oxygen partial pressure of the environment for an iron activity of 1. In air (oxygen partial pressure 0.2 bar) at 600°C, the stable oxide is haematite Fe2O3, which has a narrow homogeneity range with a low concentration of lattice defects (vacancies). Diffusion of ions through haematite is therefore relatively slow and as a result oxide growth rates are low. Oxide scales based on Fe2O3 will therefore tend to be protective as evidenced by the results of long-term exposure of the chromium steels in air (Fig. 19.1). Oxidation of iron by steam leads to the production of hydrogen, and the presence of hydrogen in the atmosphere above the oxide scale reduces the oxygen partial pressure very rapidly. Haematite formation can no longer be sustained and the next oxide, Fe3O4, magnetite, will form. This oxide has a wide homogeneity range, that is, it can accommodate excess iron and oxygen ions, maintaining electrical neutrality by an increasing concentration of lattice
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Mechanisms of oxidation and the influence of steam oxidation 900
800
700
600
Temperature (°C) 500
400
521 300
Log (oxygen partial pressure) (bar)
–5
–10 Fe3O4 magnetite –15
Fe2O3 haematite
–20 FeO wustite –25
–30
–35 8
10
12 14 Reciprocal temperature (K–1)
16
18
19.2 Thermodynamic stability of iron oxides at an iron activity of 1.
vacancies. This allows very rapid ionic diffusion and leads to high oxide growth rates. Fe3O4 scales, if they form, are therefore thick and contain a high concentration of voids and gaps.
19.2.2 Stages of steam oxidation Based on the results of short-term testing, Zurek (2004) proposed that the oxide scale formed on 9–12% chromium steels in steam-containing environments followed the sequence shown in Fig. 19.3. • •
At time t1, a thin, protective scale based on (Fe,Cr)2O3 or (Fe,Cr)3O4 forms. The Cr/Fe ratio is greater than about 0.25 (Rahmel and Tobolski, 1965). At time t2, the protective layer breaks down locally, perhaps owing to the depletion of chromium beneath the scale so that the Cr/Fe ratio decreases. Rapid growth of Fe3O4 then takes place and owing to the high outward diffusion rates of the iron ions, scale grows at the gas/oxide interface. The oxide surface is then not in equilibrium with the gas, the oxygen partial pressure at the oxide surface is lower than that in the gas and haematite cannot form. The rapid outward diffusion of ions to the scale/gas interface causes the formation of vacancies that condense into voids and gaps. Beneath the magnetite scale, the oxygen partial pressure is still sufficiently high for internal oxides FeO and Cr2O3 to be formed.
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t2
H2 H2
Protective spinel
H 2O
Alloy surface
Alloy surface
Fe3O4
Protective spinel
FeO + Cr2O3
Alloy
t3
Fe3O4 + Cr2O3
Alloy
t4
t5
Fe2O3
H 2O
H2 Fe3O4
Original alloy surface
Fe3O4
Fe3O4 + (Fe, Cr)3O4 FeO + Cr2O3 Fe3O4 + Cr2O3
Original alloy surface
Fe3O4
Alloy
FeO + Cr2O3 Alloy
Fe3O4 + (Fe, Cr)3O4 FeO + Cr2O3 Alloy
19.3 Stages in the steam oxidation of 9–12% chromium steels; the time scale varies according to the exact composition of the steel. WPNL2204
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•
• •
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At time t3, thickening of the inner scale leads to partial transformation of FeO to Fe3O4 and the Cr2O3 oxide particles become embedded in the scale. Further transformation leads to the formation of the spinel (Fe,Cr)3O4. Because Cr2O3 can only be formed internally, the place beneath which this oxide is observed coincides with the original alloy surface (Quadakkers et al., 2005). At time t4, the gap in the outer magnetite scale becomes more extensive and the area available for diffusion paths between the gas/oxide and oxide/alloy interfaces is increasingly restricted. At time t5, the gap in the outer scale is now so large that outward diffusion of iron ions practically stops. The activity of iron in the outermost scale will then decrease and haematite will be formed. If haematite is observed on the outermost surface of a steel oxidized in steam, this indicates a high concentration of macroscopic defects in the scale (voids and gaps) with associated poor scale adherence.
The actual times t1 – t5 vary widely depending on the steel composition, temperature, gas flow rates and other factors.
19.2.3 Development of defects and spalling of oxide scales The presence of defects such as voids and gaps in the oxide scales formed in steam- containing atmospheres will affect the adherence of the scales, especially during thermal cycling. The spalling (exfoliation) of scales may lead to enhanced oxidation rates as fresh material is exposed to the atmosphere. Figure 19.4 shows the scale microstructure on 10Cr–Mo–W exposed at 650°C in Ar-7% H2O for 5, 10, 50, 72 and 100 h. Formation of pores and eventually gaps in the individual stages of their evolution can be observed. After 5 h oxidation the scale is still mostly compact; owing to local vacancy condensation, some voids start to form on the original metal surface. After 10 h oxidation the voids on the original metal surface become larger and they begin to affect further oxide growth. The oxygen partial pressure in the gap probably stays constant and thus below the gap the gradient must become larger as the oxygen partial pressure at the scale/alloy interface equals the dissociation pressure of the oxide. Owing to the steeper oxygen partial pressure gradient, the transport of iron ions from the metal must increase. Thus ‘secondary’ growth of oxide scale in the gap becomes possible and the gap moves outward. Because of the enhanced iron ion outward diffusion, a secondary gap at the interface between the internal oxidation zone (FeO + Cr2O3) and the Fe, Cr spinel layer is formed. The appearance of haematite on the top of the scale is a result of the large gap in the scale. During the following hours of oxidation, the gradient of oxygen partial pressure remains
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5h
10 h
Compact oxide scale
Voids
Alloy
Pores
Alloy
50 µm
50 µm
(a)
(b)
50 h
72 h
Gap Gap Secondary gap Secondary gap Alloy
Alloy 50 µm
(c)
50 µm
(d)
100 h
Gap Nearly compact oxide scale
Secondary gap
Alloy 50 µm
(e)
19.4 Metallographic cross-sections of 10Cr–Mo–W after exposure for 5, 10, 50, 72 and 100 h at 650°C in Ar-7% H2O showing formation and evolution of the in-scale gap.
unchanged, the gap becomes larger and the external haematite layer starts to spall. The completely spalled haematite layer allows an excess of molecular oxygen to enter the external voids and/or gap and healing of gaps formed in previous stages of oxidation occurs. Figure 19.4 for 100 h clearly illustrates the situation described above, where on the left-hand side a nearly compact
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oxide scale has formed (complete healing of the gaps) while on the righthand side a very porous scale can be seen. The considerations presented above concern the formation and evolution of the ‘transient’ gap, which is formed after several minutes or hours of oxidation. The voids and gaps observed after a few thousand hours oxidation are the result of other factors such as cooling during exposure, cracking, steel microstructure and formation of Cr-rich stringers in the inner part of the oxide scale. The first voids, as already shown, start to form after very short oxidation times. The cracking and spallation of scales is correlated with the type, morphology and growth of pores and voids in the scale and could be influenced by the steel microstructure. The formation of pores and gaps is the dominating factor for the spalling characteristics of the oxide scales formed in steam. If isothermal exposure is continued, the gap may heal by growth of the inner oxide scale as gas molecules gain access through the outer layer. If isothermal exposure is continued for sufficiently long times, the gap may eventually completely close which results in a more or less compact layer. Transport of Fe cations to the surface becomes possible again and thus the outer haematite transforms into magnetite after longer times (a few hundred hours) at 650°C. If a thermal cycle is introduced in the early stages of the oxidation process, the poor adherence caused by the presence of the large gap results in spallation of the outer scale. This is illustrated in Fig. 19.5(a) for a 10CrMo–W steel exposed for 10 000 h at 625°C. Further exposure then results in growth of the freely exposed magnetite layer. Figure 19.5(b) shows the scale microstructure of the same steel after 10 000 h exposure at 650°C showing that spalling has occurred within the inward growing oxide scale, owing to the accumulation of voids there. The question of whether spallation or healing of the outer layer above the gap occurs depends on the time at which for a given temperature the thermal cycle is introduced. This can be seen from the results in Fig. 19.4 for the 10Cr–Mo–W–Si steel. If the specimen is cooled after 72 h exposure, spallation of the top layer occurs owing to the presence of the large in-scale gap. If the first temperature cycle is introduced after 250 h, sufficient time is available for (partial) healing of the gap and thus excellent oxide adherence is found even during exposure for several thousand hours.
19.3
Steam oxidation rates
19.3.1 Long-term exposure The long-term steam oxidation behaviour of the ferritic Cr steels can be rationalized on the basis of the schematic in Fig. 19.6. This graph should
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Initiation of spallation
Alloy
100 µm
(a)
Initiation of spallation
Alloy 100 µm
(b)
19.5 Different types of scale spallation during oxidation on 10Cr–Mo– W steel at (a) 625°C after 10 000 h and (b) 650°C after 10 000 h in Ar50% H2O.
only be considered as a rough approximation to the real behaviour because the borderlines between the various oxidation regimes can be strongly influenced by alloying additions, temperature and surface treatment. The three groups may be distinguished as follows: (I)
steels containing up to around 9–10%Cr form thick oxide scales, the main constituent of which is magnetite;
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Group III
log Kp
Group I
527
6
8
10 12 Cr content (mass%)
14
16
19.6 Schematic illustration showing qualitative dependence of oxidation rate on Cr content for ferritic/martensitic steels exposed in Ar-50% H2O at 550–650°C; designations ‘Group I – III’ are explained in the text.
(II) steels containing 10–12%Cr exhibit highly variable steam oxidation behaviour, the thickness and morphology of the oxide scales differing substantially as a function of test duration, temperature, minor alloying additions, grain size and surface treatment; (III) steels with more than about 12.5%Cr possess excellent steam oxidation resistance and the thin scales consist mainly of Cr2O3, (Cr,Fe)2O3 or Cr-rich (Cr,Mn,Fe)3O4 with an outermost layer of Fe2O3.
19.3.2 Influence of chromium content The commercial 12%Cr–1%Mo steel should, on the basis of Fig. 19.6, exhibit good steam oxidation resistance. However, the chromium content specified for this steel is 10–12.5%, which means that if the chromium content is at the low end of the specified range, the steam oxidation resistance will be similar to the low chromium steels. If the chromium content is at the high end of the specification, then good steam oxidation resistance would be expected. To ensure adequate steam oxidation resistance, it is therefore important that the minimum chromium content is above about 12%.
19.3.3 Effect of minor alloying additions Other minor alloying elements can significantly affect the steam oxidation resistance of the high chromium steels. Elements which promote the formation
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of spinel-type oxide scales or which enhance the diffusion rates may be beneficial. There is some evidence that small additions of silicon and manganese lead to improved steam oxidation resistance (Quadakkers and Ennis, 2002) and the steels that show the highest steam oxidation resistance, such as VM12 (Gabrel et al., 2006), contain 11–12% chromium and around 0.5% silicon.
19.3.4 Anomalous temperature dependence The formation of a protective oxide scale depends on the incorporation of chromium into the external scale and therefore the diffusion rate of chromium in the steel is an important factor. Because the diffusion rate of chromium increases with temperature, it is possible for the steam oxidation resistance to increase with increasing temperature and this has been observed by Zurek et al. (2004) for a number of chromium steels (Fig. 19.7). This effect should be taken into account in the development of steels with improved steam oxidation resistance, as testing at the highest application temperature may give a misleading assessment; the oxidation rate at lower temperatures could well be higher.
19.3.5 What is an acceptable oxidation rate? Components are generally designed on the basis of the strength needed to ensure that the planned service lifetime may be achieved. Oxidation effects 7
Relative mass change
6 10Cr–Mo–W
5 4 3
10Cr–Mo–W–Si
2 11Cr–M
o–Co
HCM12
1 12Cr–Mo–V 0 540
560
580
600 Temperature (°C)
620
640
19.7 Relative mass changes of different ferritic steels during exposure for 1000 h in Ar-50% H2O showing different types of temperature dependence of the oxidation rates.
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may lead to changes in the component dimensions and other properties and these changes can reduce the service life. Because oxidation is a surfacerelated effect and the strength properties are a bulk effect, whether a given oxidation rate will have a significant effect on the mechanical behaviour will depend on the initial component dimensions. Considering the material consumed by the oxidation process, the net section loss that can be tolerated for a component with a large surface to volume ratio, for instance a heat exchanger tube, will be much smaller than that of a component with a small surface to volume ratio, such as a turbine rotor. An indication of what is an acceptable steam oxidation rate for a given component may be obtained by looking at the long-term experience with low alloy steels in power plant at lower temperatures. The 1Cr1/2Mo steel has been successfully used at temperatures up to 550°C for durations of several hundred thousand hours. Figure 19.8 compares the mass changes for this steel at 550°C with the results obtained for P91 and P92 at 600 and 650°C. It is clear that the steam oxidation behaviour of the 9%Cr steels, with two to eight times higher mass gains in 10 000 h at 600–650°C than the 1Cr1/2Mo steel at 550°C will present a potential problem for the application of the steels in power plant. The implications of these high oxidation rates on the component service lifetime will now be considered.
50 P92 650°C
Mass change (mg cm–2)
40
30
P91 650°C P92 600°C
20 P91 600°C 10 1Cr0.5Mo 550°C 0 0
2000
4000
6000 8000 Test duration (h)
10000
12000
19.8 Mass change curves for P91 and P92 at 600 and 650°C, with comparison values for 1Cr1/2Mo at 550°C.
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Creep-resistant steels
Oxidation and service life
19.4.1 Effects of oxidation on service life For discussion of the steam oxidation interaction with the stress rupture behaviour, we shall take a mass gain of 30 mg cm–2 in 10 000 h as the benchmark (see Fig. 19.8) which is equivalent to a total scale thickness of 0.2 mm, of which about half is internal scale. Therefore a wall thickness reduction of 0.1 mm from each side in 10 000 h would be seen and if we extrapolate linearly, this would mean a reduction in wall thickness of 1.5 mm from each side in 15 years. A linear extrapolation can be justified as being conservative, because of the likelihood of scale spallation discussed above. The possible effects of the oxidation process on the mechanical behaviour are: (1)
(2)
(3)
Reduction of load-bearing cross-section: The thick oxide scales formed at 600–650°C in steam on the 9–10%Cr steels lead to a reduction in the wall thickness, which will in turn lead to increased stress (the load remains constant but the cross-sectional area is reduced). The increase in stress and the corresponding reduction in rupture life will, however, depend on the initial wall thickness of the component. Thermal insulation effect of thick oxide scales: The thick external and internal oxide scales formed on the 9–10%Cr steels in steam will have a considerable influence on heat transfer across the tube wall. In addition to reducing the efficiency of the steam generation, the decrease in the heat transfer caused by thermally insulating, thick oxide scales could lead to overheating of the tube material, unless, of course, the scales spall. Spalling of oxide scales: The spalling of the thick oxide scales would be beneficial in terms of reducing the above-mentioned thermal insulation effect. However, the spalled oxide itself can lead to tube overheating if it becomes entrapped in the system, thus reducing the flow rates inside the tubes. Erosion of turbine components may also occur. The defect nature of the thick oxide scales formed in steam will promote spalling at some stage and therefore it seems prudent to design systems so that the risk of blockages caused by spalled oxide flakes can be reduced as far as possible.
19.4.2 Quantitative estimation of the effect of steam oxidation on service life For P92, the stress rupture curves for temperatures in the range 600–700°C are shown in Fig. 19.9. From these curves it can be seen that an increase in
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300 Mean stress rupture strength of P92 250 Temperature (°C)
Stress (MPa)
200
550 150 600 612 625 637 650
100
50 700 0 1000
10 000
100 000 Time to rupture (h)
1000 000
19.9 Mean stress rupture curves for P92.
the stress from 100 to 120 MPa at 600°C results in a decrease in rupture life from 170 000 to 75 000 h, representing a life reduction of about 40%. Figure 19.9 also shows that for a given stress level, say 100 MPa, increasing the temperature from 612 to 625°C reduces the rupture life from 80 000 to 33 000 h, which is equivalent to a 65% reduction in life or a three-fold increase in the secondary creep rate for a 10 K rise in temperature. Using the steam oxidation rate of 0.1 mm steel thickness loss in 10 000 h mentioned above, the service life of components can be estimated and compared. Figure 19.10 shows the rupture life reductions for a thick-walled pipe and a thin-walled superheater tube at 600°C. The thick-walled pipe can tolerate a thickness loss of 1.5 mm, which could occur after 15 years exposure at 600°C, with only a slight reduction in service life. The thin-walled tube with the same steam oxidation rate would experience a considerable reduction in service life from over 200 000 h to 60 000 h. Laboratory tests have shown that the total scale thickness on 9–10%Cr steels exposed for 10 000 h at 600–650°C will be around 0.2 mm, which could lead to a temperature increase of 50 K and a very large associated reduction in rupture life. The thermal insulation effect of the thick oxide scales would then become the most damaging effect with respect to the component service life. Some confirmation of this has been demonstrated in field trials; the results of tests in a by-pass loop of a power station are described and the creep failure of an E911 pipe was attributed to overheating
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1000 000
Rupture life (h)
40 mm diameter, 6 mm wall thickness tube
100 000
300 mm diameter, 40 mm wall thickness pipe
P92 steel life reduction due to material loss by oxidation 600°C, pressure 300 bar 10 000 0.0
0.2
0.4 0.6 0.8 1.0 1.2 Material loss from each side (mm)
1.4
1.6
19.10 Rupture life reduction for two P92 components, a pipe of 300 mm diameter and wall thickness 40 mm and a tube of 40 mm diameter, wall thickness 6 mm.
produced by the thick oxide scale that had formed (Zabelt and Wachter, 1995). The temperature increase was estimated to have been from 630 to 675°C, which correlated roughly with the observed time to creep failure of the tube.
19.5
Development of steam oxidation-resistant steels
19.5.1 Steel composition It has been demonstrated that the steam oxidation resistance of chromium steels can be improved by increasing the chromium content to 12%. To ensure a fully martensitic microstructure, the increased chromium content has to be balanced by addition of elements that stabilize the austenite phase without reducing the ferrite/austenite transformation temperature. Cobalt and copper have been the favoured additions and the newer 12% chromium steels usually contain one of these elements. Although the steam oxidation resistance of these steels is much better than that of the 9% chromium steels and in short-term tests the stress rupture strength is at least as good, the stress rupture strength falls rapidly and falls below that of the 9% chromium steels after long test durations (Gabrel et al., 2006). This has been attributed to the formation of the Z phase, CrNbN, which in turn causes the dissolution of the fine niobium and vanadium nitrides
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and carbonitrides that provide the long-term strength. It is therefore concluded that, based on current knowledge, it is not possible to obtain the required high stress rupture strength and at the same time good steam oxidation resistance in a single steel composition. Attention has therefore been turned to coatings.
19.5.2 Coatings The status of coatings development has been summarized by Aguero (2006). The aim has been to enrich the surface layer of the steel with either aluminium or chromium, so that alumina or chromia oxide scales form. Silicon has also been considered, but as silica reacts with steam to form volatile silicon hydrides, such coatings are not suitable. There are a number of constraints concerning the coating techniques employed. The temperatures used for the coating processes should allow the basic microstructure of the steel to be fully martensitic with the fine precipitates of niobium and vanadium nitrides and carbonitrides that are essential for the long-term creep strength. The processes investigated are principally diffusion coatings and promising results have been obtained, especially for aluminium enrichment of the surface. The up-scaling of the coating processes for real components still requires investigation.
19.5.3 Surface modifications It is well known that modification of the surface of austenitic stainless steels by shot peening improves the oxidation resistance by enhancing the diffusion of chromium into the scale. Investigations of various mechanical treatments, including shot peening, grinding and polishing, on the steam oxidation resistance of martensitic steels has not revealed any significant improvement in the steam oxidation behaviour. This may be due to the fact that the initial tempered martensite microstructure consisting of fine martensite laths and the high dislocation density arising from the martensite transformation during cooling already provide the best achievable conditions for diffusion of chromium, which are however not sufficient for the formation of a protective oxide scale.
19.6
Outlook
From the point of view of the creep strength, the 9% chromium steels are suitable for applications at metal temperatures up to around 620°C and further developments are in progress to enable temperatures of 650°C to be achieved. However, the steam oxidation resistance remains a significant problem. Increasing the chromium content to 12% or more leads to excellent steam
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oxidation resistance but at the cost of reduced creep strength. At the present time and in the light of the available data, it does not appear possible to obtain both steam oxidation resistance and sufficient creep strength in a single steel composition. Coatings offer a potential solution and aluminizing has been successfully applied on a laboratory scale. Another line of development is to move to a fully ferritic steel with a chromium content of, say, 14% or more. Because of the extremely low solubility of carbon and nitrogen in ferrite, it will not be possible to produce a dispersion of strengthening carbide and nitride precipitates by conventional steel processing techniques. Instead, intermetallic phases, such as the Laves phases, could be considered for creep strengthening, provided that such precipitates are sufficiently fine and resistant to coarsening at the application temperatures.
19.7
Sources of further information
Kofstad P, High Temperature Corrosion, Elsevier Applied Science, London and New York 1988. Schuetze M, Protective Oxides and Their Breakdown, Wiley Series on Corrosion and Protection, Holmes D R (series ed.), J Wiley & Sons, Chichester, UK, 1997.
19.8
References
Aguero A (2006). ‘Coatings for protection of high temperature new generation steam power plant components; a review’, in Materials for Advanced Power Engineering 2006, Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Energy Technology Volume 53, Part II, Forschungszentrum Juelich, Germany, 949–964. Gabrel J, Bendick W, Vandenberghe B and Lefebvre B (2006). ‘Status of development of the VM12 steel for tubular applications in advanced power plants’, in Materials for Advanced Power Engineering 2006, Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Energy Technology Volume 53, Part II, Forschungszentrum Juelich, Germany, 1065–1076. Quadakkers W J and Ennis P J (2002). ‘The oxidation behaviour of chromium steels in supercritical steam power plant’, in Materials for Advanced Power Engineering 2002, Lecomte Beckers J, Carton M, Schubert F and Ennis P J, (eds), Energy Technology Series, Volume 21, Part II, Forschungszentrum Jülich, Germany, 1131–1142. Quadakkers W J, Ennis P J, Zurek J and Michalik M (2005). ‘Steam oxidation of ferritic steels: laboratory test kinetic data’, Materials at High Temperatures, 2005, 22 (1/2), 37–47. Rahmel J and Tobolski J (1965). ‘Einfluss von Wasserdampf und Kohlendioxid auf die Oxidation von Eisen in Sauerstoff bei hohen Temperaturen’, Corrosion Science, 1965, 5, 333. Zabelt K and Wachter O (1995). Ergebnisse von Feldversuchen in Kraftwerken mit 9 bis 12-%-Chromstählen, 18 meeting of the Arbeitsgemeinschaft für warmfeste Stähle,
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Presentation 6. 1, Verein Deutscher Eisenhuettenleute (VDEh), Düsseldorf December 1995. Zurek, J (2004). Oxidation and Oxidation Protection of Ferrritic and Austenitic Steels in Simulated Steam Environments at Temperatures Between 550 and 650°C, Doctoral Thesis, Rheinisch-Westfälische Hochschule, Aachen, September 2004. Zurek J, Wessel E, Niewolak L, Schmitz F, Kern T U, Singheiser L and Quadakkers W J (2004). ‘Anomalous temperature dependence of oxidation kinetics during steam oxidation of ferritic steels in the temperature range 550 – 650°C’, Corrosion Science, 2004, 46/ 9, 2301–2317.
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20 Alloy design philosophy of creep-resistant steels M . I G A R A S H I, Sumitomo Metal Industries, Japan
20.1
Introduction
High Cr ferritic steels such as ASME P91 steel have successfully been used for large diameter and thick section boiler components such as main steam pipes and headers in supercritical pressure boilers in fossil-fired power plants.1 Recent trends towards the utilization of clean energy leading to protection of the global environment have been accelerating the application of ultra supercritical pressure (USC) boilers, which are operated with higher efficiency in power generation than conventional boiler and thus release less carbon dioxide.2,3 The USC boiler requires new steels with improved creep rupture strength and steam oxidation resistance at elevated temperatures over 600°C, because of the increase in the operating temperature and pressure of the steam. Addition of W to the ferritic steel has been found to be effective in increasing creep rupture strength at high temperatures and has been already used in some developed steels such as T92/P92 and T122/P122 for the USC boilers.4 High strength austenitic stainless steels such as TP347HFG, SUPER304H and HR3C have been used extensively as superheater and reheater tubes for the latest USC boilers all over the world.5–8 In this chapter, the alloy design philosophies for creep-resistant ferritic and austenitic stainless steels for various components in USC power plants are reviewed and demonstrated in detail and the future research will be discussed with regard to further improvement of the steels for application in 700°C A-USC plants.9
20.2
Creep-resistant steels for particular components in power plants and the properties required
20.2.1 Water wall Figure 20.1 shows a schematic illustration and photographs of various boiler components such as the water wall, superheater, reheater, header and main 539 WPNL2204
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Creep-resistant steels Header 9 ~ 12% Cr ferritic steels
Superheater 18 ~ 25%Cr austenitic steels
Main steam pipe 9 ~ 12%Cr ferritic steels
Main steam pipe am Ste er t Wa
Water wall tubes Header Reheater Seawater
0.5 ~ 2% Cr steels
Water
Furnace
20.1 Schematic illustration and photographs of a fossil fired boiler and typical materials.
steam pipe and their typical materials used in a recent fossil fired power plant.10 The water wall tubes are welded in a panel structure to achieve effective heat exchange in order to produce pressurized steam from water under the super critical conditions in a furnace. Conventional steels and CrMo steels are used for the water wall tubes according to the operating temperatures and pressures, which are usually heated up to 500°C in the latest USC plants.11 The materials requirements for the water wall are, therefore, strength and corrosion resistance at elevated temperatures as well as weldability and formability to construct the water wall panel. Recently, steels with higher strength and good weldability have been developed and successfully used in the latest USC plants described below.12,13
20.2.2 Superheater/reheater The steams are superheated in a superheater (SH) up to the highest pressure at a designed temperature and reheated in reheater (RH) up to the highest temperature that will achieve the designed thermal efficiency, for instance, 42% at 600°C and 25 MPa with a SH and 605°C and 4.2 MPa with a RH in one of the latest USC plants.11 The materials requirements for SH and RH are, therefore, most severe for the components and high strength austenitic stainless steels are used for this purpose. They are required to have high creep strength
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at a designed temperature and above and to have both steam oxidation resistance for the inner surface and hot-corrosion resistance for the outer surface against coal ash containing sulfur, chlorine, vanadium and other corrosive salt-forming elements. Various types of austenitic stainless steels have been developed for SH/RH tubing application, which are described in detail later.
20.2.3 Header and main steam pipe The superheated steams are gathered into a header pipe and transferred to a turbine system through main steam piping. The header and main steam piping should, therefore, be a heavy wall thickness pipe with a large diameter to keep huge amounts of steam pressurized. They are required to have high creep strength with good ductility and toughness to prevent catastrophic failure during operation and hydro-pressure testing. They are also required to have good thermal fatigue resistance to the thermal stress imposed, because fossil fired plants are usually used to adjust the electricity supply during the day, which means they are used in a daily start and stop operation. For this reason, high Cr ferritic steels with high creep strength have been developed and have already been used in the latest USC plants, because although ferritic steels are in principle inferior in strength to the austenitic stainless steels because they have a large diffusion constant at elevated temperatures caused by the difference in their crystal structures, they are superior in thermal expansion and conductivity. High Cr ferritic steels have been developed and successfully used in the latest USC plants and will be described in detail later.
20.3
Alloy design philosophies of creep-resistant steels
20.3.1 High strength low-Cr steels Figure 20.2 shows a tree chart of developed ferritic steels.10 Strengthening of ferritic steels is mainly achieved by dislocation strengthening using C and N for martensitic and bainitic transformation and the resultant microstructure, solid solution hardening by elements such as Mo and W, and precipitation hardening by carbo-nitrides containing Cr, V and Nb and Cu phase. Corrosion resistance of the steels is mainly achieved using Cr and Si. These ferritic steels are used for various components in USC plants according to their respective strength and corrosion resistance. Figure 20.3 shows the alloy design philosophy of 2.25Cr–1.6W–V–Nb steel (HCM2S; KA-STBA24J1, T23/P23, ASME CC2199).14,15 2.25Cr–1.6W– V–Nb steel is used as a water wall, in SH and RH tubes, and in the header and main steam pipe in fossil fired boilers and heat recovery boilers. The steel has been developed to improve the creep rupture strength of 2.25Cr– 1Mo steel at elevated temperatures mainly by substituting Mo by W. High
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12 Cr steels
9Cr–1Mo–V–Nb (KA–STBA28, T91)
Strength
2.25Cr–1.6W–V–Nb (KA–STBA24J1, T23, ASME CC2199)
+ V, Nb
+ W, V, Nb, B
STB510 + Mn KA-STB480 STB410 STB340 +C
STBA 13 STBA 12
+ Cr, Mo
STBA 24 STBA 23 STBA 22 KA-STBA21 STBA20
9Cr–2Mo (KA–STAB27)
+ W, Cu, B 12Cr–1Mo–1W–V–Nb (KA–SUS410J2TB) –C, + W, Nb X20CrMoV121
– C + Mo + Cr
STBA 26 STBA 25
+ C, V + Cr
+ Mo +Cr, Ni, Cu
KA-indicates that the steel is designated by the METI standard
KA-STBA10 (CR1A)
Fe
Corrosion resistance
20.2 Tree chart of developed ferritic boiler steels.
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11Cr–0.4Mo–2W–Cu–V–Nb (KA–SUS410J3TB, T122, ASME CC2180)
542
C-steel ~ CrMo steels (2.25Cr)
Alloy design philosophy of creep-resistant steels Weldability
Toughness
max.Hv350
Hardenability : B addition
Low C-0.06 mass%
Fully tempered bainitic structure
543
Creep strength Solution strengthening : High W
Precipitation strengthening : V, Nb, B
No preheating and PWHT after welding 0.06C-2.25Cr-1.6W-0.1Mo-0.25V-0.05Nb-B Matching welding filler used without PWHT
20.3 Alloying philosophy of 2.25Cr–1.6W–V–Nb steel (HCM2S; KASTBA24J1, T23/P23, ASME CC2199).
strength is mainly achieved by the combination of a solid solution of W with (V, Nb)C dispersion hardening in a fully tempered bainitic matrix. Addition of B enhances the bainitic microstructure and is found to improve the toughness of the steel to a great extent. Low C content has been chosen to improve the weldability of the steel and as a result no preheating and post weld heat treatment (PWHT) is potentially required for some applications. Figure 20.4 shows the weldability and the crack sensitivity of 2.25Cr–1.6W–V–Nb steel. It is seen that 2.25Cr–1.6W–V–Nb steel exhibits superior weldability even without preheating, while both T22 and T91 grade steels require preheating above 200°C and 150°C, respectively. Figure 20.5 shows creep rupture data for 2.25Cr–1.6W–V–Nb steel pipes.16 It can be seen that the creep strength of this steel is estimated to be about 1.8 times higher than those of T22/P22 steel. The longest creep rupture time of 2.25Cr–1.6W–V–Nb steel pipes is about 90 000 h at 550°C. Their long-term creep strengths are very stable at temperatures between 500 and 600°C. It is, however, noted that above 600°C the longer term rupture time tends to decline over the averaged curve, which has been found to show the effect of oxidation when a small specimen was used and therefore no significant strength degradation took place, like that observed for T22 grade steels.16 Figure 20.6 shows a calculated phase diagram for 2.25Cr–1.6W–V–Nb steel at 600°C with changing Cr and C contents.17 The equilibrium phase diagram suggests that the final microstructure of 2.25Cr–1.6W–V–Nb steel consists of ferrite (α) + MX ((V,Nb)C) + M6C and may include a small amount of M23C6. Figure 20.7 shows TEM micrographs of extraction replicas of 2.25Cr– 1.6W–V–Nb steel (a) normalized and tempered and (b) crept for 12 567.6 h at 600°C. In the tempered specimen, M23C6 is formed along prior austenitic
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Cracking ratio (%)
80 70
Cracking ratio C
T22
C = h × 100 (%) H
60
T91
50 H
40
h
30 20 2. 25Cr-1.6W-V-Nb
10 0 0
50
100
150 200 250 Preheating temperture (°C)
300
350
20.4 Weldability of 2.25Cr–1.6W–V–Nb steel (HCM2S; KA-STBA24J1, T23/P23, ASME CC2199). 500 400
Stress (MPa)
300
200 500°C 550°C 100 90 80 70 60
600°C
50 40 101
650°C 10
2
3
10 Rupture time (h)
10
4
10
5
20.5 Creep rupture data for 2.25Cr–1.6W–V–Nb steel pipes (HCM2S; KA-STPA24J1, P23, ASME CC2199).
grain boundaries and MX is formed along lath boundaries and in the bainitic matrix. M7C3 is occasionally observed along lath boundaries. In the crept specimen, on the other hand, no M23C6 and M7C3 are observed and instead blocky M6C is formed along the prior austenitic grain boundaries and fine MX remains along lath boundaries and in the bainitic matrix as shown in Figure 20.7(b). It is noted that the pronounced lath structure is kept even
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0.20 0.18
C-content (mass%)
0.16 0.14 α + M23C6 + M6C + MX
0.12 0.10 0.08 0.06 0.04
α + M X + M 6C
0.02 0 0
1
2 3 Cr-content (mass%)
4
5
20.6 Calculated phase diagram for 2.25Cr–1.6W–V–Nb steel at 600°C.
MX M 6C
M6C 1 µm
1 µm (a)
M23C6
MX
MX
1 µm
M7C3 1 µm
(b)
20.7 TEM micrographs of MX in 2.25Cr–1.6W–V–Nb steel (a) normalized and tempered, (b) crept for 12 567.6 h at 600°C.
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after long-term creep deformation. Formation of M6C means that the steel loses W and/or Mo in solution, which is the key element for solid solution strengthening. However, W is superior to Mo in retarding the formation of M6C during long-term creep deformation. Figure 20.8 shows the change in precipitated W in 2.25Cr–1.6W–0.2Mo–V–Nb steel and Mo in 2.25Cr–1.0Mo– V–Nb steel during aging for up to 10 000 h at temperatures between 550°C and 650°C.18 The lines indicate fitted curves, assuming that the growth rate is controlled by the Johnson–Mehrl–Abrami theory. It is seen that in 2.25Cr– 1.6W–0.2Mo–V–Nb steel the growth rate of M6C is 10 to 100 times slower than that in 2.25Cr–1.0Mo–V–Nb steel. This is one of the major reasons for the high creep strength of 2.25Cr–1.6W–V–Nb steel.
20.3.2 Martensitic high-Cr steels for heavy-wall thickness piping
1
Fraction of Mo-precipitated (–)
Fraction of W-precipitated (–)
11Cr–0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/ P122, ASME CC2180) and KA-SUS410J3DTB) has been developed to improve the creep rupture strength and corrosion resistance of P91 type 9%Cr steels above 600°C, mainly achieved by higher Cr content and substitution of part of the Mo by W.19,20 In order to suppress δ-ferrite formation for thick wall pipes, Cu addition is chosen from the γ-forming elements shown in Fig. 20.9. Cu, unlike Ni and Mn, is a γ-forming element which does not reduce the Ac1 temperature much and does not enhance coarsening of M23C6 carbide. Cu addition enables the combination of higher Cr content with high W and Mo contents to be achieved. Figure 20.10 gives a summary of the alloy design philosophy of 11Cr– 0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/P122, ASME CC2180) and KA-SUS410J3DTB).19 This steel was originally developed
Experiment value 0.8 0.6 650°C
600°C
0.4 0.2
550°C
0 1
10
100 1000 104 Aging time (h) (a)
1 Experiment value 0.8 650°C 600°C
0.6 0.4
550°C 0.2
105
0 1
10
100 1000 104 Aging time (h) (b)
105
20.8 Changes in (a) precipitated W in 2.25Cr–1.6W–0.2Mo–V–Nb steel and (b) Mo in 2.25Cr–1.0Mo–V–N steel during aging for up to 10 000 h at temperatures between 550°C and 650°C.
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α-forming element
Base; 0.1C–11Cr steel Addition (mass%)
2W
∆Ac1 (°C)
1Cr 1Mo 0
Creq = Cr+6Si+4Mo+1.5W + 11V+5Nb+8Ti+12Al – 40C-30N-4Ni-2Mn – Cu-2Co (mass %)
2Cu 0.5Ni 0.05N
–40
γ-forming element –4
–2
0 ∆cr eq. (mass%)
+2
+4
20.9 Comparison of alloying elements with respect to changes in Ac1 temperature and Creq for 0.1C–11Cr model steels.
for two different purposes: one for large diameter, heavy wall thickness pipe and the other for a tubing application which often requires good corrosion resistance.19 The steel, therefore, has two different Cr versions: a steel with Cr content lower than 11.5% having good toughness and one with Cr content higher than 11.5% having good corrosion resistance. The steel with lower Cr content has an α′ single martensitic matrix microstructure, which is strengthened by the combination of MX carbonitrides with M23C6 along the prior austenitic grain boundaries. The steel with a higher Cr content has a δ-ferrite and α′ martensite dual matrix microstructure. This difference in the microstructure imposes variation in the creep strength shown in Fig. 20.11.21 Figure 20.11 shows creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steels with two different Cr content levels, the α′ steels with Cr less than 11.5% and the the α′ + δ steels with Cr equal to and more than 11.5%. It is obvious from these updated rupture data that the creep rupture strength should be analyzed separately for this steel according to the Cr content levels. Figure 20.12 shows the stress dependence of the minimum creep rate and the off-set strain on the acceleration creep of 11Cr–0.4Mo–2W–Cu–V–Nb steels with two different Cr content levels, the α′ and α′ + δ steels shown in Fig. 20.11.22 It can be seen that the minimum creep rate of both steels decreases with decreasing applied stress. The α′ steel exhibits a smooth change in the minimum creep rate with stress. It is found that in the α′ steel the power-law creep is relevant at a higher stress regime above 60 MPa where the stress exponents are high enough at between 5 and 15, while the viscous creep is at lower stress where the stress exponent is as low as 1. This behavior is very similar to that observed in P91 steel by Kloc and Sklenicka.23,24 The α′ + δ steel, on the other hand, exhibits a deviation from the expected
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Superior to 9%Cr VN, stable & effective Nb(C, N), grain refinement
Creep strength
W, stable
10.0~12.5%Cr V, Nb and N addition High W & Mo B addition
Stable long-term strength Weldability KA-SUS410J3J3TB/TP
KA-SUS410J3DTB
Low Ni Low C
Toughness
δ-ferrite5% (single phase)
High Cr 12%
δ-ferrite30% (dual phase)
Creq 9% Cu addition High temperature tempering 770°C
0.1C-10.5/12Cr-2W-0.4Mo-Cu-0.2V-0.05Nb-B-N
20.10 Alloying philosophy of 11Cr–0.4Mo–2W–Cu–V–Nb steel (HCM12A; KA-SUS410J3TB/TP (T122/P122, ASME CC2180) and KA-SUS410J3DTB).
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Corrosion resistance
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549
500 400 300 550°C
Stress (MPa)
200 600°C 100 90 80 70 60
650°C
50 40 30 101
Open symbols; Cr<11.5% Solid symbols; Cr≥11.5% 102
700°C 103 Rupture time (h)
104
105
20.11 Creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steels with two different Cr content levels, the α’ steels with Cr <11.5% and the α’ + δ steels with Cr ≥ 11.5%.
smooth change in the minimum creep rate around the critical stress, where the rupture life of the α′ + δ steel starts to deviate from the straight line. At lower stress a sudden drop in the minimum creep rate has been observed, indicating that the creep deformation mechanism has changed explicitly from the high stress regime and viscous creep becomes dominant in the lower stress regime. Figure 20.12(b) shows that the change in the off-set strain of the α′ + δ steel with stress is completely different from that of the α′ steel, in particular, at lower stress where a sudden drop in the minimum creep rate has been observed. This suggests that at a particular stress level in the α′ + δ steel the heterogeneous creep deformation takes place as a consequence of the prohibition of ordinary homogeneous deformation and the enhancement of localized deformation which gives rise to acceleration creep even at a very small strain and a very small creep rate. Figure 20.13 shows updated creep rupture data for 11Cr–0.4Mo–2W– Cu–V–Nb steel pipes with Cr < 11.5% (KA–SUS410J3TP, equivalent to P122, ASME CC2180).21 Three different fitted curves are based on the LMP method with original data and updated data, and the one determined by the region splitting method of Kimura25 adopted by the SHC (SHC is the committee for establishing allowable stress value, and soon, of high Cr ferritic steels for the ministry of economy, Trade and Industry, set up in Japan since 2004). Using the region splitting method, the allowable tensile stresses obtained are much lower than with those originally proposed, as given in Table 20.1.
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10–4
Creep rate (h –1)
1
10–5 α’ + δ
5 10–6 α’ 10–7 20
1 30
40
50 60 70 80 90 100 Stress (MPa) (a)
200
10–1
Creep rate (h –1)
550
10–2 α’ α’ + δ
10–3 20
30
40
50 60 70 80 90 100 Stress (MPa) (a)
200
20.12 Stress dependence of (a) the minimum creep rate and (b) the on-set strain on the acceleration creep of 11Cr–0.4Mo–2W–Cu–V–Nb steels with two different Cr content levels, the α’ and α’ + δ steels.
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500
LMP method with original data LMP method with updated data Region splitting method by SHC
400 300
550°C
Stress (MPa)
200 600°C 100 90 80 70 60 50
650°C
Open symbols; original data Solid symbols; updated data
40 700°C 30 101
102
103 Rupture time (h)
104
105
20.13 Creep rupture data for 11Cr–0.4Mo–2W–Cu–V–Nb steel pipes with Cr<11.5% (KA-SUS410J3TP, equivalent to P122, ASME CC2180). Table 20.1 Allowable tensile stress values for 11Cr–0.4Mo–2W–Cu–V–Nb steel revised according to the region splitting method adopted by the SHC in 2005. There are two steels designated according to different Cr content in the corresponding METI specification Temperature, (°C)
575
600
625
650
ASME CC2180-2(P122)* KA-SUS410J3TP pipe with 10.5/11.5 Cr KA-SUS410J3DTB tube with 11.5/12.5 Cr
107 100 94
83 68 52
61 46 25
(45) 27 16
* Allowable stress of ASME P122/T122 has been reassessed by an ASME subcommittee and will soon be revised in Japan.
To achieve longer term creep strength in advanced ferritic steels at elevated temperatures over 600°C, the creep deformation mechanism of the steels and the role of each precipitate in the respective creep deformation process have been examined using model steels with different initial microstructures consisting of an α′ single martensite matrix with M23C6 and with M23C6 and MX. Figures 20.14 and 20.15 are TEM micrographs showing the microstructural evolution during creep deformation of the α′ + M23C6 and the α′ + M23C6 + MX structures, crept and interrupted at 650°C with stresses of 70 and 100 MPa, respectively. It is seen that, in the α′ + M23C6 steel shown in Fig. 20.14, the excess dislocations are much reduced inside lath grains in the transient creep region. Around the minimum creep rate the migrations of lath and block boundaries seem to start forming equi-axed
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1µm (a) As NT
(b) After 50h, ε = 0.6%
(c) After 200h, ε = 1.35%
(d) After 872h, ε = 9.5%
20.14 TEM micrographs showing microstructural evolution during creep deformation of α’ + M23C6 steel, crept and interrupted at 650°C with 70 MPa. (a) As normalized and tempered (NT); (b) after 50 h, creep strain ε = 0.6%; (c) after 200 h, ε = 1.5%; (d) after 872 h, ε = 9.5%.
1µm (a) As NT
(b) After 216h, ε = 0.3%
(c) After 1620h, ε = 1.7%
(d) After 516h, ε = 5.0%
20.15 TEM micrographs showing microstructural evolution during creep deformation of the α’+M23C6+MX steel, crept and interrupted at 650°C with 100 MPa. (a) As NT; (b) after 216 h, ε = 0.3%; (c) after 1620 h, ε = 1.7% (d) after 5164 h, ε = 5.0%.
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Creep rate
subgrains. In the acceleration creep region, subgrain formation propagates to the whole specimen. In the α′ + M23C6 + MX steel shown in Fig. 20.15, the microstructural change during the creep deformation is very similar to that in the α′ + M23C6 steel but differs in the time taken to reach the same deformation microstructure. This means that MX carbonitrides inside lath grains serve as an effective obstacle against dislocation in motion and hence delay the rearrangement of dislocations and formation of subgrains. It is, however, noted that this seems to enhance the heterogeneous creep deformation along boundaries like the prior austenite grain boundary and packet, block and lath boundaries, since the greatly strengthened matrix by MX does not deform easily at low stress levels, although the regions around the boundaries which are softer than the matrix could deform much more easily. Figure 20.16 shows a schematic representation of creep rate versus time curves of the ferritic steels and the corresponding guiding principles for achieving long-term creep strength, which is derived from microstructural evolution during creep deformation.26 To achieve a higher creep strength the creep rate in the transient creep region can be reduced using fine dispersion of MX, α″, Cu-phase and also M23C6 and the Laves phase. All the fine precipitates serve as obstacles to dislocation in motion and hence delay the rearrangement of dislocations and formation of subgrains. In the acceleration creep region, however, subgrain formation proceeds to a great extent and homogeneous deformation becomes more difficult. In such a deformation scheme, heterogeneous creep deformation
(1) Reducing mobility of lath-, block-boundaries (2) Suppressing heterogeneous deformation
Reducing disl. mobility
Stabilization of martensite/ bainite (C-free, Co)
Solid solution (Mo, W etc) Dispersion and stabilization of boundary precipitates; M23C6, Laves (optimization of C, N, B)
Fine dispersion; MX, α”, Cu, M23C6, Laves
Something else Time
20.16 Schematic illustration of creep rate versus time curves representing creep strengthening mechanisms for the ferritic steels.
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along the prior austenite grain boundary becomes significant, sometimes resulting in creep rupture in a short period. To increase the creep resistance in the acceleration creep region, stabilization of lath, block and packet boundaries of martensite and of the prior austenite grain boundary is useful. This can be done by stabilization of the precipitates such as M23C6 and the Laves phase along these boundaries.
20.3.3 Austenitic stainless steels for superheater/reheater tubing Figure 20.17 shows creep curves of Type304H austenitic stainless steel crept at 650°C with stresses of 147, 118 and 98 MPa, compared with those of a 10–1
0.8 147MPa
0.7
147MPa
10–2
0.6 118MPa
0.4 0.3 α′+M23C6+MX 98MPa
0.2
98MPa
0.0
10–4 10–5 98MPa α′+M23C6+MX 98MPa
–6
10
0.1
10–7 0
2000 4000 6000 8000 1000 12000 Time (h) (a)
0
0.05
0.1 Strain (b)
0.15
0.2
10–3
Minimum creep rate (h–1)
Strain
0.5
Creep rate (1/h)
118MPa 10–3
10–4
304H 13 0.12C–0.002N 13
10–5
0.08C–0.05N 4
10–6
10–7 70
80
90 100 120 Stress (MPa) (c)
150
20.17 Creep curves for Type304H austenitic steel crept at 650°C with stresses of 147 118 and 98 MPa. (a) Creep rate against time, (b) creep rate against creep strain, (c) minimum creep rate against stress.
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9Cr ferritic steel.10 It is seen, unlike for the ferritic steels, that in Type304H steel the creep rate decreases gradually in the transient creep region but hardly increases in the acceleration creep region, giving a large off-set strain at the minimum creep rate. This means that homogeneous creep deformation takes place in Type304H steel. Figure 20.18 are TEM micrographs showing microstructural evolution during creep deformation of Type304H steel crept and interrupted at 650°C at 147 MPa.27 It is obvious that in the austenitic steel the initial dislocations density is in principle negligible and substantial numbers of dislocation are introduced homogeneously all over the grains during creep deformation in the transient creep region so as to suppress localized hardening. Once the dislocation starts to rearrange into a subgrain structure in some portions, the transition to acceleration creep takes place, as observed in the ferritic steels. This subgrain structure propagates to the other portions in the acceleration creep region. Note that the dislocation density does not change very much even at the end of the acceleration creep. This gives rise to a homogeneous distribution of dislocations and makes it difficult to achieve the deformation mode easily at a small strain as in the ferritic steels. This is confirmed by the specimens crept at low stresses. At lower stress levels, M23C6 forms along the dislocations inside the grain and serves as an effective obstacle to dislocation in motion and hence delays the rearrangement of dislocations and formation of subgrains as shown in Fig. 20.19.27
1µm (a)
(c)
(b)
(d)
(e)
20.18 TEM micrographs of Type304H steel crept and interrupted at 650°C and 147 MPa. (a) As ST; (b) after 50 h, ε = 2.6%; (c) after 100 h, ε = 4.8%; (d) after 246 h, ε = 13.6%; (e) after 429 h, ε = 20.0%.
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200nm
20.19 TEM micrograph of Type 304H steel crept at 650°C and 98 MPa and ruptured after 10 558 h.
Figure 20.20 shows the creep deformation mechanism and guiding principles for further strengthening of the austenitic alloys proposed.10 According to the process of microstructural evolution during creep deformation described above, all the fine precipitates such as MX, NbCrN, M23C6 , Cu, α-Cr, γ ′ and the Laves phase can serve as obstacles to dislocation in motion in the transient creep region and hence delay the rearrangement of dislocations and formation of subgrains. In the acceleration creep region, subgrain formation propagates everywhere until an equilibrium subgrain microstructure has been achieved for an easy deformation mode at a given applied stress. In such a deformation scheme, suppressing the migration of the subgrain boundaries and stabilization of the grain boundary are necessary to increase the creep resistance of the alloys. This can in principle be done by stabilization of the fine particles such as M23C6 and the Laves phase inside grains and along boundaries. Solution hardening by elements like Mo, W and N is an essential fundamental strengthening mechanism to be used for the austenitic alloys, while phase stability against σ embrittlement is required for long-term creep ductility and the resultant creep strength.28 These strengthening methods have successfully been used in the developed steels such as SUPER304H with M23C6, Cu and NbN and HR3C with M23C6, MX and NbCrN as shown in the tree chart of the developed austenitic stainless steels in Fig. 20.21. TP347HFG (fine-grained 18Cr–12Ni–Nb steel) is widely used as superheater and reheater tubes in fossil fired boilers.28 The steel has been developed to improve the steam oxidation resistance of conventional TP347H stainless
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Creep rate
Suppressing (1) rearrangement of dislocations (2) migration of sub GBs (3) GB embrittlement
Hardening and introduction of dislocations Solution hardening (H,Mo,W etc) Fine particles such as MX, NbCrN, Cu, M23C6, α-Cr, γ’, Laves
Solution hardening (Mo, W etc) Phase stability (N, Ni etc) Stabilization of boundary precipitates such as M23C6, Laves Long-term stabilizatin of fine particles
20.20 Guiding principles for further strengthening the austenitic alloys.
steel by grain refinement through a specially established thermomechanical process. The microstructure of the steel consists of an austenitic fine-grained matrix strengthened by M23C6 carbide mainly along the grain boundary and finely dispersed NbC carbide in the matrix. NbC is fine and stable even after long-term creep exposure at high temperatures. Figure 20.22 shows creep rupture data for TP347HFG tubes with average curves by the Larson–Miller parameter method.29 The longest creep rupture time for TP347HFG tubes is about 60 000 h at 600°C. Their long term creep strength is very stable and no degradation in creep strength is expected at the temperatures up to 750°C. Figure 20.23 shows the manufacturing process for establishing fine-grained microstructures in 18Cr–9/12Ni–Nb steels, which is characterized by double stage heat treatment in order to achieve a fine grain microstructure with a very fine dispersion of NbC in the matrix, compared with that for the conventional TP347H steel.30 In the developed process, NbC resolves into the matrix during the presolution treatment at higher temperatures and reprecipitates finely in the matrix during the subsequent solution treatment at lower temperatures. This gives rise to a fine grain microstructure with a very fine dispersion of NbC in matrix. Figure 20.24 shows the initial microstructures of TP347HFG tubes and conventional TP347H steel. A homogeneous and very fine grain structure has been achieved by using a double stage heat treatment in the developed process.30 Figure 20.25 shows the change in the amount of Nb precipitated as NbC in TP347HFG after presolution treatment at various temperatures.30 NbC precipitated more when using the lower temperature solution treatment after a higher temperature presolution treatment.
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+Cu,Nb
High-Cr austenitic steels (22 ~ 25Cr)
25Cr–20Ni–Nb–N; HR3C (KA-SUS310J1TB, ASME CC2111, TP310HCBN)
18Cr–9Ni–3Cu–Nb–N; SUPER304H (KA-SUS304J1HTB, ASME CC2328)
+ Nb, N
Strength
Fine Grained 18Cr–12Ni–Nb; TP347HFG KA-SUSTP347HTB
SUS321HTB
SUS347HTB (Fine grain) + Ti
+ Nb
(KA-SUS309J2TB) –C, + Mo, N
16Cr–12Ni–Mo (SUS316HT, AISI316H) 18Cr-8Ni (SUS304HTB, AISI304H)
25Cr-20Ni (SUS310TB, AISI310) +Cr, Ni +Cr, Ni 22Cr-12Ni (SUS309TB, AISI309)
18Cr-8Ni (AISI302)
Corrosion resistance
20.21 Tree chart of the developed austenitic steels.
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18-8 type austenitic steels (18Cr)
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559
500 400 300
Stress (MPa)
200 600°C 650°C 100 90 80 70 60 50
Test Temp.
700°C
600°C 650°C
750°C
700°C
40
750°C
30
800°C
800°C
20 101
102
103 Rupture time (h)
104
105
20.22 Creep rupture data for TP347HFG (fine-grained 18Cr–12Ni–Nb steel) tubes. Developed process
Conventional process
precipitation
Fine NbC dissolution
Cold working CW 30%
Solution
Pre-solution
Fine NbC
SUPER304H TP347HFG
CW 20~30%
Fine grain & fine NbC
TP347H
Coarse grain & coarse NbC G.S.No.; 4~5
G.S. No.8
20.23 Manufacturing process to establish fine-grained microstructures in 18Cr–9/12Ni–Nb steels compared with the process for conventional TP347H steel.
Figure 20.26 shows the effect of final-solution treatment temperature on the grain size of TP347HFG tubes and conventional TP347H steel.30 The grain size of TP347HFG tubes is much smaller than that of the conventional steel even with an increase in the solution treatment temperature. This was achieved by more precipitation of fine NbC carbide during the presolution treatment at lower temperatures.
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100µ (a)
(b)
20.24 Optical micrographs of fine grain microstructure for (a) TP347HFG attained using a double stage heat treatment compared with (b) treatment for TP347H. 1.0 Pre-solution treatment 1250°Cx 10min 1300°Cx 10min
Extracted Nb content (%)
0.8
0.6
0.4
0.2
Precipitated Nb as NbC during final-solution treatment
0 As presolution
1100 1150 1200 1250 Solution treatment temperature (°C)
20.25 Change in the amount of Nb precipitated as NbC in TP347HFG after presolution treatment at various temperatures.
Figure 20.27 shows the oxidation behavior of TP347HFG with different grain size tested at 650°C and 700°C. Steam oxidation resistance improves with reducing grain size of the steels. This steam oxidation resistance is achieved by the thin and tight protective Cr2O3 corundum-type inner oxide layer formed in the fine-grained steel (see Figure 20.28)30. SUPER304H (18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) is used as superheater and reheater tubes in fossil fired boilers.31 The steel has been developed to substitute for conventional Type304H and Type321H steels by the addition of copper and nitrogen to increase creep strength at elevated temperatures and toughness after long-term exposure at high temperatures. The microstructure of the steel consists of an austenitic matrix strengthened by M23C6 carbide mainly along the grain boundary and a finely dispersed Cu-phase and NbCrN nitride in matrix. The Cu-phase is
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12
Grain size (ASTM No.)
10 8 6 4 2
TP347HFG Conventional TP347H
0 1100 1200 1300 Solution treatment temperature (°C)
Inner-scale thickness (µm)
20.26 Effect of solution treatment temperature on grain size of TP347HFG and conventional TP347H steel.
650°C 700°C
100
TP347H GS6) TP347HFG(GS8) TP347HFG(GS9)
50
20
In steam
10 500
1000
3000 Time (h)
20.27 Steam oxidation resistance of TP347HFG tested at 650°C and 700°C. Outer scale Inner scale
Outer scale
Inner scale
Metal
Metal
Fe3O4 Fe3O4 Cr2O3 (Fe, Cr)3O4 Cr2O3 (a)
(Fe, Cr)3O4 (b)
20.28 Schematic illustration of oxide layers formed in (a) fine grained and (b) coarse grained 18Cr austenitic steels.
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fine and coherent is the matrix giving rise to a significant increase in creep strength at elevated temperatures. NbCrN is also fine and stable even after long-term exposure at high temperatures. No significant sigma phase is expected to form even after 10 000 h in the temperature range between 600 and 800°C, mainly achieved by stabilization of the austenitic matrix with copper and nitrogen. Figure 20.29 shows creep rupture data for SUPER304H (18Cr–9Ni–3Cu– Nb–N steel; KA-SUS304J1HTB, ASME CC2328) with average curves by the Larson–Miller parameter method.31 The longest creep rupture datum of SUPER304H tubes is over 85 000 h at 600°C. Their long-term creep strength is very stable and no degradation in creep strength is expected at the temperatures up to 750°C. Figure 20.30 shows the allowable stress determined for SUPER304H(18Cr– 9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) compared with that for the conventional steel, TP347H and the corresponding strengthening mechanisms used.32 The microstructural change in SUPER304H tubes after aging for up to 10 000 h in the temperature range between 600°C and 750°C is available in the literatures,32,33 There is no significant microstructural change observed even after aging for 30 000 h at 750°C. A detailed TEM observation of the specimens aged for 3000 h in the temperature range between 600°C and 750°C has shown that a fine coherent Cu phase is dispersed in the matrix as well as fine NbCrN nitrides and no harmful blocky precipitation such as a sigma phase is found. This fine dispersion of the precipitates is the major 500 400 300 600°C
Stress (MPa)
200
650°C 100 90 80 70 60 50
700°C
Test Temp. 600°C
750°C
650°C
40
700°C
30
750°C 800°C
800°C 20 101
102
103 Rupture time (h)
104
105
20.29 Creep rupture data for SUPER304H (18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328).
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Allowable tensile stress (MPa)
180 160 140 120
Tensile region
563
Creep region
Solid solution strengthening (N) 18Cr–9Ni–3Cu–Nb–N
100
Precipitation strengthening Nb (C,N) NbCrN M23C6 Cu phase
80 TP347H 60 40 20 0 0 100 200 300 400 500 600 700 800 Temperature (°C)
20.30 Allowable tensile stress determined for SUPER304H(18Cr–9Ni– 3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) according to the Japanese METI standard.
strengthening mechanism of this steel in the creep region at higher temperatures, schematically depicted in Fig. 20.30.34 Figure 20.31 shows change in toughness of SUPER304H tubes after aging for up to 30 000 h in the temperature range between 500°C and 750°C. Toughness is reduced by aging for a short period but keeps at a high level even after aging for 30 000 h at 750°C.34 HR3C (25Cr–20Ni–Nb–N steel; KA-SUS310J1TBÅCTP310HCbN, ASME CC2115) is used for superheater and reheater tubes in fossil fired, black liquor recovery and refuse fired boilers.35 An improved TP310 steel has been developed by the addition of niobium and nitrogen with an increase is creep strength at elevated temperatures. The microstructure of the steel consists of an austenitic matrix strengthened by M23C6 carbide mainly along the grain boundary, and fine dispersed NbCrN nitride in matrix. NbCrN is fine and stable even after long-term creep exposure at high temperatures. No significant sigma phase has been found after aging for more than 30 000 h in the temperature range between 600 and 800°C, achieved by optimizing the Nibalance. Figure 20.32 shows creep rupture data for HR3C tubes with average curves by the Larson–Miller parameter method.36 The longest creep rupture datum of HR3C tubes is for about 90 000 h at 700°C. Their long-term creep strength is very stable and no degradation in creep strength is expected at the temperatures up to 750°C. The microstructural change in HR3C tubes after aging for up to 10 000 h in the temperature range between 600°C and 750°C is reported in the
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Charpy impact value (J cm–2)
250
200
150
100
500°C 550°C 600°C
50
650°C 700°C 750°C
0 as sol.
1
10
104
100 1000 Aging duration (h)
105
20.31 Change in toughness of SUPER304H(18Cr–9Ni–3Cu–Nb–N steel; KA-SUS304J1HTB, ASME CC2328) after aging for up to 30 000 h at 500°C and 750°C. 500 400 300 600°C
Stress (MPa)
200
650°C 100 90 80 70 60 50
700°C
Test Temp. 600°C
750°C
650°C
40
700°C
30
750°C
800°C
800°C 20 101
102
103 Rupture time (h)
104
105
20.32 Creep rupture data for HR3C (25Cr–20Ni–Nb–N steel; KASUS310J1TB, TP310HCbN).
literature.37,38 There is no significant microstructural change observed even after aging for 10 000 h at 750°C. A detailed TEM observation of the specimens aged for 3000 h in the temperature range between 600°C and 750°C has shown that fine NbCrN nitrides dispersion is identified in the matrix and no harmful blocky precipitation such as the sigma phase is found.
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20.3.4 Fe–Ni based austenitic alloys used for 700°C AUSC plant In Europe and the USA, new research projects, THERMIE AD70039 and DOE USC project40 have been initiated to achieve much higher temperature and pressure conditions such as over 700°C at 35 MPa and 760°C at 35 MPa, respectively. This is, in fact, a challenging program for the metallurgists since none of the ferritic and austenitic steels developed seems to survive in such a hostile environment, and there is no experience of using Ni–Co base alloys for large diameter and heavy wall thickness piping applications mainly owing to their poor toughness, fatigue resistance and workability. Figure 20.33 shows creep strength data of the candidate austenitic alloys for 700°C A-USC boilers, compared with those of the advanced 9/12%Cr ferritic steels.27,40 It can be seen that only Ni–Co base alloys such as Alloys 740 and 617 strengthened by γ ′ are applicable with respect to creep strength at 700°C, while HR6W, Fe–Ni alloy is marginal and SUPER304H, austenitic stainless steel cannot be used at present. Figure 20.34 shows creep rupture data for the Fe–Ni alloy, 23Cr–43Ni– 7W at temperatures between 650 and 800°C. Unlike Alloys 740 and 617, this alloy is strengthened by a combination of the dispersion of precipitates such as M23C6, MX and the Laves phase and a solid solution of W with stabilization of the precipitates by B. This is confirmed by a TEM observation showing that a fine dispersion of the Laves phase as well as M23C6 and MX which serve as effective obstacles to the dislocation in motion remains even after long term-creep deformation for 58 798 h at 700°C with a stress of 98 MPa, as shown in Fig. 20.35.
Stress (MPa)
1000
600°C × 105h 650°C × 105h
700°C × 105h
Ni-Co base allosy
Alloy 740 Alloy 617 mod 100 Alloy 617
9.12Cr
HR6W Alloy 800H SUPER304H 10 20
22
24 T(log t + 20) × 103
26
28
20.33 Creep strength of the candidate materials for 700°C USC boilers.40
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650°C 700°C 750°C 800°C
300 250 200
Stress (MPa)
566
150
100
50 101
102
103 Rupture time (h)
104
105
20.34 Creep rupture data for 23Cr–43Ni–7W alloy.
Laves
MX 1µm
M23C6
250nm
27.35 TEM micrographs of the specimen of 23Cr–43Ni–7W alloy crept and ruptured after 58 798.4 h at 700°C.
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Note that the phase stability against σ phase formation is a key to achieving both high strength and enough ductility in these austenitic alloys. This has been demonstrated by 23Cr–43Ni based model alloys with Mo or W. Figure 20.36 shows the Larson–Miller parameter plot of the creep rupture data for 23Cr–43Ni based model alloys with 3%/5%Mo or 5%/7%W crept at 700, 750 and 800°C. It can be seen that the alloys with W are superior to those with Mo for longer term creep strength at high temperatures. To analyze this difference between the alloys with Mo and those with W, the microstructural evolution during creep deformation at high temperatures has been extensively examined using a thermodynamic calculation and TEM observation. Figure 20.37 shows a comparison of the equilibrium phase diagrams for the alloys with 5%Mo and 7%W calculated by Thermo-Calc. In the alloy with 5%Mo, the σ phase is the dominant phase with M23C6 and a small amount of MX at the temperatures between 700 and 800°C, while in the alloy with 7%W, the Laves phase is the major precipitate with M23C6 and MX. This characteristic difference in the phase equilibria of the alloys has been confirmed by optical micrographs and extraction replicas from specimens aged for 3000 h at 750°C, shown in Fig. 20.38. In the alloys with Mo coarse blocky precipitates identified as σ phase form along grain boundaries, with a small amount of fine precipitates, M23C6 and coarse Laves phase inside the grains. It is, however, noted that in the alloys with W, no blocky σ phase has been identified and instead fine Laves phase forms homogeneously with fine precipitates, M23C6 and MX both along the grain boundary and inside the grains. It is thus considered that in the alloys with W long-term creep strength 200 M3 M5 W5 W7
Stress (MPa)
150 700°C × 105h
100
80 21.5
22.0
22.5
23.0 23.5 T(log t + 20) × 103
24.0
24.5
25.0
20.36 Larson–Miller parameter plot of creep data for 23Cr–43Ni alloys with 3%/5%Mo or 5%/7%W.
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Creep-resistant steels 0.20 M5(5%Mo)
Mole fraction of precipitates
0.18 0.16 0.14 0.12 σ
0.10 0.08 0.06 0.04
M23C6
0.02 0 500
MX
αCr 1000 Temperature (°C) (a)
1500
0.10 W7(7%W)
Mole fraction of precipitates
0.09 0.08 0.07 0.06 0.05 0.04
Laves αCr
0.03 0.02 0.01 0 500
C23C6 MX µ 1000 Temperature (°C) (b)
1500
20.37 Equilibrium phase diagrams for (a) 23Cr–43Ni–5Mo and (b) 23Cr–43Ni–7W calculated by Thermo-Calc.
has been achieved by fine dispersion of the Laves phase along with M23C6 and MX, while in the alloys with Mo, coarse precipitates such as the σ phase along the grain boundary and the Laves phase inside grains have been detrimental to the long-term creep strength. To explore further increase in creep strength of these alloys, the possibilities of increase in the Laves phase and of introduction of γ ′ and α-Cr phases have been examined using phase diagram calculations based on 0.08C– 23Cr–43Ni–7W–0.1Ti–0.2Nb–0.003B, as shown in Fig. 20.39. The amount of Laves phase can be increased to a great extent with increasing W without
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Alloy design philosophy of creep-resistant steels M5
M5
M3
σ
σ
569
M5
M23C6 1µm
5µm Laves W7
W5
W7
W7
M23C6 Laves
5µm
20µm
1µm
Mole fraction of precipitates
0.08
Laves
0.07 0.06 0.05 0.04 0.03
µ
αCr
M23C6
0.02
MX
0.01 0 500
0.20
Mole fraction of precipitates
0.20
0.10 23Cr-43Ni-10W 0.09
0.18
1000 Temperature (°C) (a)
0.14 αCr
0.10 0.08 0.06 0.04
µ M23C6
0.02 0 500
σ MX 1000 Temperature (°C) (c)
0.14
1500
γ′
0.12 0.10 0.08 0.06
αCr
0.04 0.02
0.20
0.16
23Cr-53Ni-7W-2Ti-0.5Al
0.16
1500
30Cr-43Ni-7W
0.12
0.18
M23C6 µ
0 500
Mole fraction of precipitates
Mole fraction of precipitates
20.38 Optical microstructures and extraction replicas from 23Cr–43Ni with 3/5Mo or with 5/7W alloys aged for 3000 h at 750°C.
0.18
MX 1000 Temperature (°C) (b)
1500
Alloy 120
0.16 0.14 0.12
σ
0.10 0.08 0.06 0.04 0.02 0 500
αCr Laves MX M23C6 1000 Temperature (°C) (d)
1500
20.39 Calculated phase diagrams for stabilizing Laves, γ’ and α-Cr in 23/30Cr-43/53Ni-7/10W-Ti-Al alloys and MX phases in Alloy120. WPNL2204
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any other phase changes. γ ′ phase can be introduced by increasing Ti and Ni, while, on the other hand, the Laves phase decreases and α-Cr phase and MX increase at higher temperatures. The α-Cr phase can be stabilized by increasing Cr and Ni, while the Laves phase decreases to a great extent. MX is found to be used effectively in Alloy120, although the σ phase is dominant at the temperatures around 700°C. It is thus concluded from these calculations that there are several possibilities for further strengthening of the austenitic alloys. It is, however, noted that phase stability against σ phase formation is a key to achieving both high strength and enough ductility in these austenitic alloys to be used for 700°C A-USC boilers.
20.4
References
1 V. K. Sikka, C. T. Ward and K. C. Phomas, ASM International Conference Production, Fabrication, Properties and Application of Ferritic Steels for High-Temperature Applications, ASM, Washington DC, 1981. 2 F. Masuyama, I. Ishihara, T. Yokoyama and M. Fujita, Thermal Nuclear Power, 1995, 46, 498. 3 K. Muramatsu, Proceedings Advanced Heat Resistant Steels For Power Generation, R. Viswanathan and J. Nutting (eds), The University Press, Cambridge, 1998, 543. 4 F. Masuyama, Proceedings Advanced Heat Resistant Steels For Power Generation, R. Viswanathan and J. Nutting (eds), University Press, Cambridge, 1998, 33. 5 Sumitomo Seamless Tubes and Pipe Creep Data Sheets, Sumitomo Metal Industries, 1993. 6 Y. Sawaragi, N. Otsuka, H. Senba and S. Yamamoto, Sumitomo Search, 1994, 56, 34. 7 T. Kan, Y. Sawaragi, Y. Yamadera and H. Okada, Proceedings 6th International Conference on Materials for Advanced Power Engineering 1998, Liege, Forschungszentrum Julich GmbH, 1998, 60. 8 Y. Sawaragi, H. Teranishi, A. Iseda and K. Yoshikawa, Sumitomo Search, 1990, 44, 146. 9 M. Igarashi, H. Okada and H. Semba, Proceedings 8th Workshop on the Innovative Structural Materials for Infrastructure in 21st Century, NIMS, Tsukuba, 2004, 194. 10 M. Igarashi, M. Yoshizawa, H. Okada, H. Matsuo, Y. Yamadera and A. Iseda, CAMP ISIJ, 2003, 17, 336 (in Japanese). 11 T. Otsuka, Y. Yamaji, S. Takenaka, M. Ichiryu, H. Momma and S. Takano, Thermal Nuclear Power, 2006, 57, 734. 12 N. Komai, F. Masuyama, I. Ishihara, T. Yokoyama, Y. Yamadera, H. Okada, K. Miyata and Y. Sawaragi, Advanced Heat Resistant Steels For Power Generation, University Press, Cambridge, 1998, 96. 13 Y. Sawaragi, K. Miyata, S. Yamamoto, F. Masuyama, N. Komai and T. Yokoyama, Advanced Heat Resistant Steels For Power Generation, University Press, Cambridge 1998, 144. 14 F. Masuyama, T. Yokoyama, Y. Sawaragi and A. Iseda, Materials for Advanced Power Engineering, Part 1, Kluwer Academic Publishers 1994, 173. 15 Y. Sawaragi, T. Kan, Y. Yamadera, F. Masuyama, T. Yokoyama and N. Komai, Proceedings of the 6th International Conference on Materials for Advanced Power Engineering 1998, Liege, Forschungszentrum Julich GmbH, 61.
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Alloy design philosophy of creep-resistant steels 16 17 18 19 20 21 22
23
24 25 26
27 28 29 30
31 32 33 34 35 36 37 38 39
571
M. Yoshizawa, A. Iseda and M. Igarashi, CAMP ISIJ, 2005, 18, 1549. (in Japanese). K. Miyata, M. Igarashi and Y. Sawaragi, ISIJ Inte., 1999, 39, 947. K. Miyata and Y. Sawaragi, ISIJ Inter., 2001, 41, 281. A. Iseda, Y. Sawaragi, S. Kato and F. Masuyama, Proceedings of the Fifth International Conference on Creep Materials, Florida, 1992, 389. Y. Sawaragi, A. Iseda, K. Ogawa and F. Masuyama, Materials for Advanced Power Engineering, Part 1, Kluwer Academic Publishers 1994, 309. M. Yoshizawa and M. Igarashi, International Journal of Pressure Vessels and Piping, 2007, 84, 37. M. Igarashi, M. Yoshizawa, A. Iseda, H. Matsuo and T. Kan, Proceedings the 8th Liege Conference on Materials for Advanced Power Engineering, J. Lecomte-Beckers, F. Schubert and P.J. Ennis (eds), Liege, Forshungszentrum, Jülich GmbH, 2006, Volume II, 1095. L. Kloc and V. Sklencka, Proceedings 6th Liege Conference on Materials for Advanced Power Engineering, J. Lecomte-Beckers, F. Schubert and P. J. Ennis (eds), Liege, Forshungszentrum, Jülich GmbH, 1998, Volume, I, 215–222. L. Kloc and V. Sklencka, Mater. Sci. Eng. A, 1977, 234–236, 962. K. Kimura, Proceedings of ASME Pressure Vessel and Piping Conference, Denver, CO, 17–21 July, ASME, 2005, Paper 71039. M. Igarashi, S. Muneki, H. Kutsumi, T. Itagaki, N. Fujitsuna and F. Abe, Proceedings 5th International Charles Parsons Turbine Conference, A. Strang, W.M. Banks, R.D. Conroy, G.M. McColvin, J.C. Neal and S. Simpson (eds), University Press, Cambridge, 2000, 334. M. Igarashi, H. Okada and H. Semba, Proceedings 9th Workshop on the Innovative Structural Materials for Infrastructure in 21st Century, NIMS, Tsukuba, 2005, 96. Y. Sawaragi, N. Otsuka, H. Senba and S. Yamamoto, Sumitomo Search, 1994, 56, 34. Creep Properties of Heat Resistant Steels and Superalloys, Landolt-Bornstein New Series VIII-2B, Springer, 2004, 251–257. H. Teranishi, K. Yoshikawa, H. Fujikawa, M. Kubota, K. Tokimasa and M. Miura, Proceedings of International Conference on Coatings and Bi-Metallics for Energy Systems Chemical Process Environment, ASM Conference, South Carolina 1984. Creep Properties of Heat Resistant Steels and Superalloys, Landolt-Bornstein New Series VIII-2B, Springer, 2004, 260–264. Y. Sawaragi, N. Otsuka, K. Ogawa, S. Kato and S. Hirano, Sumitomo Search, 48, 1992, 50. Y. Sawaragi, N. Otsuka, K. Ogawa, S. Kato and S. Hirano, Sumitomo Metals, 1991, 43, 24 (in Japanese). H. Senba, Y. Sawaragi, K. Ogawa, A. Natori and T. Kan Materia, 2002, 41, 120 (in Japanese). Creep Properties of Heat Resistant Steels and Superalloys, Landolt-Bornstein New Series VIII-2B, Springer, 2004, 292–296. Y. Sawaragi, Y. Teranishi, A. Iseda, K. Yoshikawa, Sumitomo Search, 1990, 44, 146. Y. Sawaragi, H. Teranishi, H. Makiura, M. Miura and M. Kubota, Sumitomo Metals, 1985, 37, 66. Y. Sawaragi, Y. Teranishi, A. Iseda and K. Yoshikawa, Sumitomo Metals, 1990, 42, 260 (in Japanese). R. Blum and R. W. Vanstone, Proceedings 6th International Charles Parsons Turbine Conference, 16–18 September 2003, Dublin, Ireland, A. Strang, R. D. Conroy, W.
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M. Banks, M. Blackier, J. Leggett, G. M. McColvin, S. Simpson, M. Smith, F. Starr and R. W. Vanstone (eds), 2003, 487–510. 40 R. Viswanathan: Proceedings EPRI Conference on Materials and Corrosion Experience for Fossil Power Plants, US Program on Materials Technology for USC Power Plants, South Carolina, USA 2003.
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21 Using creep-resistant steels in turbines T.- U. K E R N, Siemens AG Power Generation Group, Germany
21.1
Introduction
Energy is the source of general well being and the standard of living in each country. Recent history has shown that a safe and sufficient energy supply is the basis for the development of the whole population of a country and region. The world situation today is characterised by a constantly increasing population and the desire for improved living conditions. The industrialisation of a country is automatically connected to the availability of electricity. It is the key to progress. In the decades to come, there will continue to be heavy reliance on fossil fuels, such as coal, oil and natural gas, because nuclear technologies are always under discussion. Research into fusion reaction is still under way and could take a further 20 years to reach industrial importance. Because reserves of oil and natural gas are unlikely to be sufficient to satisfy fully the projected increased power demands and to fill the gap until the introduction of further advanced renewable energy sources, coal may be the only fuel available in substantial quantities around the world. This will force power plant utilities to use the most advanced technologies available to increase the efficiency of their power plants to meet the increasingly stringent emission regulations for safeguarding health and preserving the environment for future generations. One of the best known and widely used standard technologies is the steam power plant using a steam generator (e.g. boiler), a steam turbine, condenser, generator and the thermodynamics of the steam–water cycle for the energy conversion of heat to electricity. A typical steam turbine turbo-set arrangement is shown in Fig. 21.1. The components are, from left to right: high pressure turbine HP with valves, intermediate pressure turbine IP with valves, two low pressure turbines LP, the generator and, underneath, the steam condenser. This chapter deals with the application of creep-resistant steels in the steam turbine components in the past, today and in the future. 573 WPNL2204
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IP + valves 610°C/60bar
2 × LP 350°C
Generator
HP + valves 600°C/300bar
Condenser
21.1 Steam turbine arrangement, example for a 800 MW unit.
21.2
Implications for industries using creep-resistant steels
There are different turbine arrangements for a power plant possible depending on the customer requirements and the technology of the turbine supplier. But one thing remains: the extensive use of steels in the overall construction. For example, with a turbine arrangement for 800 MW as shown in Fig. 21.1, an overall weight of steels of 1800 tonnes is required, comprising 130 tonnes HP, 350 tonnes IP and 1300 tonnes LP parts. For temperatures higher than 400°C, mainly creep-resistant steels are applied because of their excellent mechanical, chemical and physical properties, and easy manufacturability. Depending on the function, the requirements in service and the specifically applied design, the design criteria for different turbine parts and components in high pressure HP, intermediate pressure IP and low pressure LP turbines or combined turbine variants are also different. The main critical components are covered in Table 21.1 from rotors and casings, to blades and bolts. The property profile requested covers all areas of the material behaviour. It ranges from static strength, creep rupture strength, toughness, fatigue properties and crack growth to the influence of the environment characterized by corrosion, erosion and oxidation behaviour. For a safe design, different strength and toughness criteria have to be applied to ensure the safe operation of the turbo-set during its designed lifetime. Material properties have to be determined by standardized tests with specific specimens. Transfer of the results to the actual components is made by models and rules based on official standards, whereas special aspects
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Table 21.1 Material data requirements for application in steam turbine design Requirements
HP- & IP-rotors
LP-rotors, LP-discs
HP- & IP-casings
stress corrosion cracking corrosion fatigue Erosion behaviour Oxidation behaviour HP = high pressure turbine, IP = intermedium pressure turbine, LP = low pressure turbine.
LP-blades
HP- & IP-bolts
Using creep-resistant steels in turbines
Static strength: tensile strength Creep rupture strength: creep behaviour Toughness: fracture toughness Fatigue properties: low cycle fatigue (LCF) high cycle fatigue (HCF) Crack growth: static – creep CG alternating – fatigue CG Corrosion: local corrosion
HP- & IP-blades
575
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are investigated in research and development (R&D) projects and are brought into industrial practice. The summary in Table 21.2 shows the different criteria that play a key role for engineers. Depending on the service temperature, the material properties are sensitive to the time of application. Typical material behaviour for steels at T > 400°C is the phenomenon of creep. It is one of the most critical properties in application because the design life is increasing from 100 000 h to 200 000 h operation time owing to current customer requests. Long-term testing has to be performed and accurate extrapolation methods have to be developed and validated. This point is discussed in more detail in other chapters of this book. The description in Table 21.2 helps to clarify the different loading data during real operation and to mirror these with the possible material data. The loading data are: • • • • • •
the different types of stress σi stress changes ∆σi strain changes and differences ∆εi operation temperature Ti start-up and shut-down operation cycles Ni the overall operation time ti.
As an example, the different types of stress sources for rotors during operation, start-up and shut-down are schematically shown in Fig. 21.2. Centrifugal force and temperature gradients cause stresses which the component has to withstand during its operation time and start-up and shut-down cycles. Additional local loading occurs if there are natural or manufacturing defects in the components volume. Different kinds of analysis have to be performed to ensure reliable and safe operation of the components. Starting with the stress analysis, fracture mechanic evaluation and fatigue evaluation have to be performed. The respective material properties used are determined in the following material tests: • • • • • • • •
Tensile: 0.2 yield strength (YS), ultimate tensile strength (UTS), elongation Creep rupture: creep rupture strength Ru for a time t at temperature T, Ru/t/T Creep elongation: creep elongation limit Rpε for a time t at temperature T, Rpε/t/T Fracture toughness: linear–elastic static fracture toughness KIC Fatigue crack growth: amplitude of stress intensity ∆K, crack growth rate da/dN Creep crack growth: stress intensity K, crack growth rate da/dt Creep crack initiation: stress intensity for static crack initiation KIid and energy integral C* Low cycle fatigue: amplitude of strain ∆ε, cycles to failure Nfi.
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Table 21.2 Summary of design criteria for steam turbines Material data
σi
Stress evaluation 0.2YS, UTS, elongation Creep rupture strength Creep elongation limit
Limit values
Allowable values
Required material properties
0.2YS 0.8*Ru/t/T 0.8*Rpε/t/T
σi ≤ 0.2YS/S σi ≤ Ru/t/T/S σi ≤ Rpε/t/T/S
0.2YS/UTS/Elong. = f (T) Ru/t/T = f (T, t) Rpε/t/T = f (T, t)
a0 ≤ acrit,1/S
KIC after Long term exposure
∆σi
Fracture mechanic evaluation (a0 = start defect size) – short term safety a0, KIC acrit
∆ε
–
Ti Ni ti
long term safety cyclic : a0, da/dN static : a0, da/dt
Fatigue evaluation – with defect (a0) cyclic : da/dN = f (∆K) static : da/dt = f (Klid) – without defect (with notch) cyclic : Nfi = f (∆ε) static : tui = f (Ru/t/T)
∆ai ∆ai
} a0 + ∑ ∆ai ≤ acrit,2/S
Nfi (tui)
∑ N i /Nfi + ∑ t i /tu i ≤ Eallow
Nfi tui
da/dN = f (∆K, T) at low stress rate Creep crack initiation tA = f (Klid) = f (C*) ∆ε = f (N, T) Ru/t/T = f (t, T) Rpε/t/T
}
i, index for different service loading conditions; UTS, ultimate tensile strength; YS, yield strength; IC, mode I critical; Iid, mode I ideal elastic; u/t/T, ultimate strength for time t at temperature T.
Using creep-resistant steels in turbines
Imposed load data
577
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Pressure loading
Thermal stress Temperature Ta distribution αk · σ∆T
Ti ∆T
σ∆T
Thermal stress at the nut of turbine blade fixture Stress caused by pressure Stress caused by centrifugal force
21.2 Loading conditions for a turbine rotor, schematically.
The initial crack sizes a0 have to be determined by appropriate non-destructive testing during quality assurance in the material procurement process. One of the most applied material classes for critical turbine components operating at temperatures T > 400°C is the low alloyed CrMoV steel which is also used for forgings and castings. The Cr content ranges from 0.5% up to 2.25%. The balance of properties described in Tables 21.1 and 21.2 can be achieved in practice. This is demonstrated for a typical 1CrMoNiV rotor forging in Fig. 21.3 in terms of basic properties.1 The strength was investigated for several parts with so-called radial cores going through the whole rotor body in the radial direction. It was found to be stable whereas the toughness characterized by the 50% fracture appearance transition temperature (FATT) shows clear changes from the surface at 20°C to the centre area at about 80°C. The reasons for this are the cooling conditions during heat treatment of large parts with thermal gradients and the material type itself which results in an annealed martensitic structure at the surface and an upper bainite structure in the centre. The tougher martensite shows disadvantages for long-term creep behaviour, as shown in Figs 21.4 and 21.5, with investigations simulating the different possible microstructures of 1CrMoV steels.2 The lowest creep rate was achieved with the upper bainite structure. The results from a real rotor forging at the surface and in the centre are also given. They demonstrate that the cooling conditions have to be optimized, for example by slowing down, to get a homogenous upper bainite structure with the required best creep properties in the forging. The creep rupture strength achieved for times up to 90 000 h confirms the advantages of upper bainite. The fatigue strength was also determined, resulting again in a better behaviour for the upper bainite structure (Fig. 21.6).
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1.3 m
Heat treatment contour
30 Mg
Radial core
Rm 70 Z 700
50
Rp 0.2
40 30
600
20
A 500
10
Surface
Temperature (°C)
Ductility (%)
Stress (N mm –2)
800
Centre
100 80 FATT several rotors
60 40 20 Surface
Position in the forging
Centre
21.3 Balance of properties for a 1CrMoNiV rotor with respect to through-hardenability.1 101
Ferrite – Pearlite Martensite Lower bainite Upper bainite Outside As delivered Centre transverse
Strain (%)
100
10–1
10–2 101
102
103 Time (h)
104
105
21.4 Influence of microstructure on creep behaviour of 1CrMoV rotor steel at 530°C and 127 MPa.2
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Ferrite – Pearlite Martensite Upper bainite Lower bainite Outside As delivered Centre transverse Notched Kt = 4.5 Smooth
Stress (MPa)
800 Hot tensile test
300 200
Scatterband of 1% CrMoNiV-steels Mean value ±20% (SEW555) Specimen still running Specimen withdrawn
100 101
102
103 Time (h)
104
105
21.5 Influence of microstructure on creep strength behaviour of 1CrMoV rotor steel at 530°C.2
Ferrite – Pearlite Upper bainite Lower bainite Martensite With hold time Without hold time
4
Strain ∆εt (%)
2
. ε (% min–1) 5.5–6
± 20 min
100 8 6 4
2 Time: 2 · 10–1 102
2
5
R mt =ˆ ∆ε t creep E
103 2 5 104 Cycles to crack initiation N i
2
5
105
21.6 Influence of microstructure on fatigue behaviour of 1CrMoV rotor steel at 530°C.2
Summarizing these points, the martensitic structure at the surface of rotor forgings has to be avoided because the rim position later on holds the rotating blades of the turbine shaft which are highly loaded during operation. Heat treatment and manufacture have to be performed in an accurate way to ensure the right microstructure at the relevant points.
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Another important aspect is shown in Fig. 21.7. The investigations here were made to obtain information about the notch weakening behaviour of this type of steel. Notch weakening means that notched creep specimens fail earlier than the respective smooth specimens. The importance in practice is confirmed by the fact that the surface of a rotor always has areas similar to notches such as in the steam inlet area as well as in the blade attachment grooves. Using the wrong quality heat treatment, the steel can show a significant drop in rupture ductility and therefore also notch weakening. Furthermore, the application of 1CrMoV steel in bolts is common, and here the notches are clearly introduced by the thread itself. Cracks caused by the effect of notch weakening can form before the end of designed lifetime resulting in failure or damage. The delivery specification of the turbine makers has to give advice about the right heat treatment, also implying that they know their materials very well. Figure 21.8 demonstrates that not only does the creep rupture strength play an important role in the integrity of components but also the creep elongation and rupture ductility.1 Here the stresses of a stop and control valve in the cast design are calculated over the service lifetime. Starting with high elastic stresses at the inner wall (position A) and lower stresses at the surface (Position B) caused by the internal pressure, the equivalent stress σv has reduced after 100 000 h operation time at the inner surface but increased 900°C 1 h/oil + 750°C 2 h
Stress (N mm–2)
1000
1000°C 1 h/oil + 570°C 2 h
αk = 4.5
500 200 α = 4.5
100
䊊 Elongation (%) 䊉 Reduction of area (%)
50
100 80 60 40 20 0 0.1
1
10 102 103 104 Time to rupture (h)
105
0.1
1
10 102 103 104 105 Time to rupture (h)
21.7 Creep rupture behaviour of 1CrMoV steels at 550°C for two different heat treatments.1
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10
20 30 50
20
40
30 40 70
50
110
60
σv Elastic σv after 105 h Mises reference stress (N/mm2) Mean stress σ = 47 N/mm–2 according to design codes
B
Tangential creep strain (%)
Location A A
0.6
250 bar 540 °C
0.5 φ500 φ900
0.4 0.3 0.2
Location B
0.1 103
2 × 103
104 2 × 104 Time (h)
105 2 × 105
21.8 Creep stress and strain of a high pressure stop and control valve body from 1CrMoV cast steel.1
outside owing to relaxation and plastic deformation via creep. The tangential creep strain has significantly increased for the inner position A to 0.4% after 100 000 h which is, compared to the outer surface position B with 0.1%, a factor of four. If the real creep elongation reaches the rupture ductility for the appropriate stress state at any position, cracks can occur and propagate by creep crack growth. Cracks introduced by creep will propagate faster during the operational cycle via fatigue crack growth or creep crack growth mechanisms. Multiaxiality, reducing the toughness reserve of the material, can also strongly influence the creep and fatigue component behaviour. Design rules have to take into account the different operational, customer specific conditions and provide sufficient exact residual life tools to ensure the safe operation of the components for the designed lifetime.
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Improving the performance and service life of steel components
Up until the 1980s, based on the available materials and their property profile, especially the achievable creep strength, the thermodynamic parameters of steam turbines were restricted to main steam temperatures T of 540°C. The optimization in design and materials yielded an increase in T to 565°C in the 1990s. The material classes used then were creep-resistant low and high alloyed steels of the 1-2CrMo(V) and 11-12CrMoV type available of the time. Examples of typical creep-resistant steels for steam turbines in Europe are given in Table 21.3. The increasing demand of the energy market for further improved efficiencies in power plants and political efforts for climate protection on Earth by reducing CO2 emissions have resulted in extensive R&D initiatives for material development worldwide. The aim has been to increase the creep rupture strength of materials by keeping the other properties stable. At the beginning, single types of material were reviewed to identify the most promising ones. Different aspects of component application were taken into consideration, for example that rotors for high temperatures can reach large diameters and weight, or casings have to be weldable for repair and construction purposes. Finally, work was started with the martensitic 9– 12%CrMoV steels. The advantages are clearly visible: • • • • • • •
method of manufacture (forging, casting) good through hardenability up to 1200 mm diameter balance in strength and toughness (long and short term) weldability potential for high oxidation resistance ease of fabrication lower cost than austenites and Ni-base alloys
Table 21.3 Examples of creep-resistant steels for power plant application up to 565°C (in wt%) Component type
C
Cr
Mo
V
Rotor
0.30 0.22 0.19 0.22 0.20 0.17 0.17 0.18 0.17
1.20 2.10 11.0 12.0 13.0 1.20 1.30 2.25 1.30
1.10 0.85 0.60 1.00 1.00 0.50 1.00 1.00 1.00
0.30 0.30 0.20 0.30
Blade
Valve body Casing
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W
N
Nb
0.10
0.30
0.65
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In Europe efforts have been concentrated on the COST programme (Fig. 21.9). This is a long established European programme aimed at coordinating pre-competitive research activities in numerous areas of science and technology. Participants include turbine and boiler makers, manufacturers of forgings, castings, pipework and welding consumables, as well as testing institutes and utilities. The work has spanned the entire range, from basic research and development, through non-destructive and destructive testing of small experimental batches and large scale components, to standardization of specific steels and full-scale practical application in currently operating power plants. Action started with COST501 (1986–1997), was continued in COST522 (1998–2003)3–5 and is currently run as COST536 (2004–2009).6,7 It has established a strong trans-European network in this field. The effectiveness of past COST actions is exemplified by the introduction of newly developed martensitic steels for forgings, castings and pipework. These improved steels are in commercial operation in advanced European power stations and have made it possible to increase the operating steam temperatures from 530– 565°C to 580–600°C with a corresponding increase in thermal efficiency. In Japan a R&D programme was initiated by the Electrical Power Development Company (EPDC) in 1981 to explore the possibilities of developing and applying new materials for high-temperature steam power plants. The programme started with basic research into strengthening mechanisms in 9–12%Cr steels and has led to the introduction of a number of steels with much improved creep strength at temperatures of 600°C and above.8 Full scale components, such as rotors, casings and boiler tubes and pipes have been manufactured and employed in a series of advanced steam power plants ordered by EPDC. Further basic research was started 1997 with the NRIM-STX21 programme which aimed to find new alloy design concepts for boiler components. Current activities are concentrated on the transformation of the results from trial melts to real industrial applications. In the USA, collaborative work has been performed within the EPRI programme RP1403, which began in 1978 with basic studies and started practical work in 1986. Activities were focused on the development of steels for applications such as thick-walled pipes. This programme was international and work was carried out by companies in the USA, Japan and Europe. The project was successful in that the materials P92 (NF616) and P122 (HCM12A) have been code approved by ASME. In addition, thick-walled pipes and headers made from these steels have been manufactured and full-scale tested and are in operation in power plants in Europe. The criterion for success in all of these developments for new 9–12CrMoV steels was set as an improved creep rupture strength tested by conventional creep specimens. The other properties determined in standard tests have been accepted if they were not worse than those of 1CrMoV steels. A summary
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USA
EUROPE
R & D : EPDC
R & D : EPRI
COST 50/501
Manufacturers, utilities, EPDC
Manufacturers
1981 – 1991
Study 1978 – 1980
Manufacturers, steelworks, utilities and R & D – institutes
316 bar 566/566/566°C 314 bar 593/593/593°C 343 bar 649/593/593°C 50 MW Pilot power plant 1989 – 1990 1991 – 1993 1994 – 2000 300 bar 630/630°C
310 bar 566/566/566°C 310 bar 593/593/593°C 345 bar 649/649/649°C EPRI – RP 1403 – 15 300 – 900 MW R & D : 1986 – 1993 Steels and components for Boiler + Turbine (USA, Japan, ALSTOM + MAN)
NRIM – STX 21 Project: USC 650°C/350 bar Boiler
EPRI – RP 1403 –50 –WO9000 –38
1997 – 2012
1990 – 1999
Thick wall components
Thick wall pipes: P 92 + P 122 (USA, Japan, UK + Denmark)
1983 – 1997 300 bar 600/600/600°C 300 bar 600/620°C Rotors Cast components Bolting materials Pipes, tubes, Welds COST 522
COST 536
1998 – 2003
2004–2009
300 bar 620/650°C
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21.9 Development activities for the new 9–12CrMoV class steels.
Rotors, cast components, bolts, pipes, tubes, welds
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Japan
586
Creep-resistant steels
for the new European 600°C rotor materials is given as an example in Fig. 21.10.4 In comparison with the known 1CrMoV and 12CrMoV steels, important changes were made in the chemical composition: the carbon content is reduced and the elements nitrogen, niobium and sometimes tungsten have been added. For the most promising alloy, 9Cr–MoCoVB steel, the elements boron and cobalt play an important role in highly increased creep strength. Compared with conventional 1Cr steel, a 50–70 K increase in application temperature without any significant design changes is possible. The improvement within the class of 12Cr steels is also demonstrated at a 30–50 K higher temperature for the same 100 000 h creep rupture strength if we consider the 9–10Cr steel data. The basis for the extrapolation to a relevant design life of 100 000 h is to perform real long-term tests in the laboratory which so far in Europe have reached up to 80 000 h for different melts from real components.9 This is of great importance for the reliability of the extrapolation because it is known that the new materials can show changes in the precipitation status after 30 000–50 000 h resulting in a strong decrease in the long-term strength. A recent example is the pipe steel P122 with strongly reduced creep rupture data.10,11 The advantageous creep behaviour of the new steels enables the design of high temperature turbo-sets with the same design rules as for the conventional machines.
Steel 1CrMoV 12CrMoV 10CrMoV 10CrMoWV 9CrMoCoVB
C 0.28 0.21 0.12 0.12 0.12
Cr 1.0 12.0 10.0 10.0 9.0
Mo W 0.9 – 1.0 – 1.5 – 1.0 1.0 1.5 – (Weight%)
V 0.30 0.30 0.20 0.20 0.25
Nb – – 0.05 0.05 0.05
N – – 0.05 0.05 0.02
B – – – – 0.010
Co – – – – 1.0
100 000 h creep rupture strength
200 10%CrMo(W)VNbN10-1(-1)
9%CrMoCoVNbB 1%CrMoV 100 12%CrMoV appr. 30°C appr. 70°C
0 500
550
600
650
Temperature (°C)
21.10 Creep rupture strength at 100 000 h for new 600°C steels in Europe.
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Similar to results for forgings, the development of cast materials for valves and casings has been successful.7 It has resulted in alloys of 9Cr type with W, V, Nb, N and boron. Service temperatures can be increased up to 620°C depending on the material data. The advantage of having stationary parts like these loaded mainly by internal pressure lies in the possibility of reducing the service stresses by increasing the wall thickness. For this, the thermal flexibility characterized by loading via low cycle fatigue during start-up and shut-down operation and thermally introduced secondary stresses has to be maintained. In comparison, rotor stresses can only be handled when the material has a higher creep strength or the temperature is reduced by direct or indirect cooling. But this is not easily possible for casings and valves. The compositions of industrially applied cast materials for higher temperatures are given in Table 21.4. Component examples for a high temperature rotor forging and an inner casing of a steam turbine with 600°C main steam are shown in Figs 21.11 and 21.12. In Germany, the application of the new steels to real turbine components is accompanied by extensive research work in industry together with several institutes to determine the relevant service behaviour. An overview of the Table 21.4 High temperature cast materials applied in Europe (in wt%) Component type
C
Cr
Mo
V
W
N
Nb
Casing or valve
0.12 0.12
9.0 10.0
1.0 1.0
0.20 0.20
1.0
0.05 0.05
0.06 0.06
21.11 Rotor for high temperature application.
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21.12 Inner casing in 10Cr steel for high temperature application.
main activities is given in Fig. 21.13. It shows that programmes have been run or are still ongoing to check: • • • • •
the influence of component size and technical ranges for the chemical analysis on long-term creep strength9 the influence of the test location of a real power plant component on long-term creep strength up to 100 000 h for a safe lifetime extrapolation9 the creep rupture strength and low cycle fatigue behaviour under servicerelevant load changing conditions12,13 the high-temperature crack growth behaviour on fracture mechanics14 the influence of multiaxiality on material behaviour15
With the knowledge gained in these programmes the design can be optimized for new apparatus. In addition, lifetime evaluation during service with its continuously changing loading conditions is significantly improved.16,17 In the next part, two examples of material application for 565°C and 600°C steam turbines of the Siemens type are discussed: high pressure (HP) and intermediate pressure (IP) turbine classes. The barrel type design of the HP turbine is a typical feature of Siemens steam turbines and was introduced about 50 years ago (Fig. 21.14). The inner parts guide the stationary blades in a near perfect axis-symmetric body. No flanges are required as this vertically split and bolted casing is compressed by high pressure steam, which is contained in the inlet part of the outer casing and which has no horizontal or vertical split at all. Although over the years many different HP turbine designs
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Determination of component relevant material data
Uni-axial behaviour
Creep rupture strength Standard
Multi-axial
High temp. crack behaviour
Low cycle fatigue
Random Standard
Long term Influence of 100 000 h manufacturing
Influence of multi-axiality
Long term LCF
Single step
As-service relevant
21.13 R&D activities for qualification of 600°C steels in Europe.
1CrMoV cast steel (Divided if 600°C)
565°C: 1CrMoV cast 600°C: 10CrMoWV cast
565°C: 1CrMoV cast 600°C: 10CrMoWV cast
565°C: 1CrMoV 600°C 10CrMoWV
11–12CrMoV steel
12CrMoV steel
21.14 HP turbine material application for a 565°C and 600°C design.
have been used, the principle of pressure containment in a barrel has never changed. The barrels of today are for different applications, either built using an inlet and exhaust casing bolted together (for 600°C application, a material combination of 1CrMoV–10CrMoWV is used), or consist of one big barrel type outer casing of 1CrMoV cast material, closed with a bolted-on cover. The monoblock HP turbine rotor is made of either 1CrMoV steel or newly developed 10CrMoWV steel.
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In contrast to the single flow, barrel type HP turbine, the IP turbine consists of a top and a bottom half for inner and outer casings and is either a single or a double flow cylinder, depending on the power range (Fig. 21.15). Nowadays, the design of IP turbines allows reheat temperatures of more than 600°C. The double flow IP turbine shown in Fig. 21.15 was applied to a project with 610°C using a special internal cooling.18 With this cooling scheme, a temperature reduction of up to 16 K at the critical rotor surface section is achieved. The rotor is either made from 1CrMoV or 10CrMoWV steel depending on the steam temperature and loading conditions. The inner casing material choice shows the same temperature dependence as the rotors and is either made from 1CrMoV or 10CrMoWV cast steel. The blades in the high temperature region above 400°C are made from 10–12CrMoV(Nb) steel. Steels are not used in the highest inlet temperature areas. Ni-based alloys have to be applied here depending on the service stresses. For lower temperatures, a 13Cr steel is applied to take account of corrosion. The outer casing design requirements are well met with iron-based globular cast material at lower costs than steel casts. Other turbine manufacturers use a welded design for the rotors with 10Cr material in the steam inlet area and low alloyed steel in the area with T < 500°C.19–22 The blades in the steam inlet area with highest temperatures are often made with nickel base alloys or austenites.23,24 These examples show how deeply the design of turbo-sets is influenced by and connected to the class of creep-resistant steels. They allow flexible design and operation modes in service. Therefore continuous material
565°C: 1CrMoV 600°C: 10CrMoWV
Globular cast iron 565°C: 1CrMoV cast 600°C: 10CrMoWV cast
13CrMo
12CrMoV
11CrMoVNb
21.15 IP turbine material application for a 600°C design.
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improvements and development of this material class are required and driven by the industry.
21.4
Next steps into the future
In the 1980s and 1990s there were plans to design ultra super critical power plants using austenites for very high stressed components.25,26 The idea was to apply austenites up to temperatures of 650°C filling the gap between the martensitic materials (up to 600°C) and the Ni- or Co- based superalloys (>650°C). In Europe this concept was influenced by the experience already gained during the design, installation and service of small power plants with steam temperatures of T = 600–650°C and an installed power output of 7– 107 MW.27,28 The high temperature materials used here were austenites. These machines were designed in the early 1950s. Some of them are still in operation, for example at Eddystone in the USA.29,30 They have been operated as base load units with a low flexibility during operation, that is, fewer startup and shut-down operations. For today’s applications, the main market requirements for power plants have radically changed towards a higher power output of >300 MW, flexible operating conditions, short delivery times and cost-effective solutions. Austenites do not fulfil the design requirements resulting from the new market conditions. The main disadvantages of austenites are the physical properties and lower strength resulting in high thermal stresses in large components during start-up and shut-down. Therefore this material class is very susceptible to strain-induced cracking and thermomechanical fatigue characterized by the R-value (resistance-to-crack), see Fig. 21.16.18 A high R-value will result in a lower probability of thermal induced crack initiation in thick walls. Prototype tests with an austenite rotor have shown a significant reduction in service life owing to start and stop operation at the 50 MW Wakamatsu demonstration plant in Japan.31–33 The Ni-base alloys as well as the martensitic steels show higher R-values and are therefore better candidates for higher steam temperatures for thick walled components. At the same time, the manufacturing of large components like rotors and casings is a very special task which requires much more R&D effort to obtain fully realised components with the required homogeneous properties. As the success of the development of new martensitic creep-resistant steels increases the applicability of these materials, the development of new austenitic alloys seems to be less attractive as it was in the 1990s, especially if the superior material properties of the martensites are taken into account.34 They are very cost effective and offer a high flexibility during operation. Turbine companies today are applying the 9–10Cr steels in current designs not only for very high temperature machines but also in efforts to increase
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8000
R = 0.2YS*K/(E*α)
High R-value → low probability of thermal induced crack initiation
1Cr (550°C)
5000
Ni-base alloy
New 10Cr (600°C)
6000
K – thermal conductivity E – E-module at T α – linear coeff. of expansion
NiCr20TiAl
Martensites
Austenites
2000
1000
Type Type Type Type 347 304 316 321
Type 800
NF 709
X8CrNi MoVNb 16–13
3000
X8CrNi MoVNb 16–16
A 286
X8CrNi MoBNb 16–16
Temperature: 650°C 4000
Esshete 1250
Material parameter R
7000
0
21.16 Susceptibility to crack initiation for austenites and Ni-base alloys at 650°C.
the performance of highly stressed components for steam parameters T < 600°C. On the other hand, steam power plants with temperatures of up to 600–610°C can be supplied for today’s customer by applying a special design. The new generation of 9–11CrMoVCoB steels under development will enable further advanced steam parameters up to 630°C and therefore additionally increased efficiency. A survey of applicable material classes for steam turbine components is given in Table 21.5. It shows the expectations that the limit for creep-resistant steels is expected to be at about 650°C. The next technological step is already on the way. There are now concrete plans to design and build 700°C steam power plants. The basic work was started at the end of the 1990s in Europe.35,36 For 700°C steam power plants, steels no longer serve as candidate materials for the high temperature areas. Nickel- and cobalt-based superalloys play an important role. 37,38,39 Unfortunately the manufacture of large components using them is restricted.40 Therefore new design concepts including multi-material components have to be developed. Steels will serve for the parts where the temperatures are low enough to design with their properties. The connection to the Ni-based components will be made by welding or other comparable technologies. This means that the higher the application temperature of new steels, the smaller the amount of expensive Ni-base superalloys that have to be used.41,42 With this approach, the chances of realisation of 700°C steam power plants are greatly increased. The current plan of a large German utility is to build a 700°C demonstration power plant starting in 2010, to begin operation in 2014.43 This challenge
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Table 21.5 Materials for power plant components applied at high temperatures Temperature range Component
max. 565°C
600–610°C
620–650°C
≥ 700°C
Valve bodies
9–10Cr steel 9–10Cr steel
Turbine casing
1–2CrMo steel
9–10Cr steel
New 9–11Cr steel New 9–11Cr steel New 9–11Cr steel
Superalloys
Rotor
1–2CrMo steel 9–10Cr steel 1–2CrMo steel
Superalloys
Superalloys
has been taken up by European boiler and turbine makers. Strong efforts are now required to achieve this ambitious aim.
21.5
Summary
The importance of creep-resistant steels for steam turbine application is very great. Many of these steels are used in different compositions and for different components. The well-balanced property profile of these steels enables a safe and sound design of turbo-sets with lifetimes of >100 000 h. The development of new creep-resistant steels in the last decade has opened up the market for steam temperatures up to 620°C, resulting in higher efficiencies for fossil-fueled steam power plants and therefore meeting efforts for climate protection by reducing CO2 emissions. The next generation of 700°C power plants is already being planned. Here also steels will keep their important role for design purposes and cost efficiency, enabling special designs and optimized costs. In this way, an economically viable energy supply can be achieved despite the use of expensive nickel- and cobalt-based superalloys. Future R&D will concentrate on the development of new creep-resistant weldable steels for temperatures up to 650°C, improved superalloys for large and thick-walled components in turbine and boiler and cost-optimized superalloys to reduce overall costs for future steam turbines operating at temperatures of 700°C or above.
21.6
References
1 Muehle E E and Ewald J, ‘High-reliability steam turbine components – materials and strength calculation aspects’, 4th International Conference High Temperature Materials for Power Engineering COST-501 and COST-505, Liege, Belgium, 1990, Research Centre Juelich, Germany.
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2 Wiemann W, Ewald J, Niel K and Reiermann D, ‘Influence of microstructure of 1% CrMoV steel on creep and strain controlled fatigue behavior’, International Conference LCF of Materials, Munich, 1987, Deutscher Verband fuer Materialpruefung DVM, Germany. 3 Staubli M, Mayer K H, Kern T-U and Vanstone R W, ‘The European joint development program COST522’, 5th Proceedings International Charles Parsons Conference, Cambridge, UK, 2000, IoM Communication, UK. 4 Kern T-U, Staubli M, Mayer K H, Escher K and Zeiler G, ‘The European efforts in development of new high temperature rotor materials up to 650°C–COST522’, 7th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2002, Research Centre Juelich, Germany. 5 Scarlin B, Kern T-U and Staubli M, ‘The European efforts in material development for 650°C USC power plant – COST522’, 4th International Conference Advances in Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA. 6 Kern T-U, Staubli M, Mayer K H, Donth B, and Zeiler G, ‘The European efforts in development of new high temperature rotor materials–COST536’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany. 7 Staubli M, Hanus R, Weber T, Mayer K H and Kern T-U, ‘The European efforts in development of new high temperature casing materials - COST536’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany. 8 Masuyama F, ‘Alloy development and material issues with increasing steam temperature’, 4th International Conference Advances in Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA. 9 Mayer K H, Blum R, Hillenbrand P, Kern T-U and Staubli M, ‘Development steps of new steels for advanced steam power plants’, 7th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2002, Research Centre Juelich, Germany. 10 Igarashi M, Yoshizawa M, Iseda A, Matsuo H and Kan T, ‘Long term creep degradation in 12%Cr ferritic steel tubes and pipes’, International Conference ECCC Creep and Fracture in High Temperature Components – Design and Life Assessment, London, UK, 2005, IoM Communication, UK. 11 Iseda A, Yoshizawa M, Igarashi M and Kan T, ‘Long term creep strength degradation in T122/P122 steels for USC power plants’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany. 12 Schwienheer M, Haase H, Scholz A and Berger C, ‘Long term creep and creepfatigue properties of the martensitic steels of type (G)X12CrMoWVNbN10-1-1’, 7th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2002, Research Centre Juelich, Germany. 13 Berger C Schwienheer M and Scholz A, ‘Creep and fatigue properties of turbine steels or application temperatures up to 625°C’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany. 14 Mueller F, Scholz A and Berger C, ‘Crack behavior of 10Cr steels under creep and creep-fatigue conditions’, International Conference ECCC Creep and Fracture in
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23
24
25 26 27
28
29
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High Temperature Components – Design and Life Assessment, London, UK, 2005, IoM Communication, UK. Ringel M, Roos E, Maile K and Klenk A, ‘Constitutive equations of adapted complexity for high temperature loading’, International Conference ECCC Creep and Fracture in High Temperature Components – Design and Life Assessment, London, UK, 2005, IoM Communication, UK. Bagaviev A, ‘Steam turbine component integrity analysis based on high temperature fracture mechanics’, International Conference ECCC Creep and Fracture in High Temperature Components – Design and Life Assessment, London, UK, 2005, IoM Communication, UK. Bagaviev A and Sheng S, ‘Industrial application of creep/fatigue crack initiation and growth procedures for remaining life analysis of steam turbine components’, International Conference ECCC Creep and Fracture in High Temperature Components – Design and Life Assessment, London, UK, 2005, IoM Communication, UK. Feldmueller A and Kern T-U, ‘Design and materials for modern steam power plants’, Proceedings 5th International Conference Charles Parsons Conference for Advanced Materials, Cambridge, UK, 2000, IoM Communication, UK. Gerdes C and Bartsch H, ‘Steam turbine shafts of combined high pressure/low pressure rotors’, 14th International Conference International Forgemasters Meeting, Wiesbaden, Germany, 2000, German Iron and Steel Institute. Magoshi R, Nakano T, Konishi T, Shige T and Kondo Y, ‘Development and operating experience of welded rotors for high-temperature steam turbines’, International Conference Joint Power Generation Conference, Miami Beach, Florida, USA, 2000, EPRI, USA. Nakano T, Tanaka K, Nakazawa T and Nishimoto S, ‘Development of large-capacitiy single-casing reheat steam turbines for single-shaft combined cycle plant’, Mitsubishi Heavy Industries Ltd, Technical Review, 2005, 42 (3), 23–29. Magoshi R, Tanaka Y, Nakano T, Konishi T, Nishimoto S, Shige T and Kadoya Y, ‘Development of welded rotors for high-temperature steam turbines’, International Conference Power Engineering (ICOPE), Chicago, Illinois, USA, 2005, EPRI, USA. Scarlin B, ‘Advanced high-efficiency turbines utilizing improved materials’, International Conference Advanced Steam Plant, IMechE Conference Transactions, 1997–2, Steam Plant Committee IMechE, UK. Roberts B W and Vanstone R W, ‘Materials for today’s fossil and nuclear steam turbines’, International Conference Steam Turbine Retrofit Conference, Chicago, USA, 2006, EPRI, USA. Zoerner W, ‘Steam turbines for power plants employing advanced steam conditions’, 10th International Conference CEPSI, Christchurch, New Zealand, 1994. Drosdziok A and Feldmueller A, ‘High-efficiency steam turbines for coal-fired power plants’, International Conference Power Gen Asia, Singapore, 1995. Haas H, Engelke W, Ewald J and Termuehlen H, ‘Turbines for advanced steam conditions’, International Conference American Power Conference, Chicago, Illinois, 1982, EPRI, USA. Haas H, Zimmermann A and Termuehlen H, ‘Turbines for advanced steam conditions – operational experience and development’, 1st International Conference Improved Coal-Fired Power Plants, Palo Alto, California, USA, 1986, EPRI, USA. Campbell C B, Frank C C and Sphar J C, ‘The Eddystone superpressure unit’, ASME paper 56-A-156, 1957.
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30 Silvestri G J Jr, ‘Eddystone station, 325MW generating unit 1 – A brief history’, American Society of Mechanical Engineers, USA, 2003, special issure, 30 pp. 31 Yugami H et al., ‘Operating experience of Wakamatsu high temperature turbine, Step 1’, Mitsubishi Juko Giho, 1990, 27 (1), 3–16. 32 Miyashita K, ‘Overview of advanced steam plant development in Japan’, International Conference Advanced Steam Plant, IMechE Conference Transactions, 1997, Steam Plant Committee IMechE, UK. 33 Muramatsu K, ‘Development of ultra-super critical plant in Japan’, International EPRI Conference Advanced Heat Resistant Steels for Power Generation, San Sebastian, Spain, 1998, EPRI, USA. 34 Fukuda M, Tsuda Y, Yamashita K, Shinozaki Y and Takahashi T, ‘Materials and design for advanced high temperature steam turbines’, 4th International Conference Advances in Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA. 35 Vanstone R W, ‘Advanced (700°C) pulverised fuel power plant’, 5th International Conference Charles Parsons Conference for Advanced Materials, Cambridge, UK, 2000, IoM Communication, UK. 36 Blum R and Hald J, ‘Benefit of advanced steam power plants’, 7th International Conference Materials for Advanced Power Engineering, Liege, Belgium 2002, Research Centre Juelich, Germany. 37 Blum R and Vanstone R W, ‘Materials development for boilers and steam turbines operating at 700°C’, 6th International Conference Charles Parsons Conference for Advanced Materials, Dublin, Ireland, 2003, IoM Communication, UK. 38 Blum R, Vanstone R W and Messelier-Gouze C, ‘Materials development for boilers and steam turbines operating at 700°C’, 4th International Conference Advanced Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004. EPRI, USA. 39 Scarlin B, Vanstone R and Gerdes R, ‘Materials development for ultra-supercritical steam turbines’, 4th International Conference Advances in Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA. 40 Kern T-U, Wieghardt K and Kirchner H, ‘Materials and design solutions for advanced steam power plants’, 4th International Conference Advanced Materials Technology for Fossil Power Plants, Hilton Head Island, South Carolina, USA, 2004, EPRI, USA. 41 Blum R and Vanstone R W, ‘Materials development for boilers and steam turbines operating at 700°C’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany. 42 Edelmann H, Effert M, Wieghardt K and Kirchner H, ‘The 700°C steam turbine power plant – status of development and outlook’, International J Energy Technol Policy, 2007, in print. 43 Tschaffon H, ‘The European way to 700°C coal fired power plant’, 8th International Conference Materials for Advanced Power Engineering, Liege, Belgium, 2006, Research Centre Juelich, Germany.
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22 Using creep-resistant steels in nuclear reactors S. K. A L B E R T, Indira Gandhi Centre for Atomic Research, India and S. S U N D A R E S A N, Maharaja Sayajirao University, Baroda, India
22.1
Introduction
Nuclear power plants differ from fossil power plants mainly in the source of heat for converting water into steam, which is subsequently used to run the turbine and produce electricity. In the former, the source of heat is nuclear fission (or fusion, in future fusion reactors), while in the latter, it is the burning of the fossil fuels like coal, oil or gas. In general, therefore, the structural materials chosen for nuclear reactors should also meet the requirements of fossil power plants in terms of good creep resistance, oxidation resistance, low-cycle fatigue strength, thermal conductivity, and so on. In addition to these, the elements present in the structural materials should also have a low neutron absorption cross-section, that is the probability of neutrons produced in the reactor being absorbed by these elements should be low. Further, the properties of these materials should not degrade under the high levels of radiation that exist in nuclear reactors. Such degradation is generally referred to as radiation damage and includes irradiation embrittlement, irradiation creep, swelling, helium embrittlement, and so on, which are described briefly later in the chapter. In a nuclear reactor, heat is produced by fission of the heavy elements like U or Pu, which is achieved by bombardment of the nuclei of these elements by neutrons. A typical fission reaction is given below: U92 + 1n0 →
235
94
Sr38 +
140
Xe54 + 2 1n0 + ≈160 MeV
Fission reactions produce, on an average, two or three neutrons per fission which can, in turn, take part in further fission, thus sustaining the reaction and generating energy continuously. However, not all neutrons would be available for the fission reaction: some of them would escape from the reactor and be permanently lost, others would be absorbed by the structural materials present, and still others would be simply scattered away by nuclei of various elements present in the nuclear fuel and the structural materials. It is thus important that absorption of neutrons by structural materials should 597 WPNL2204
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be minimized in order to sustain the fission reaction. It is in this context that the neutron absorption cross-section of various elements and their isotopes present in the structural materials becomes important in material selection for reactor core application.1 The core is that part of the reactor where the nuclear reaction takes place; the structural materials that contain the fuel and fission products and facilitate the removal of heat from the fission products are part of the reactor core. The unit of neutron absorption cross-section is the barn (1 barn = 10–28m2/nucleus)1 and the higher the cross-section the greater is the probability of neutron absorption. Minimizing the concentration of elements with a high absorption cross-section is an important criterion in material selection for the reactor core. The absorption cross-section is a function of the energy of the neutrons. In thermal nuclear reactors which use low-energy thermal neutrons (≤0.025 eV) to sustain the fission reactions, structural materials based on metals that have a low absorption cross-section in this energy range only can be used for the core applications. There are only a few metals like Zr, Al and Be that have absorption low enough1 to make them suitable for core application in the thermal reactors and among them only Zr-based alloys have the required mechanical properties. Average neutron absorption cross-section for iron in this energy range is high (2.55 barn compared with 0.19 barn for Zr) and hence ferrous alloys are not used for core application in thermal reactors. However, in the case of fast breeder reactors (FBRs), the energy of neutrons used for both fission and breeding is significantly higher and in this energy range the absorption cross-section for iron and most of the major alloying elements present in the steels is quite low.2 In these reactors, therefore, alloy steels are the main structural materials even for the core applications. Further, steels are also actively considered for core application in future fusion reactors.
22.2
Radiation damage
Radiation damage is a general term employed for material deterioration in a radiation environment and includes radiation swelling, irradiation creep, irradiation embrittlement, helium embrittlement, and so on. It occurs when high-energy particles displace atoms from their normal lattice sites to form Frenkel defects (vacancies and interstitials)3 and is usually expressed as displacements per atom or dpa, i.e. the number of times an atom is displaced during the irradiation. Additionally, neutrons can cause transmutation reactions with atoms of the irradiated material, resulting in solid or gaseous products. The former are generally considered harmless to material properties,4 but they could be highly radioactive isotopes with long half-lives. This aspect is relevant to the development of reduced-activation steels in which elements that can produce harmful radioactive isotopes in the reactor environment are either removed or maintained at very low levels. The gases produced are
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helium (through an (n,α) reaction) and hydrogen (through an (n,β) reaction). Small amounts of helium so produced can result in serious property deterioration which is known as helium embrittlement.
22.2.1 Radiation swelling Irradiation of some structural materials like austenitic stainless steels at temperatures in the range 0.3–0.5 Tm (Tm is the absolute melting temperature) and with intermediate to high neutron doses produces a significant density decrease and volume increase – a phenomenon known as radiation swelling. The point defects that form from displacement events are quite mobile at reactor temperatures and most vacancies and interstitials annihilate themselves by recombination. The surviving defects migrate to sinks and get absorbed. However, dislocations which are the major sinks for these defects have a slightly higher preference for interstitials and hence over a period of time excess vacancies appear in the material that migrate to form clusters. These clusters are stabilized as three-dimensional voids by innate gases in the material and/or transmuted gas such as helium produced by (n,α) reaction. Voids are not formed below 0.3 Tm because of dominant mutual recombination of the interstitials and the slow-moving vacancies and above 0.5 Tm because thermal vacancy concentration then exceeds that induced by irradiation.5 The relationship between neutron dose and void swelling in materials is of much practical interest. It is characterized by an incubation dose below which no swelling occurs. Above this, there is a linear increase in swelling with increase in dose.6 The extensive experimental data for austenitic stainless steels have shown that the magnitude of swelling depends on major and minor elemental contents and the initial thermomechanical treatment. These factors affect only the incubation or transient dose before swelling accelerates to a constant rate (~ 1% per dpa in austenitic stainless steel) independent of the irradiation temperature in the peak swelling range.7 Austenitic stainless steels, presently employed in the fabrication of reactor core components (clad tubes and wrapper) of FBRs, have poor resistance to radiation swelling. The composition of the conventional austenitic steels like AISI 316 and AISI 304 and the temperature range and radiation dose in FBRs are such that these steels experience considerable swelling in the operating environment of these reactors. Swelling in austenitic steels becomes unacceptable at dose rates higher than 50 dpa. However, a burn-up of 100 000 MW days/tonne corresponds to 85 dpa and the target burn-up required for FBRs to become economically viable is 200 000 MW days/ tonne, equivalent to fission of 20 atom% of the heavy metal present in the fuel. Hence the desired levels of burn-up cannot be achieved if clad tubes and wrapper made from conventional austenitic stainless steels are used in FBRs. This has led to the modification of both microstructure and composition
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of these alloys to improve the swelling resistance. Meanwhile, it was also found that steels with a ferritic/martensitic microstructure are much more resistant to void swelling than those with an austenitic microstructure and this resulted in the development of a large number of high-Cr ferritic steels for FBR core application. The basis for the development of austenitic stainless steels with improved resistance to void swelling has been the fact that void swelling can be reduced if the defects are pinned down by dislocations or solute atoms. As a result of these defect–dislocation or defect–solute atom interactions, coalescence of vacancies to form voids is delayed and recombination of interstitials and vacancies formed owing to neutron bombardment is facilitated. Accordingly, it was found that cold-worked 316 stainless steel exhibits much higher resistance to void swelling than the solution-annealed material. Further, increasing Ni above the levels present in AISI 316 steel and controlled additions of Ti and Si can also improve void swelling resistance in austenitic steels. Figure 22.1 shows the variation with dose of the maximum deformation in the diameter of a fuel pin irradiated in the Phenix reactor for various austenitic cladding materials.8 Later it was found that controlled addition of P and trace amounts of B can improve the creep resistance of this class of alloys. These findings led to the development of a new class of austenitic stainless steels with a base composition of 15Cr–15Ni with controlled additions of Ti and P in the USA,9 Europe8 and Japan10,11 for use in core components of FBRs. Many commercial and test reactors have clad tubes and wrapper materials made of this class of austenitic stainless steels subjected to an optimum level of 20%
8
Diametral deformation (%)
7 6 5 CW 316
4
CW 316 Ti
3
CW 15.15 Ti
2 1 0
CW Si-mod 15.15 Ti
40
60
80 100 Dose (dpa)
120
22.1 Variation in the maximum diametral deformation of the fuel pins of different austenitic steels irradiated in Phenix with neutron dose.8
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cold work. It has been found that a burn-up of 100 000 MWd/t can be achieved safely with core components of this class of alloys. For burn-up levels higher than the above, one has to depend on ferritic alloys which are much more resistant to void swelling than the austenitic alloys. Several explanations have been put forward for the superior swelling resistance of ferritic steels relative to austenitic steels. They include mechanisms that depend on solute trapping, the character of the dislocation loop structure, the lower relaxation volume of the body centred cubic (bcc) structure and the tempered martensitic structure.6 However, none of the proposed mechanisms completely explains the superior swelling resistance of the ferritic alloys3,5,12 and it is felt that a combination of lower dislocation bias (for solute atoms) in bcc alloys, high self-diffusion, low helium generation rates and high subgrain boundary sink strength might be contributing to the higher swelling resistance of ferritic steels in comparison with the austenitic alloys. Kim et al.13 compared swelling in type 316LN stainless steel (SS), 9Cr– 2WVTa steel and three oxide dispersion strengthened (ODS) steels using dual beam (3.2 MeV Fe+, 330 keV He+) irradiation to 50 dpa and 260 appm of He at 650°C. The ODS alloys studied were: Fe–17Cr–0.25Y2O3 (17Y3), Fe–12Cr–0.25Y2O3 (12Y1) and Fe–12Cr–3W–0.4Ti–0.25Y2O3 (12YWT). The microstructures of the steels were quite varied, with the SS having very clean grains and a low dislocation density (ρ<< 1013m m–3) and the 9Cr– 2WTa steel having large M23C6 precipitates and a dislocation density of ~1013 m m–3. The dislocation density in ODS alloys was typically in the range 1015–1016 m m–3 with 12YWT having a higher dislocation density than the others. Size and number densities of the oxide particles varied among the three ODS alloys. They were 3–5 nm and 1–2 × 1023 m–3, respectively, in 12YWT, 10–40 nm and 1021–1022 m–3 for 12Y1 and a larger particle size and lower number density in 17Y3 than in 12Y1. The results showed that a bimodal distribution of bubbles (r < 0.5 nm) and voids (0.5 nm ≤ r < 10 nm) developed in the alloys in decreasing order of 316LN SS, 9Cr–2WVTa, 17Y3 and 12Y1. In the 12YWT alloy, the bubbles formed were typically 1nm in size and associated with fine particles and a high density of dislocations with a number density of ~1023–1024m–3. These results indicated that the more complicated the microstructure, the less is the swelling. The density of available sinks for vacancies and interstitials (dislocations, precipitates, bubble, etc.) determines the recombination of these defects and thus the void formation. A change in dislocation bias caused by a change in microstructure could also be important in void formation. Microstructure affects the length of the transient stage (shift from incubation to steady state) of the swelling with irradiation dose. Once sufficient voids are nucleated, ferritic steels also swell and the steady state swelling rate is ~ 0.2%/dpa, which is around one-fifth of that observed in austenitic stainless steels.14
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The results of neutron irradiation studies conducted in different FBRs of various countries into swelling in ferritic steels have been reviewed by Klueh and Harries.6 Steels studied included 2.25Cr–1Mo, modified 9Cr–1Mo, HT9 (12Cr–1MoVW), EM12(9Cr–2MoVNb), EM10 (9Cr–1Mo), DIN 1.4914 (12Cr MoNiVNb), EP-450 (12Cr–1.5MoVNbB) and also some ODS alloys. The range of temperatures used for irradiation varied from 400–650°C in different reactors and these alloys were irradiated to different levels ranging from 30– 200 dpa. Among the alloys studied, swelling was a minimum in 12Cr steels with HT9 irradiated to 200 dpa at 420°C exhibiting a swelling of only 1.02% compared to 1.76% of 9Cr–1MoVNb steel. Steels with 9%Cr generally exhibited more swelling than those with 12%Cr, which is in agreement with the maximum swelling observed for 9%Cr binary alloys of iron and chromium. The degree of swelling was particularly high in the EM12 alloy, which was attributed to its duplex microstructure and the presence of δ-ferrite. Further, the swelling resistance of ODS alloys was found to be comparable with that of conventional alloys. Results also indicated that prior microstructure or heat treatment history also influenced the swelling behaviour of these alloys. The fusion reactor irradiation environment differs from that of a fission reactor in the high-energy component in the neutron spectrum. Neutrons with energy up to about 14.1 MeV are present in a fusion reactor, while the neutron energies in a fission reactor are lower than 2 MeV. The higher energies in a fusion plant imply a higher cross-section for the (n,α) reaction with elements like Fe, Cr and Ni. Significantly greater amounts of helium and atomic hydrogen will form in steels exposed to the fusion environment, a consideration that is important in developing steels for the first wall and blanket structures in a fusion power plant.
22.2.2 Irradiation creep Discovered in the 1950s, first in uranium and a few years later in stainless steel, irradiation creep refers to the strain undergone by a material when irradiated under stress. However, in the absence of either stress or irradiation the creep becomes insignificant.15 Thermal creep becomes prominent for irradiation at temperatures ≥ 0.5 Tm, while irradiation creep can be significant at much lower temperatures. Similar to the case of thermal creep, dislocation climb and glide play an important role in the deformation processes that occur during irradiation creep.16 Much of the theoretical work on irradiation creep has been concerned with steady-state point defect concentrations. More recently, it has been realised that large contributions to creep can occur as a result of several types of transients in point defect populations.15 The point defects generated by irradiation can be annihilated through stress-induced processes, for example by absorption on dislocations. If the absorption is asymmetric or preferential,
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it can cause the dislocations to climb, which can subsequently lead to glide of the dislocations. It must be noted that, if vacancies and interstitials were absorbed equally, annihilation would occur without climb, and there would be no creep.16 The presence of an applied stress can change the capture efficiencies of oriented dislocations for the irradiation-induced point defects. In particular, there is a slight bias for self-interstitials to be absorbed by dislocations whose Burgers vectors are more nearly aligned with the stress axis. This causes dislocation climb and the different rates of climb for dislocations with Burgers vectors at different orientations to the applied stress can result in climb creep,17 which has been termed stress-induced preferential absorption (SIPA) creep. When gliding dislocations become pinned at obstacles, the bias-driven interstitials at dislocations can cause them to climb around the obstacles by the SIPA process, after which the dislocations can glide again under the applied stress until they are pinned by another obstacle. This results in an increment of strain, referred to as preferred absorption glide (PAG) creep. The climb-enabled glide mechanism can also operate as a result of swelling: as cavities grow, dislocations absorb excess interstitials and climb.17 Swellingdriven creep has also been called I-creep. The various mechanisms of irradiation creep have been reviewed by Mathews and Finnis.18 Irradiation creep can help accommodate stresses induced by differential swelling. However, excessive irradiation creep can lead to buckling and the stress relaxation can render devices such as clamps and bolts ineffective.15 The initial experimental studies on irradiation creep of ferritic/martensitic steels showed that at temperatures of the order of 300°C, at which thermal creep was negligible, considerable irradiation creep could occur. However, the steady-state irradiation creep rate for a 12Cr–MoVNb ferritic/martensitic steel was nearly an order of magnitude lower than for austenitic stainless steels.16 A study of the behaviour of the 12Cr–MoVW steel (HT9) irradiated at temperatures of 540°C and 595°C19 showed that this steel is superior to several austenitic stainless steels, although precipitation hardened Ni-base alloys were significantly more creep resistant. It should be remembered, however, that at these temperatures thermal creep could also occur in addition to irradiation creep. The creep behaviour of an in-reactor tested specimen in the HT9 steel was compared by Chin20 with that of a thermal control specimen of the same material, as illustrated in in Fig. 22.2. While the steady-state creep rates of both the specimens are nearly the same, the difference in total strain arises from the existence of a primary creep stage in the thermal control specimen; irradiation appears to eliminate any detectable primary creep region. Figure 22.320 show the variation in the creep coefficient B (which relates creep strain to creep stress) as a function of temperature. The results illustrate that
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0 0.9
2
0.8
HT-9 540°C 50 MPa
0.7
4
6
ε (%)
0.6
Time (103 h) 8 10 12
14
16
18 20
10
11 12
Thermal
0.5 0.4 0.3 In-reactor
0.2 0.1 0.0 0
1
2
3 4 5 6 7 8 9 Fluence (1022 neutrons cm2)
B (10–28 cm2/neutron × MPan)
22.2 Comparison of irradiation creep with thermal creep for HT9 alloy.20
10 HT-9 STA 8
Fluence (1022 n/cm2)
6
4.3 6.7 10.2
4 2 0 350
450
550 650 Temperature (°C)
750
22.3 Variation of creep coefficient B with temperature for HT9 alloy.20
in-reactor creep is insensitive to temperature changes below about 500°C, that is, under conditions in which thermal creep makes an insignificant contribution. Above 550°C, thermal creep becomes prominent and, as the creep strengthening mechanisms rapidly deteriorate, large creep rates are observed.20 While a transient creep stage was not found in Chin’s in-reactor creep tests, the results obtained with 12 Cr–MoVNb steel by Wassilew et al.21 did show a transient stage. This difference has been attributed to the fact that the fluence range in the latter tests (using uniaxial specimens) was well below that in Chin’s pressurized-tube studies. Apparently, the first measurements in the pressurized-tube tests were made after the primary stage was completed.16
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Irradiation creep studies on several 9Cr–Mo ferritic/martensitic steels have demonstrated their superiority to austenitic stainless steels such as Type 316 and 15-15Ti, especially at temperatures up to about 500°C.22,23 Later work comparing the behaviour of HT9 and T91 showed that modified 9Cr–1Mo steel underwent even lower irradiation creep than 12Cr–1MoVW.24 Klueh and Harries16 have estimated the irradiation creep coefficients (expressed in units of MPa–1dpa–1) of a number of Cr–Mo steels and the result shows the essential similarity in behaviour of this class of ferritic/martensitic steels. The coefficient falls in the range of 1.25–10 × 10–7 MPa–1dpa–1 for the different Cr–Mo steels that were examined. It is of some interest to compare the performance of high-Cr ferritic/ martensitic steels with that of ODS ferritic steels. The latter are candidate materials for fuel cladding for fast reactors. Comparative tests under similar stress, temperature and irradiation conditions have shown that the creep resistance of the ODS steels was superior by a factor of 2–5 times. The Cr– Mo steels, however, performed 3–4 times better than Type 316 austenitic stainless steel.16 If the out-of-pile creep component is subtracted from the inpile creep component, the true irradiation creep behaviour is revealed. Such an analysis shows that, in conventional Cr–Mo steels, the total creep increases continuously with increasing temperature, since above 450°C thermal creep is dominant. With the ODS steels, on the other hand, irradiation creep rate decreases above 450°C and, since thermal creep is also low in these steels, the total creep rate actually decreases with increasing temperature above 450°C. The superior irradiation creep behaviour of the ODS steels might be based on the presence of a large number of dispersed particle interfaces which act as sinks for the point defects and thus reduce the fraction going to the dislocations to cause irradiation creep.16
22.2.3 Irradiation embrittlement An area of important concern in the use of ferritic/martensitic steels in light water reactors, fast reactors and future fusion reactors is the adverse effect of irradiation on fracture toughness.25 Irradiation is known to cause a large increase in the ductile-to-brittle transition temperature (DBTT) and a significant reduction in the upper shelf energy (USE). The practical implication is that the post-irradiation DBTT may rise above ambient temperature, although it might have been well below room temperature prior to irradiation. Irradiation embrittlement of ferritic steels is related to the lattice hardening caused by precipitation, induced or accelerated by the irradiation. These effects arise from irradiation below about 0.4 Tm, where Tm is the absolute melting temperature of the steel (Tm ≈ 1500°C). Hardening raises the flow stress and, if the fracture stress is assumed to be unaffected by irradiation, the DBTT is increased.25
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On account of their relatively superior void swelling resistance, the ferritic/ martensitic steels are potential candidate materials for the first wall and blanket structures of future magnetic fusion reactors. The first wall of a fusion reactor will experience displacement damage from the high-energy neutrons of the fusion reaction. Additionally, large amounts of transmutation helium will form: for one year of irradiation of a martensitic steel, about 110 appm (1 appm = 1 atom per million) of He is produced in a fusion reactor first wall compared with < 10 appm He in a fast reactor.26 Thus, the irradiation conditions expected in a fusion reactor cannot be simulated in a fast reactor. Investigators have therefore attempted to produce simultaneously displacement damage and transmutational helium by adding nickel to the steel under study and irradiating it in a mixed-spectrum reactor, such as the high-flux isotope reactor (HFIR). The fast neutrons in the spectrum cause displacement damage, while the thermal neutrons produce helium from 58Ni in a two-step transmutation reaction 58Ni (n, γ) 59Ni followed by 59Ni (n, α) 56Fe. This technique has been used to investigate the effect of helium on mechanical properties and swelling resistance.26 In the study of irradiation effects on impact behaviour, the limited irradiation space available in most of the fission test reactors has dictated the use of small-size Charpy specimens 1/2 and 1/3 the standard size.27 Such tests have shown that the shifts in DBTT have been greatest at the lowest irradiation temperature, which is usually below 400°C.28 Below 400–500°C, the magnitude of the shift caused by displacement damage varies inversely with irradiation temperature.25 Most irradiation embrittlement studies have been conducted in the temperature range 360–600°C. Although the Charpy curves are shifted by irradiation hardening, the fracture mode is generally unaltered between the unirradiated and irradiated steels: cleavage or quasi-cleavage on the lower shelf and ductile void coalescence on the upper shelf.25 However, in the situation where helium is produced by transmutation, features associated with intergranular fracture have been observed. This aspect is discussed subsequently. Investigations of high-chromium ferritic steels in fast-neutron environments have indicated that the shifts in DBTT tend to saturate with increase in fluence, as illustrated in Fig. 22.4.27 The hardening caused by irradiation also tends to saturate in a similar manner. The saturation effect is, however, not observed in irradiation studies in mixed-spectrum reactors. This has been attributed to the larger amount of helium generated by irradiation in environments containing thermal neutrons. Only small quantities of helium are produced in fast reactors, whereas considerably larger amounts are formed in mixed-neutron environments.27 Similarly, a study using spallation neutron sources has also demonstrated the absence of the saturation effect in T91 steel.29 The data showed that the DBTT increases continuously with dosage in the range investigated (0–9.4 dpa) as shown in Fig. 22.5.29 At the highest
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40 12 Cr-1 MoVW Steel
Energy (J)
30
Normalized and tempered 10 dpa, 365°C, FFTF 17 dpa, 365°C, FFTF
20
10
0 –200
–100
0 100 Temperature (°C)
200
300
22.4 Charpy curves for half-size specimens of 12Cr–1MoVW steel in the normalized and tempered condition and after irradiation to 10 and 17 dpa at 365°C in FFTF.27
0.25 –30
DBTTSP (°C)
0.15
–90
0.10
–120 DBTTSP USESP
–150
–190
0
2
4 6 Displacement (dpa)
8
USESP (J)
0.20
–60
0.05
10
0.00
22.5 Irradiation dose dependence of the DBTT from small punch (SP) test.29 USE is the upper-shelf energy.
dosage of 9.4 dpa, the increase in DBTT was as high as 295°C, when the shift was recalculated to standard Charpy conditions. The large increase in DBTT was also associated with intergranular fracture. These effects have also been ascribed to the large quantities of helium produced during irradiation
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with spallation sources. In fact, the increase in DBTT showed a linear dependence on helium content, see Fig. 22.6.29 The microstructural changes caused by irradiation occur primarily over the range 300–500°C and almost no change occurs at 600°C.30 Depending upon the temperature of irradiation, some recovery of the as-tempered, cellular subgrain structure might occur, as during long-term annealing. However, of greater relevance to the embrittlement process is the production of dislocation loops from the irradiation-induced point defects and, at temperatures of 400°C and above, the formation of voids and micropores (which are presumably helium bubbles) along prior-austenite and lath boundaries.21,30 No new precipitation was observed in T91 after irradiation, but there were changes in the as-tempered carbide distribution and elemental concentration.30 However, in several other Cr–Mo steels, extensive precipitation has been observed over a range of irradiation temperatures, for example, the G phase observed in HT-9 and identified to be a nickel silicide, α′ and Mo2C.30 Such precipitation has been considered to be related to the observed delays in the onset of swelling in irradiated ferritic alloys, although it is also held to be responsible for irradiation hardening and embrittlement.31 These microstructural changes result in hardening and embrittlement depending upon the temperature of irradiation. Investigations of several Cr– Mo steels have demonstrated that the degree of embrittlement is considerable and is a matter for serious concern. Irradiation of 12Cr–1MoVW steel in a fast-flux test facility (FFTF) to 10 and 17 dpa at 365°C confirmed the saturation effect under fast reactor irradiation. The same investigation also showed that the increase in DBTT in 2.25Cr–1Mo steel after irradiation under identical conditions in the FFTF was similar in magnitude to the increase observed for
125
∆DBTTSP (°C)
100
75 Slope = 0.135°C/appm He 50
25 100
200
300 400 500 600 Helium content (appm)
700
800
22.6 Helium content dependence of ∆DBTT on T91 steel obtained in small punch test.29
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12Cr–1MoVW steel. However the final DBTT for the 2.25 Cr–1Mo steel was lower than that for 12Cr–1MoVW because of the lower starting DBTT for the 2.25 Cr–1Mo material.27 Investigations of 12Cr–1MoVW and 9Cr–1MoVNb irradiated in a fast reactor have revealed that, while the DBTT shifts tend to saturate with increasing fluence, the hardening and embrittlement are significantly decreased at irradiation temperatures of 450°C and above.32 Klueh and Alexander irradiated specimens of 9Cr–1MoVNb (T91) and 12Cr–1MoVW (HT9) steels with and without 2% Ni addition in the mixedspectrum high-flux isotope reactor (HFIR) at 300 and 400°C.33 While no saturation was observed up to a displacement of ~ 42 dpa, the shifts in DBTT measured after irradiation at 400°C were the largest ever observed for these steels. The Ni-added steels exhibited a greater effect than those without nickel enhancement. Also, the shifts were much larger after irradiation in a mixed-spectrum reactor than when the same steels were irradiated in a fast reactor, where little helium formed during irradiation.32 Many attempts have been made to study the possibility of improving preirradiation toughness by optimizing prior heat treatment, minor impurity concentration and fabrication processes without significantly affecting the strength, swelling resistance and creep resistance. Work on 9Cr–1MoVNb steel has demonstrated that heat treatment can be optimized to reduce prioraustenite grain size and lower the pre-irradiation DBTT.34 The shift in DBTT owing to irradiation was, however, not affected by the prior heat treatment. This means that such optimized heat treatment can be used to ameliorate the effect of irradiation on DBTT in 9Cr–1MoVNb steel. In the case of 12Cr– 1MoVW steel, however, the same investigation showed that austenitization temperature and austenite grain size had little effect on the unirradiated DBTT. The implication is that prior heat treatment cannot be used as a means to offset the DBTT shift caused by irradiation in 12Cr–1MoVW. It is of interest to mention here that in the same steel the use of a relatively high austenitization temperature (to dissolve all carbides and δ-ferrite present) in conjunction with a moderate temperature age at 700°C has been reported to cause a 50°C improvement in the pre-irradiation DBTT.35 The embrittlement caused by irradiation has also been studied in the case of weldments. The results, however, have not been consistent. Hu and Gelles28 found with 12Cr–1MoVW steel that the weld and heat-affected zone samples exhibited the same or slightly lower DBTT and similar upper shelf energy (USE) as the base metal. This would mean that ferritic/ martentisitic steels are not expected to suffer any penalty in performance caused by the presence of welds. However, the same investigators found in a subsequent study on 9Cr–1Mo steel that the DBTT for the weldment was about 60°C higher than for the base metal exposed to the same irradiation conditions.32
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During the development of reduced-activation steels, systematic studies of the effects of various elements were made. An important result was that the boron content seemed to be a controlling factor for DBTT. The 10B isotope (20% of the natural composition) is a strong thermal neutron absorber and transmutes to He and Li even at moderate neutron fluxes. Helium bubbles have been detected on the crack surfaces of irradiated material. The data showed that the higher the boron concentration the steeper was the slope of the ∆DBTT versus irradiation dose curve. The boron effect overrode all other compositional factors.36 Another study with reduced-activation steels showed that, while the presence of boron generated He during irradiation in all the steels investigated, significant shifts in DBTT resulted only in some cases.36 In general, such studies demonstrate that at least with low neutron fluxes, the advanced reduced-activation alloys provide much improved impact properties after irradiation.37
22.2.4 Helium embrittlement In addition to the irradiation hardening and embrittlement taking place at irradiation temperatures below about 0.4 Tm , another form of embrittlement is exhibited at elevated temperatures higher than about 0.5 Tm following irradiation at ambient (or elevated) temperature. Such high-temperature embrittlement is manifested by austenitic stainless steels and some other materials and has been attributed to the presence of helium gas generated during irradiation. The subject has been comprehensively reviewed by Klueh and Harries.38 The production of helium is attributed to nuclear reactions involving B and Ni bombarded by thermal neutrons and Ni, Cr and Fe by fast neutrons in the reactor environment. Helium bubbles forming at grain boundaries of the irradiated austenitic stainless steels are believed to nucleate cavities that grow under stress and coalesce to form intergranular cracks in a manner analogous to cavity growth during thermal creep. Although several solutions have been suggested for mitigating the problem, such as grain refinement, trapping of helium at particles inside the grains and promoting grain boundary precipitation to reduce the production of vacancies, a fully satisfactory answer to the problem has yet to be found.38 In addition to the austenitic stainless steels, ductility reductions in elevatedtemperature tensile testing have been observed in several nickel-base alloys, a niobium alloy and, to a lesser extent, in ferritic stainless steels. In contrast, however, the high-chromium Cr–Mo ferritic/martensitic steels have been found to be much more resistant to this form of embrittlement. In a relatively early investigation,39 nickel-doped 9Cr–1MoVNb and 12Cr– 1MoVW steels were irradiated in a mixed–spectrum reactor, which generated up to 49 appm of helium, but subsequent tensile testing at 700°C showed no
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detrimental effects of irradiation or helium concentration on ductility. On the other hand, cold-worked Type 316 austenitic stainless steel exhibited a significant reduction in ductility on similar irradiation and testing. Helium implantation in a compact cyclotron with a beam of α-particles has enabled helium concentrations of up to 3000 appm to be developed in a high-Cr ferritic/ martensitic steel corresponding to DIN 1.4914.40 However, tensile testing at temperatures up to 750°C revealed no helium embrittlement. In-beam creep tests at 600 and 700°C also showed no embrittlement, the fracture mode remaining transgranular except that, at the highest helium concentration and test temperature, portions of mixed trans- and intergranular fracture were observed. Helium implantation studies on 9Cr–0.5V and 9Cr– 2W reduced-activation steels confirmed the resistance of high-Cr martensitic steels to elevated temperature embrittlement, the tensile fracture always being ductile and transgranular.41 The relative immunity of ferritic/martensitic steels to such embrittlement is believed to be related to their lath microstructure containing fine carbide precipitates in the ferrite matrix and on the low-angle lath boundaries. It is postulated that such a structure enables the partitioning of helium atoms to the lath and precipitate boundaries, so that the helium concentration on any high-angle grain boundaries is relatively low.39 Microstructural evidence supporting the above reasoning has been provided by subsequent investigations on high-Cr martensitic steels, which showed helium bubbles forming on the lath boundaries, sub-boundaries and dislocations in the laths, but no preferential bubble growth at prior-austenite grain boundaries that could cause intergranular fracture.40,41
22.3
Embrittlement caused by ageing
An important limitation of high-Cr ferritic/martensitic steels is their marginal fracture toughness and their transition from ductile to brittle behaviour at a temperature not much below 0°C. A matter for further concern to that the toughness tends to deteriorate during long-term exposure to elevated temperature and/or aggressive environments. Since the presence of impurities can adversely affect toughness, it is necessary to impose tight specifications to control impurity concentration. Investigations of the impact toughness of 9Cr–1Mo steels have shown that phosphorus and silicon raise the DBTT and decrease the upper-shelf energy (USE) as a consequence of segregation effects and increase in the deltaferrite content, respectively. The effect of sulphur is more pronounced than that of phosphorus and has been attributed to the formation of non-metallic inclusions which reduce toughness.42 Several studies have shown that in 9Cr–1Mo steels a loss of room temperature ductility and toughness and an increase in DBTT can occur as
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a result of thermal exposure in the range 500–600°C. Much of the mechanical property deterioration that takes place during long-term thermal ageing can be related to microstructural changes arising from such exposure. The pretempered material exhibits a tempered martensitic microstructure consisting of lath-shaped ferrite subgrains within prior-austenite grain boundaries, the subgrains containing a concentration of dislocation networks. Precipitation is generally distributed along grain and subgrain boundaries, consisting of coarse M23C6-type and finer MC-type carbides. Microstructural changes observed during long-term ageing (typically, up to 25 000 h) of 9Cr–1MoVNb steel can be separated distinctly into two temperature regimes of quite different ageing effects.30,43 At lower ageing temperatures in the range 482–593°C, little change occurs in lath/boundary and carbide precipitate structures, except for a transient increase in dislocation density within the subgrains. However, the more important effect of lower temperature ageing is the formation of significant amounts of the Fe2Motype Laves phase, which occurs to a great extent in steels containing a higher amount of silicon.30,43 The Laves phase forms preferentially at the lath and grain boundaries. Abundant, fine vanadium carbide (VC) needles also form at temperatures of 593°C and below. In contrast, during ageing at temperatures in the range 593–704°C, there is detectable recovery and coarsening of the lath/subgrain boundary structure into equiaxed grains and coarsening of the carbide precipitates, but new precipitation seems to be much reduced, especially in low-silicon steels.43 The DBTT increases sharply and USE decreases on ageing at temperatures up to 593°C, reaching a maximum and a minimum, respectively, after 25 000 h of thermal exposure. It is considered very likely that precipitation, especially of the intermetallic compound Laves phase, is responsible for the property deterioration, although the exact mechanism has not been identified. The amount of precipitates and relative amount of the Laves phase reach a maximum after ageing at 538°C, corresponding well with the maximum shift in DBTT. At ageing temperatures of 650°C and above, structural coarsening softens the material and the USE increases rapidly. In low-alloy steels, a process of temper embrittlement is known to occur by the segregation of residual impurity elements P, Sn and Sb to the prioraustenite grain boundaries. High-Cr ferritic/martensitic steels also exhibit this type of embrittlement on ageing, especially during short-term ageing.44,45 It has been shown that, in a high-phosphorus (>0.017%) 9Cr–1Mo steel, ageing at 550°C for 1000 h leads to significant reduction in ductility and increase in DBTT, which is believed to be associated with the segregation of phosphorus to the carbide–matrix interfaces at which de-cohesion occurs. For longer ageing durations of 5000 h and above, the formation of the Laves phase becomes the dominant effect, leading to shifts in the DBTT of more than 60°C.44 Even in such cases, it has been proposed that phosphorus could
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play a secondary role by being absorbed in the Laves phase and further embrittling it.45 Since conditions promoting ageing embrittlement and hydrogen embrittlement may co-exist during the operation of a fusion reactor, the synergism between the two has been studied in a 9Cr–1Mo steel.45 This investigation showed that a strong interaction existed between hydrogen embrittlement and temper embrittlement. Both acting together resulted in a greater degree of embrittlement, in which the dominant factor controlling the interaction was the precipitation of the Laves phase containing phosphorus in solution.
22.4
Use of heat-resistant steels in major reactor types
Thermal reactors, which use low-energy thermal neutrons to sustain nuclear fission, are the most widely used reactors in nuclear energy production. There are different variants of the thermal reactors depending on fuel, moderator (used to slow down the high-energy neutrons produced in the reactor) and the fluid used for the removal of the heat. They include boiling water reactors (BWR), pressurized water reactors (PWR), pressurized heavy water reactors (PHWR), Russian VVER type reactors. In all these reactors, the core structural materials (used to contain fuel and extract heat from the fuel) are Zr-based alloys. However, the reactor vessel, heat transport system, steam generators, and so on employ different classes of steels. As the steam temperature is typically below 300°C, creep-resistant steels are not extensively used in the major components of these reactors. These are employed only for those components dealing with steam, like steam piping and turbines. Austenitic stainless steels find application as piping material, especially in BWRs, because of their corrosion resistance. In PHWRs, austenitic stainless steel has replaced carbon steel in recent reactors as the material of construction for calendria, a major reactor component supporting the coolant tube, because of the embrittlement of the latter in service. High temperature gas cooled reactors (HTGR) are a special case of thermal reactors that operate at relatively high temperatures. They use gases like helium or carbon dioxide for heat transfer and graphite to moderate the neutron energy. In one HTGR type known as advanced gas cooled reactors (AGR), uranium oxide clad with stainless steel is used as fuel. As the coolant temperature could be as high as 687°C, superalloys are used for the structural components of the reactor. Steam temperatures are lower at ~540°C.46
22.4.1 Fast breeder reactors Fast breeder reactors use fast neutrons to sustain the fission reaction as well as breeding. The non-fissile isotopes 238U and 232Th are converted to fissile
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isotopes of 239Pu and 233U, respectively, thus producing fresh fuel during reactor operation. Use of such reactors as a source of energy is being actively pursued in India, Russia, France and Japan and these reactors are expected to play a major role in the revival of the nuclear industry worldwide. Since hydrogen nuclei present in water, the normally used medium to extract heat, absorb and slow down the neutrons, water or any hydrogen-containing fluids is not suitable for extracting heat from these reactors. Hence, liquid metals (predominantly sodium) are used for extraction of heat produced from fission. Subsequently, in the secondary side of the reactor, liquid metal transfers the heat to water to produce steam which finally drives the turbine. Hence, compatibility with liquid metals is a major consideration in the selection of the structural materials for this class of reactors. A heat transport flow sheet for a prototype fast breeder reactor (PFBR) that is under construction is shown in Fig. 22.7.47 Fast breeder reactors (FBRs) use heat-resistant steels extensively both in the reactor core as well as in the conventional steam side of the reactor. As already mentioned, in the fast neutron spectrum, the neutron absorption cross-sections for Fe and other alloying elements are very small and hence Surge Tank 2 Loops
613 K 3.2MPa 753 K 17.2MPa
4 loops 6 LMW
795 K
Air
Steam generator Secondary sodium pump
Sodium
628
HP
IP
LP
500 MWt
Generator
TR
0.01MPa
Reheater
Turbine
To feed water heaters Condenser
820 K
CEP
Core PSP 670 K
SEA
301 K – 305 K
Deaerator
HP heaters IHX BFP
LP heaters
22.7 Schematic heat transport flow sheet of India’s PFBR. PSP – Primary Sodium pump; IHX – Intermediate heat exchanger; HP – High pressure; LP – Low pressure; IP – Intermediate pressure; TR – Transmission; BFP – Boiler feed pump; CEP – condenser extraction pump.
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neutron absorption is not a concern in the selection of the materials. Further, the normal operating temperature range for these reactors is typically ~ 600°C – a range that is ideally suited to the use of heat-resistant (both austenitic stainless and high-Cr ferritic) steels. Regarding the choice of materials, the major reactor components of FBRs can be classified into core components, structural components like the reactor vessel, safety vessels, piping, etc. and steam generators. Most of the FBRs for which construction is completed and those which are under construction use austenitic stainless steels for the first two sets of components. However, for steam generators, ferritic steels are more appropriate. In the immediate future, ferritic steels might even be chosen even for core components.
22.4.2 Austenitic steels in FBRs Core components Although, for the reactors currently in operation and under construction, both core components and other reactor parts, except steam generators, are made predominantly from austenitic stainless steel, the steels chosen for core components differ from those chosen for other components. Clad tubes which contain the fuel and the sub-assembly wrapper which houses the clad tube are the major components of the reactor core. Figure 22.8 shows a schematic of a fuel sub-assembly and clad tube for India’s PFBR programme. These components are exposed to intense neutron irradiation which can in turn result in serious problems like void swelling, helium embrittlement and irradiation creep. Deformation of the sub-assembly components occurs owing to void swelling and creep induced by irradiation and internal sodium pressure. In fact, dimensional changes associated with void swelling limit the resident time of these components in the reactor. The fuel clad tubes experience temperatures in the range of 400–700°C under steady-state conditions. The wrapper tubes operate in the range 400–600°C. Besides adequate tensile strength, ductility and creep strength, compatibility with liquid sodium which is used as the coolant is also an important material property requirement. The first generation materials for fast reactor core components belonged to Type 304 and Type 316 austenitic stainless steels, especially the latter.47 These steels were soon found to be inadequate because of unacceptable swelling at doses higher than 50 dpa and the first attempt was to introduce cold work to improve the resistance to swelling. Subsequently, the alloy composition was modified, as already reported, by increasing the nickel content, lowering the chromium content, raising the phosphorus level and adding titanium. It was found that there is a synergistic effect between titanium and carbon and this has been related to the amount of free carbon available for trapping vacancies, the effect being smallest if titanium is totally absent
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Head
7 shield pins, OD 36±0.15, ID 33.4,±0.15 pellet stack 03110.1
Steel V shielding
300
V
5
B4C shielding
200
Width A/F 131.3
Adaptor
25
616
Section-MM Top blanket
X
Bottom blanket
2580
X
E
4500
E
CD
Fuel
CD 1000 (Active core)
210 FUEL PMS
Section-EE
300
Clad OD-LB10.00 0.4710.02
10
Spacer wire 51.4510.00 Pallet OD-LB10.00 0.1310.3
Section-BB
H
30
608
H
710
Coolant entry tube
Section-HH
Clad tube
Discriminator Section XX Fuel subassembly
22.8 Schematic of fuel sub-assembly and clad tube for a prototype fast breeder reactor.
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or if all the carbon atoms are bound by the titanium. On this basis, the need to maintain the optimum Ti/C ratio was identified.48 The beneficial effect of phosphorus, which delays significantly the irradiation-induced swelling and creep, is attributed to several possible mechanisms: strong binding between phosphorus in solution and vacancies, phosphide precipitates acting as traps for helium atoms, and so on.49 A small addition of boron is made in some of these grades to improve creep ductility.47 The composition of some of the commercially available austenitic stainless steels developed for FBR core components is given in Table 22.1. It is necessary to know how the addition of elements such as phosphorus and titanium, done to increase the resistance to irradiation effects, influences the creep properties, both in the reactor environment and outside. Phosphorus additions to Type 316 SS have been shown to increase the thermal creep rupture strength at 650°C, but not at 750°C.50,51 Similarly, titanium addition to a 15 Cr–15 Ni stainless steel was found to improve thermal creep behaviour, the effect reaching a maximum in the range 0.2–0.3%Ti. However, silicon is reported to exercise an adverse effect on the creep performance of coldworked 15Cr–15Ni–Mo–Ti steel owing to accelerated precipitate coarsening effects. Another systematic study on the influence of Ti, Si, and P on thermal creep properties has been carried out during the development of D9I in the USA.49 This mainly involved adding controlled amounts of phosphorus and boron to a base D9 composition and optimizing their levels. The investigation clearly showed that both these elements increased rupture life at a test temperature of 700°C. The creep ductility was also found to exceed 18% in all the compositions studied. Based on the results of this work, the chemical composition of D9 was modified and the new composition was designated as D9I, as given in Table 22.1. Long-term biaxial creep rupture tests performed in the temperature range 650–760°C have confirmed the superior creep behaviour of D9I in relation to D9. The in-reactor creep strength of D9 and D9I is compared in Fig. 22.9.52 It can be seen that the improvement in creep rupture life observed in out-of-reactor tests with controlled additions of phosphorus and boron is also maintained during irradiation. Structural components Austenitic stainless steels with composition slightly modified from the standard AISI 304 or 316 steels are the major structural materials for most FBR components (other than core components and steam generators). The temperatures during operation of these components range from 400°C to 550°C and the environment is liquid sodium or argon with sodium vapour or nitrogen gas depending on the location of the component. Except for the regions near the core, irradiation is not a matter of serious concern. These
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Alloy
Cr
Ni
Mo
Mn
Si
C
Ti
Type 316
18.0
13.0
2.6
1.9
0.80
0.05
0.05
P
Others
Fe Balance
D9 (USA)
13.8
15.2
1.5
1.7
0.9
0.052
0.23
0.003
–
D9I(USA)
13.5
15.0
1.8
1.9
0.8
0.043
0.26
0.030
0.005 B
Balance
15–15Ti (France)
14.7
14.7
1.2
1.6
0.43
0.096
0.43
0.007
–
Balance
Balance
Mod. 15–15Ti (France)
14.9
14.8
1.5
1.5
0.9
0.085
0.50
0.007
–
Balance
PNC 316 (Japan)
16.0
14.0
2.5
1.8
0.8
0.055
0.10
0.028
0.08 Nb
Balance
PNC 1520 (Japan)
15.0
20.0
2.5
1.9
0.8
0.06
0.25
0.025
0.11 Nb
Balance
DIN 1.4970
15.1
15.1
1.3
1.3
0.5
0.088
0.48
0.004
–
Balance
D9 (India)
14.0
15.0
2.0
2.0
0.4
0.04
6×%C
0.02
0.05 Nb
Balance
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Table 22.1 Nominal chemical compositions of some currently used austenitic steels for FBR core applications
Using creep-resistant steels in nuclear reactors
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Hoop stress (MPa)
1000
100
D9 575°C 605°C 670°C 750°C 10 12
D91
D91 575°C 630°C 695°C 775°C 14
D9
16 18 LMP, T (13.5 + log tp) × 10–3 (K-h)
20
22.9 Comparison of in-reactor creep strength of D9 and D9I alloys.52 LMP is the Larson–Miller Parameter, temperature in Kelvin and time in h.
steels are chosen because of their high-temperature properties, compatibility with liquid sodium, creep strength, weldability, availability of design data and long experience in their use of these steels in the application temperature range in other industries. Low-cycle fatigue and creep-fatigue interactions are important considerations in the selection of these materials. Reactor designers prefer monometallic construction for liquid sodium systems for fear of migration of interstitial elements like carbon through liquid sodium owing to differences in thermodynamic activity in bimetallic systems. Austenitic stainless steels are thus used in the entire liquid sodium system even if the temperatures of some components are low enough to permit the use of less expensive ferritic steels.47 Austenitic stainless steels like 304, 316, 321, 347, 304L and 316L have been used in various reactor components in the reactors built in different countries in the early stages of their FBR programmes, but the current trend is to use low-carbon, nitrogen-added 316 or 304 generally designated as 316L(N) or 304L(N). Low-carbon varieties are chosen to minimize the risk of sensitization. Nitrogen is added to compensate for the loss of strength caused by reduction in carbon content. Compositions of different 316L(N) steels considered for the European Fast Reactor (EFR), the Demonstration Fast Breeder Reactor (DFBR), the Superphenix and Prototype Fast Breeder Reactor (PFBR) are given in Table 22.2. These steels have been extensively studied for their properties and compatibility with liquid sodium. Figures 22.10 and 22.11 show a comparison of creep rupture strength and fatigue behaviour, respectively, of various 316L(N) steels.47,53,54
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Table 22.2 Chemical composition of major austenitic stainless steel structural materials used in different FBRs (wt%)47 Element
316L(N) SS (EFR)
316FR (DFBR)
316L(N ) SS (Superphenix)
316L(N) SS PFBR
C Cr Ni Mo N Mn Si P S Ti Nb Cu Co B (ppm) Nb+Ta+Ti
0.03 17–18 12–12.5 2.3–2.7 0.06–0.08 1.6–2.0 0.5 0.025 0.005–0.01 NS* NS 0.3 0.25 10–20 0.15
0.02 16–18 10–14 2–3 0.06–0.12 2.0 1.0 0.015–0.04 0.03 NS NS NS 0.25 10 NS
0.03 17–18 11.5–12.5 2.3–2.7 0.06–0.08 1.6–2.0 0.5 0.035 0.025 0.05 0.05 1.0 0.25 15–35 NS
0.024–0.03 17–18 12–12.5 2.3–2.7 0.06–0.08 1.6–2.0 0.5 0.03 0.01 0.05 0.05 1.0 0.25 20 NS
NS, not specified
500
400
Stress (MPa)
300
200
316L(N) – Superphenix, France 316FR – DFBR, Japan 316L(N) – Germany 316L(N) – PFBR, India 316 – ORNL, USA 316L(N) – RCC-MR design curve 100 101
102
103 Rupture time (h)
104
105
22.10 Comparison of creep rupture strengths of 316 and 316L(N) stainless steels at 873 K from various countries.47
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101
Total strain range (%)
Japan type 316FR, 1 × 10–3 s–1, 50 mm plate CRNL type 316 FR, 1 × 10–3 s–1 or higher, 50 mm plate US type 316, 4 × 10–3 s–1, 538–566°C
100
10–1 101
550°C
102
103
104 Cycles to failure
105
106
107
22.11 Comparison of fatigue behaviour of 316 and 316FR stainless steels.54
22.4.3 Ferritic steels for FBR core application Though at present no commercial FBR is operating with a ferritic steel core, these steels have been extensively tested for this application in various reactors like FFTF in the USA,55 Phenix in France,56,57 PFR in the UK, and BN-350 and BN-600 in Russia.58,59 Extensive out-of-reactor irradiation studies have also been carried out on ferritic steels. The incentive for the use of ferritic steels is the high burn-up of fuel that can be achieved in reactors with such cores, which would in turn reduce the fuel-cycle cost considerably. With the austenitic stainless steels, which are the presently used clad and wrapper materials in FBR, the burn-up that can be targeted is ~100 000 MW days/ tonne, which corresponds to a neutron dose of 85 dpa and a life of ~ 2 years for the fuel sub-assembly. The life is limited by the radiation swelling that the austenitic stainless steel undergoes at these dose rates. With the use of ferritic steels, on the other hand, there is virtually no radiation swelling and hence the target burn-up can be at least doubled. Though swelling resistance of ferritic steels is superior to that of austenitic stainless steels, the rise in the DBTT under irradiation, poor formability and low high-temperature strength limit the choice of ferritic steels for FBR core application. There is also concern with respect to transfer of carbon from the ferritic steel side (high activity of carbon) to the austenitic steel (low activity of carbon) with flowing liquid sodium acting as transfer medium, which can further reduce the strength of the ferritic steel. However, the recent development of many advanced ferritic steels, some of them specifically for nuclear
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applications, and extensive in-reactor testing of these steels have provided enough confidence for the use of these materials in the core components of the FBRs. Table 22.3 gives the chemical composition of various ferritic steels considered for the core components of FBRs. Swelling behaviour and irradiation embrittlement of different ferritic steels have already been discussed in the previous sections. They indicate that swelling is lowest for steels with 12%Cr, while shift in DBTT is minimum for steels with 9%Cr. Figure 22.12 shows the influence of chromium content on DBTT shift in ferritic/martenistic steels by neutron irradiation.60 However, the swelling resistance of both 9%Cr and 12%Cr ferritic steels is far lower than that of austenitic stainless steels. Hence, for wrapper application, 9%Cr steels seem to be more appropriate than 12%Cr steels. Since the DBTT of 9%Cr steels is below ambient temperature even after irradiation, it will ensure that, during handling of the fuel sub-assembly removed from the reactor at ambient temperature, the risk of fuel pins coming out owing to failure of the wrapper will be minimum. Material selection for the clad tube is more challenging than for the wrapper. The operating temperature of the clad tube is high (600–700°C), with some transient that can go even higher, though the neutron dose experienced is comparable to that in the wrapper. Ferritic steels made through the conventional processing route of melting, casting and forming do not possess adequate creep resistance in this temperature range and hence are not ideal replacements for the presently used austenitic alloys. This has led to the development of a new class of ODS ferritic/martensitic steels for use as clad material. Y2O3 is the most widely used type of oxide particles added in these alloys and they react with Ti present in the steel to form complex oxides of yttrium and titanium. Ukai et al.61 studied the creep properties of ODS alloys containing Y2O3 (12Ce–2WTiY2O3) and have demonstrated that at 700°C, the creep properties of these alloys are superior to those of conventional ferritic steels like HT9 and PNC-FMS, austenitic stainless steel SUS316 and another ODS alloy MA957 (14Cr-0.9Ti–0.3MoY2O3) developed in the USA.62 However, fabrication of thin clad tubes from such steels, which are produced through the powder metallurgical route, is a major technological challenge that has to be overcome before these alloys are used in FBRs. At present, mixed oxides (a mixture of uranium and plutonium oxides) are used as fuels for FBRs and a ferritic wrapper with austenitic clad seems to be the appropriate choice for these fuels. However, there is a strong incentive to switch from oxide fuels to metallic fuels to achieve higher burnup levels. The alloys that are under active consideration are based on the U– Pu–Zr system. Use of metallic fuel necessitates shifting clad material from austenitic to ferritic steels. This is because in addition to reduced swelling in the ferritic steels, the damage caused by fuel-clad interaction (FCI) is less for ferritic steels than the austenitic steels. FCI can be either mechanical or
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Table 22.3 Chemical composition of major ferritic steels developed and studied in different countries for FBR core applications (wt%) Cr
Ni
Mo
Si
Mn
V
Nb
Ti
P
S
N
B
Others
UK FI FV607 CRM–12 FV448
0.15 0.13 0.19 0.10
13.0 11.1 11.8 10.7
0.47 0.59 0.42 0.64
– 0.93 0.96 0.64
0.30 0.53 0.45 0.38
0.45 0.80 0.54 0.86
– 0.27 0.30 0.16
– – – 0.30
– – – –
– – – –
– – – –
– – – –
– – – –
– – – –
France F17 EM10 EM12 T91
0.05 0.10 0.10 0.10
17.0 9.0 9.0 9.0
0.10 0.20 0.30 <0.40
– 1.0 2.0 0.95
0.30 0.30 0.40 0.35
0.40 0.50 1.00 0.45
– – 0.40 0.22
– – 0.50 0.08
– – – –
≤0.008 ≤0.008 ≤0.008 ≤0.008
≤0.008 – ≤0.008 ≤0.008
0.020 – – 0.050
– – – –
– – – –
Germany 1.4923 1.4914 1.4914 mod
0.21 0.14 0.16– 0.18
11.2 11.3 10.2– 10.7
0.42 0.70 0.75– 0.95
0.83 0.50 0.45– 0.65
0.37 0.45 0.25– 0.35
0.50 0.35 0.60– 0.80
0.21 0.30 0.20– 0.30
– 0.25 0.10– 0.25
– – –
– – –
– – 0.010
– 0.029 0.0015 max.
– 0.007 – max.
– –
USA HT9 403
0.20 0.12
11.9 12.0
0.62 0.15
0.91 –
0.38 0.35
0.59 0.48
0.30 –
– –
– –
– –
– –
– –
– –
0.52 (W) –
0.2
11
0.4
0.5
–
–
0.2
0.05
–
–
–
0.05
–
–
0.2
11.0– 13.5
0.3– 0.5
1.2– 1.8
–
–
0.1– 0.3
0.3– 0.6
–
–
–
–
0.004
–
Japan PNC–FMS Russia EP450
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Using creep-resistant steels in nuclear reactors
ALLOY
624
Creep-resistant steels
250
2.25 CrV JLF-4
DBTT shift (°C)
200
JLF-6 12Cr-6MN-1W
2.25 Cr1WV 10 dpa; 365°C 2Cr-1.5V
12Cr-6MN-1V 36dpa; 410°C
150
12Cr2WV
2.25Cr-2W JLF-3
100
2.25Cr-2WV
5Cr2WV
F82H 9Cr-2WV
50
0 0
7dpa; 365°C 7dpa; 365°C 10dpa; 365°C 38dpa; 410°C
Cr-1V JLF-1 9Cr-1W 9Cr-2WVTs
2
4
6 8 Chromium content (wt%)
10
12
22.12 Variation in ∆DBTT with Cr content in ferritic steels.60
chemical. Fuel clad mechanical interaction (FCMI) arises from the applied stress when element design restrains fuel swelling and results in plastic deformation of the clad. Fuel clad chemical interaction (FCCI) is a complex multi-component diffusion problem; U and Pu in the fuel and rare-earth fission products generated during irradiation have a propensity to interact metallurgically with the clad material at high temperature.63–67 FCCI is an important consideration in the choice of the clad material, as the composition of the clad material has a significant effect on the above phenomenon. Ni is the major element that contributes to FCCI; it forms low-melting eutectics with U and enhances diffusion of U and Pu into the clad, thus increasing the diffusion layer thickness. As Ni is an important alloying element in austenitic stainless steel, FCCI is understandably greater in this class of steel than in ferritic steels. Further, loss of Ni from the clad surface can also change the structure at the surface from austenitic to ferritic, introducing additional uncertainty with respect to its performance. As a result, when using metallic fuel in FBRs, it is essential that both clad and wrapper should be of ferritic steels. However, presently available commercial ferritic steels produced via the ingot route do not possess the required high-temperature strength. This again brings to the fore the importance of ODS ferritic alloys as a clad tube material for FBRs. An alternative option is to design the reactors with lower temperatures of operation than in the present design with oxide fuels so that ferritic steels chosen for the wrapper application can themselves be used.
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22.4.4 Fusion reactors68 In fusion reactors, energy is produced from the fusion of nuclei of light elements and the most suitable fusion reaction occurs between the nuclei of two heavy isotopes of hydrogen, viz., deuterium and tritium, to form a helium nucleus, accompanied by the release of a neutron and energy: 2
D1 + 3T1 → 4He2 (13.5 MeV) + 1n0 (14.1 MeV)
The temperature required for this reaction is of the order of millions of degrees Celsius and at such temperatures fuel changes from the gaseous state to plasma. The hot plasma is magnetically contained in a vacuum vessel and isolated from the walls of the vessel. The most promising magnetic confinement systems are toroidal (ring-shaped) and the most advanced of these is the Tokomak reactor. Inertially confined fusion systems in which energy is produced by repeated ignition of D-T pellets by focused laser beams are also being examined. The deuterium fuel is abundant and is easily extracted from water. Tritium can be produced by bombarding Li with neutrons. Lithium is kept as a blanket surrounding the vacuum vessel for breeding tritium and a breeding ratio of more than one can be achieved with this blanket. The overall fusion reaction can thus be represented as follows: 2
D1 + 6Li3 → 2 4He2 + 22.4 MeV
The inner wall of the fusion reactor vessel, which faces the plasma and is exposed to nuclear radiation and fusion products, is often referred to as the first wall. Selection of materials for this component is an important aspect in the development of fusion reactors. The first wall is subjected to the following severe conditions during service: • • •
mechanical and electromagnetic loading, and alternating thermal stresses induced by surface and volumetric heating owing to the pulsed nature of the operation; irradiation with high-energy (14.1 MeV) fusion neutrons producing displaced atoms and helium, hydrogen and transmutation products, leading to changes in bulk properties; and bombardment with ions and energetic neutral atoms from the plasma, resulting in surface and near-surface damage.
Potential structural materials considered for this application include austenitic steels based on the Fe–Cr–Ni and Fe–Cr–Mn systems, Cr–Mo ferritic/martensitic steels, alloys based on vanadium, niobium, molybdenum, titanium, and so on and SiC/SiC composites. The Cr–Mo and high-Cr ferritic steels initially considered for this application are the same as those envisaged for clad and wrapper applications of fast breeder reactors.
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Unlike in fission reactors, there are no highly radioactive products resulting from nuclear fusion. However, the high-energy neutrons can be absorbed by nuclei of various elements present in the structural material and these nuclei undergo transmutation producing radioactive isotopes of various elements. There is thus a specific interest in eliminating or minimizing those elements that can produce radioactive products as a result of neutron absorption and subsequent transmutation. A major limitation in evaluating structural materials for fusion reactor application is the inability to conduct in-reactor irradiation studies on the candidate materials. One has to rely heavily on the data on neutron crosssection, radioactive isotopes, emission from these isotopes, half-life, and so on. Accordingly, inventory codes and cross-section and decay libraries have been developed in various countries to predict the radionuclide inventories of materials exposed in fusion reactors.69,70 Using these codes and libraries, the type of radiation, activity levels after stipulated cooling time, decay heat, and so on are predicted for the reference neutron spectrum to which various reactor components like first wall, blanket, shield and magnetic coils are exposed. In the case of ferritic steels considered for fusion reactor applications, the effects of various alloying elements on contact γ-dose rate, induced activity and decay heat have been evaluated using the codes and libraries for a neutron loading of 2 MW m–2 for 2.5 years. These results show that Cr (any concentration), V (≤8%), Mn (≤1%), Ta (~1%) and Si (<0.4%) are acceptable, while Mo (>100 ppm), Nb (>1 ppb) and Ni (>50 ppm) are unacceptable alloying additions. Further, C, B and Ti in the concentration that is generally present in these steels do not detrimentally affect the activation parameters and W is not an ideal replacement for Mo in these steels.71 In addition to major and minor alloying elements, the effect of impurity elements in these steels on the γ-dose rate contribution has been evaluated using these libraries and codes and the results show that impurities like Pd, Ag, Bi, Hf, and so on should be maintained below 1 ppm.72 These results clearly show that compositions of the various ferritic steels presently considered for fast breeder reactor core application will not satisfy the stringent requirement on contact γ-dose rate (≤25 µSv h–1 after 100 y cooling), induced activity (<103Bq kg–1 for unrestricted release) and decay heat (<1 W m–3 for low level waste after 50 years of interim storage).73–75 This has resulted in the development of reduced activation ferritic steels specifically for fusion reactor applications. The principal approaches adopted in the development are the replacement of radiologically undesirable Mo, Nb and Ni in the commercial steels by elements such as W, V, Mn, Ta and Ti which have equivalent or similar effects on the constitution and structures and the removal of impurities that adversely affect the induced activities and dose rates, even when present in low concentrations in steels. Typical examples
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of reduced activation steels are EUROFER (Europe), F82H and JLF-1 (Japan) and 9Cr–2WVTa (USA). Compositions of these steels are given in Table 22.4.68
22.4.5 Ferritic steels for steam generators and turbines in nuclear plants As mentioned in the introduction, in most nuclear reactors, the steam produced using the heat generated from nuclear fission is used to drive the turbine. As in the case of fossil power plants, the materials for construction of turbines are the ferritic steels. For rotors, steels like Cr–Mo–V steel are used and for turbine blades and shrouds 12%Cr martensitic stainless steels are used. The steam generators for PWRs and PHWRs do not use creep-resistant steels as the steam temperatures are low (<300°C) in these reactors. However, in the case of FBRs, steam temperatures are comparable to those in fossil power plant and creep-resistant steels are employed. FBR steam generators have liquid sodium in the shell side and high temperature water and steam in the tube side. Since any leak in the tube would bring water or steam into direct contact with liquid sodium, resulting in a very aggressive reaction, the choice of the material and the design of the steam generator must be such that the possibility of tube rupture is eliminated. Hence, though austenitic stainless steel is chosen as the main structural material for FBRs, ferritic steels are considered for steam generators. The ferritic steels possess better resistance to stress corrosion cracking in the presence of high-temperature steam or water, good thermal conductivity (required for transfer of heat from liquid Na to water) and a low thermal expansion coefficient and have sufficient creep strength at the operating temperature of ~500°C.47 One matter of concern, however, is the possibility of carbon transfer from ferritic steel to austenitic steel through liquid sodium owing to the difference in carbon activities between the two steels. This can result in carburization of the latter making it brittle and decarburization of the former making it soft. In fact, it is interesting to note that modified 9Cr–1Mo steel was originally developed by Oak Ridge National Laboratory specifically for use in steam generators of FBRs.76 The attempt was to reduce the carbon activity in the steel by adding Nb and V and introducing N to form carbonitrides instead of just carbides as in conventional 2.25Cr–1Mo and ordinary 9Cr–1Mo steels. As the FBR programme in the USA came to a standstill, this alloy did not find immediate application in FBRs, but its superiority over the conventional Cr–Mo steels was recognized by power industries worldwide and this alloy emerged as the preferred material for fossil power plant replacing the 2.25Cr–1Mo and 12Cr–1MoVNb steels that were in wide use in the industry at that time. Modified 9Cr–1Mo steel led the new generation of advanced ferritic steels
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Programme
Name
C
Si
Mn
Cr
W
V
Ta
N
B
CEC
LA12TALC
0.09
0.03
1.0
8.9
0.8
0.40
0.10
0.02
–
EUROFER
0.10–
0.05
0.4–
8.0–
1.0–
0.20–
0.06–
0.02–
0.004–
0.12
(max)
0.6
9.0
1.2
0.30
0.10
0.04
0.006 0.003
Japan
USA
F82H
0.10
0.20
0.50
8.0
2.0
0.20
0.04
<0.01
JLF-1
0.10
0.08
0.45
9.0
2.0
0.20
0.07
0.05
9Cr–2WVTa
0.10
0.30
0.40
9.0
2.0
0.25
0.07
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Table 22.4 Typical/nominal compositions of reduced activation martensitic steels with a favourable combination of properties (wt%)68
Using creep-resistant steels in nuclear reactors
629
pursued actively by the fossil power industries to increase the efficiency of power plants and reduce CO2 emissions. At present, modified 9Cr–1Mo steel is the first choice as structural material for steam generators of FBRs; for India’s PFBR, steam generators made from it are under fabrication. The Demonstration Fast Breeder Reactor (DFBR) and European Fast Breeder Reactor (EFR) also chose this steel for steam generators, although the reactors themselves were not constructed. However, FBRs constructed or conceived in earlier years like BN-600 and BN-800 (Russia), SNR 300 and MONJU ( Japan) and Phenix (France) had either 2.25Cr–1Mo steel or a combination of 2.25Cr–1Mo steel (for evaporator) and austenitic stainless steel (for superheater) for their steam generator. For Superphenix of France, Alloy 800 was the material of construction for the steam generator.47 Though many advanced ferritic steels with better high-temperature properties than modified 9Cr–1Mo steel are currently available, it is unlikely that any of the new steels would be considered for FBR steam generators in the immediate future. This is because the steam temperature in FBRs is limited by the maximum temperature of liquid sodium, which is in turn depends on the maximum temperature of the reactor core. The reactor core temperature is limited by the properties of the structural material, and with both the currently used austenitic stainless and the ferritic steels of the immediate future not much increase in the reactor temperature is expected. Further, at ~ 500°C, type IV cracking, resulting from the poor creep strength of the weld joint, is not a serious problem in modified 9Cr–1Mo steel. However, once metallic fuel and ODS clad tubes are introduced in FBRs, it would be possible to increase the steam temperature and hence new alloys could be considered.
22.5
Fabrication and joining considerations
From the above discussion, it is clear that austenitic stainless steels and advanced ferritic steels are the main creep-resistant steels considered for nuclear applications. Among these, austenitic stainless steels are widely used in FBRs and advanced ferritic steels are the future structural material for core components of FBRs and fusion reactors. Components of FBRs vary widely in size and dimensions; for Indian Prototype Fast Breeder Reactor (PFBR) clad tubes have a wall thickness <0.5 mm with a diameter of ~ 6 mm and a length of a few metres, while the grid plate, which supports the fuel sub-assemblies, is a massive plate that is 60–80 mm in thickness and ~6 m in diameter. Fabrication of components from plates, pipes, tubes, and so on made from austenitic stainless steels have not been a serious technological challenge because of their good formability and weldability. However, there have been specific cases of reactor component fabrication that needed extensive development even with the use of austenitic stainless
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steels. One such case is the end cap welding of fuel clad tubes made from the 15Cr–15Ni–Ti class of steels. The hot cracking susceptibility of D9 alloy (15Cr–15Ni–Ti steel) chosen for clad and wrapper and 316L(N) steel chosen for all the major structural components of PFBR has been studied in detail using the Varestraint test. The results have demonstrated the much greater susceptibility of D9 alloy to solidification cracking than 316L(N) steel.77–79 The hot cracking problem in stainless steel welding is generally avoided by ensuring some minimum volume fraction of δ-ferrite in the weld metal, which in turn is achieved by modification in the composition of the welding consumables employed. However, end cap welding of the clad tube is done using autogenous gas tungsten-arc welding with no filler addition and the D9 alloy solidifies in the fully austenitic mode with no possibility of δferrite formation. Elements like Ti, P, S and N present in the steel also contribute to cracking by forming low-melting eutectics. It has been found that cracking in the end cap welding of the clad tube can be avoided only by the use of an end plug made of a different material like 316L(N) steel. Further, welding parameters need to be carefully controlled in a narrow range to ensure defect-free welds. Another weld joint in FBRs that has been subjected to extensive study is the dissimilar joint involving austenitic stainless steel piping and the steam generator header made from ferritic steel. Operating experience from many fossil power plants has shown that these joints fail prematurely in the ferritic steel side of the joint80 and the reasons attributed to the failure are the secondary stresses generated by the difference in thermal expansion coefficients of the two steels, carbon migration from the ferritic steel side to the austenitic steel owing to difference in the carbon activities, strain accumulation at the ferritic steel owing to the difference in creep strengths of the two steels, oxide notching in the ferritic steel side, and so on. A trimetallic transition metal joint involving the use of an intermediate Alloy 800 piece between the austenitic and ferritic steels has been suggested for improving the life of the joint. Austenitic stainless steel is welded to Alloy 800 using ER 16-8-2 filler wire (an austenitic stainless steel consumable), and Alloy 800 is welded to the ferritic steel using ER NiCr–2/ENiCrFe–3 (Ni-base) consumables.81,82 The properties of these joints have been studied extensively and the results have shown that the life of the trimetallic joints is many times longer than that of the direct joint between austenitic stainless and ferritic steels.81–86 In addition to Alloy 800, another Ni-base alloy Alloy 600 has also been used as a transition piece in the dissimilar joints between austenitic and ferritic steels developed for fusion reactor applications.87,88 A recent trend in the transition joints involving these two classes of steels is the use of inserts of graded composition manufactured by powder metallurgy using the hot isostatic pressing (PM-HIP) process.89 As the formability and weldability aspects of ferritic steels are considerably
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different from those of austenitic stainless steels, changing the material of construction for core components from the latter to the former in the near future would introduce additional issues of fabrication and joining. Drawing thin clad tubes and forming a hexagonal wrapper from ferritic steels are much more difficult than producing them from austenitic stainless steels. Similarly, welding ferritic steels calls for preheating and postweld heat treatment and this would necessitate localized heat treatment of the weld joint without having any adverse effect on the rest of the component. At present, ODS ferritic steels developed specifically for core components of FBRs and fusion reactors suffer from anisotropy in properties and difficulties in forming and joining. Processing them into thin clad tubes for FBR application has proved to be a difficult task. Hamilton et al.62 describe the technological efforts taken up to develop successfully the processing route for an ODS alloy MA 957 from rods to clad tubes. Tubes were made from bar stock using a combination of processes involving gun drilling, rod drawing, rerodding and final plug drawing. The product was also subjected to annealing in the different stages of the forming operation. It was required to control the annealing and forming temperatures and reduction in each stage of forming very carefully to complete the forming operation satisfactorily. Later, end cap welding of the tube was also successfully demonstrated. In spite of these developments and the significant R&D activities currently taking place, it is unlikely, however, that ODS alloys will be used in commercial reactors in the near future. High-Cr ferritic steels produced by the conventional processing route and their low-activation counterparts are more likely to be the materials that may find application in the nuclear reactors in the near future.90
22.6
Summary
Austenitic stainless steels and high-Cr ferritic steels are the major classes of heat-resistant steels considered for nuclear applications. Among these, the austenitic steels are already in use as major structural materials for FBRs. These steels are also used as a major piping material in BWRs and are being considered for fusion reactor application. However, the high susceptibility of austenitic steels to radiation swelling, helium embrittlement, irradiation creep, and so on have necessitated replacement of these alloys for the reactor core components with ferritic steels which are much more resistant to radiation swelling. Extensive in-reactor studies on advanced ferritic steels have shown that these alloys can tolerate neutron doses as high 200 dpa and energy production as high as 200 MW days/tonne can be achieved with the use of ferritic wrapper and clad material, compared to <100 MW d ton–1 currently achieved with austenitic clad and wrapper in FBRs. However, the creep strength of these steels is inferior to that of austenitic stainless steels and they also exhibit irradiation embrittlement (reflected as a large shift in DBTT
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Creep-resistant steels
under irradiation). Creep strength is important for ferritic steels considered for clad tubes, for which the operating temperature is high, and a relatively low shift in DBTT is important for ferritic steels considered for wrapper fabrication. ODS alloys with improved creep resistance have been developed for clad application using the powder metallurgy route and thin clad tubes have been produced successfully from them. For wrapper application, the choice is between 12%Cr steels like HT9, which exhibit good resistance to radiation swelling and 9%Cr steels like EM10 or 9Cr-1Mo steels which show minimum shift in DBTT; between them, the preference is for the latter. Ferritic steels considered for fusion reactor application (first wall of the fusion chamber) have been derived essentially from those studied for FBR core application. Emphasis is also being given to minimizing or avoiding those elements in the steel that produce radioactive isotopes with long life when exposed to high-energy neutrons present in fusion reactors. This has led to the development of reduced-activation steels specifically for use in fusion reactors. Elements like Ni, Mo, Nb, which produce radioactive isotopes when exposed to the reactor environment, are replaced by W, V and Ta. Further, trace elements that can produce radioactive isotopes are identified and upper limits are specified in these alloys. The purpose is to ensure safety of the personnel during maintenance, waste disposal and recycling of the components. There is renewed interest worldwide in nuclear energy and FBRs are going to play a major role in the revival of nuclear energy as an alternative to fossil fuel. Commercial reactors operating on nuclear fusion are also likely to be a reality in the future. It may therefore be expected that the use of creep-resistant steels, especially advanced ferritic steels, in nuclear industries is going to increase in a manner similar to their use in fossil power plants.
22.7
References
1 S. Glasstone and A. Sesonskei, Nuclear Reactor Engineering, Vol. 1, CBS Publishers and Distributors, New Delhi, 2001, 72. 2 V. McLane, C.L. Dunford and P.F. Rose, Neutron Cross Section, Vol. 2, Neutron Cross Section Curves, Academic Press, 1988, 184. 3 L.K. Mansur, in Kinetics of Non-Homogeneous Processes, G.R. Freeman (ed.), John Wiley & Son, New York, 377. 4 R.L. Klueh and D.R. Harries, High Chromium Ferritic and Martensitic Steels for Nuclear Applications, Chapter 8, ASTM, Philadelphia, 2001, 81. 5 E.A. Little, J. Nucl. Mater., 1993, 206, 324. 6 R.L. Klueh and D.R. Harries, High Chromium Ferritic and Martensitic Steels for Nuclear Applications, Chapter 9, ASTM, Philadelphia, 2001, 90. 7 F.A. Garner, J.F. Bates and M.A. Mitchell, J. Nucl. Mater., 1992, 189, 201. 8 J.L. Seran, V. Levy, P. Dubuisson, D. Gilbon, A. Maillard, A. Fissolo, H. Touron, R. Cauvin, A. Chalony and E. Le Boublbin, in Effect of Radiation in Materials: 15th
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10 11 12 13 14 15 16 17 18 19 20
21
22
23
24 25 26 27
28
29 30
31
32
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International Symposium ASTM STP 1125, R.E. Stoller, A.S. Kumar and D.S. Gelles (eds), ASTM, Philadelphia, 1992, 1209. M.I. Hamilton, C.D. Johnsion, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yang and N. Abraham in Residual and Unspecified Elements in Steel, ASTM STP 1042, A.S. Melilli and E.G. Nisbett (eds), ASTM, Philadelphia, 1969, 124. Y. Yateishi, J. Nucl. Sci. Tech., 1989, 26, 132. I. Shibahara, N. Akasaka, S. Onose, H. Okada and S. Ukai, J. Nucl Mater., 212–215 1994, 487. G. R. Odette; J. Nucl. Mater., 1988, 155–157, 921. I.S. Kim, J.D. Hunn, N. Hashimoto, D.L. Larson, P.J. Maziasz, K. Miyahara and E.H. Lee, J. Nucl. Mater., 2000, 280, 264. F.A. Garner, M.B. Toloczko and B.H. Sence, J. Nucl. Mater., 2000, 276, 123. L.K. Mansur, ‘Theory and experimental background on dimensional changes in irradiated alloys’, J. Nucl. Mater., 1994, 216, 97–123. R.L. Klueh and D.R. Harries, ‘High chromium ferritic and martensitic steels for nuclear applications Chapter 9, ASTM, Philadelphia, 2001 p. 113. J.O. Stiegler and L.K. Mansur, Ann. Rev. Mater. Sci. 1979, 9, 405. J.R. Mathews and M.W. Finnis, J. Nucl. Mater. 1988, 159, 257. E.R. Gilbert and B.A. Chin, Nucl. Tech., 1981, 52, 273. B.A. Chin, ‘An analysis of the creep properties of a 12Cr–1MoWV steel’, Topical Conference on Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J. Michel (eds), AIME, Warrendale, PA, USA, 1984, 593. C. Wassilew, K. Herschbach, E. Materna-Morris and K. Ehrlich, in Topical Conference on Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J. Michel (eds), AIME, Warrendale, PA, USA, 1984, 607. R.J. Puigh and G.L. Wire, in Topical Conference on Ferritic Steels for Use in Nuclear Technologies, Eds. J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA, 1984, 601. J.M. Dupuoy, Y. Carteret, H. Aubert and J.L. Boutard in Topical Conference on Ferritic Steels for Use in Nuclear Technologies, Eds. J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA, 1984., 125. M.B. Toloczko, F.A. Garner and C.R. Eiholzer, J. Nucl. Mater. 1994, 212–215, 604. R.L. Klueh and D.R. Harries, High chromium ferritic and martensitic steels for nuclear applications, Chapter 9, ASTM, Philadelphia, 2001, 139. R.L. Klueh and D.J. Alexander, J. Nucl. Mater, 1992, 187, 60. R.L. Klueh and D.J. Alexander, in Effect of Radiation in Materials: 15th International Symposium ASTM STP 1125, R.E. Stoller, A.S. Kumar and D.S. Gelles (eds), ASTM, Philadelphia, 1992, 1256. W.L. Hu and D.S. Gelles, in Topical Conference on Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J. Michel (eds), AIME, Warrendale, PA, USA, 1984, 631. Y. Dai, X.J. Jia and K. Farrell, J. Nucl. Mater., 2003, 318, 192. P.J. Maziasz and R.L. Klueh, in Effect of Radiation in Materials: 15th International Symposium ASTM STP 1125, R.E. Stoller, A.S. Kuma and D.S. Gelles (eds), ASTM, Philadelphia, 1992, 1135. D.S. Gelles and L.E. Thomas, Topical Conference on Ferritic Steels for Use in Nuclear Technologies, J.W. Davis and D.J. Michel (eds), AIME, Warrendale, PA, USA, 1984, 559. W.L. Hu and D.S. Gelles, in Influence of Radiation on Materials: 13th International
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39 40
41 42 43
44 45 46 47 48 49
50
51 52
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Creep-resistant steels Symposium Part II, ASTM STP 956, F.A. Garner, C. H. Heneger Jr. and N. Igata (eds), ASTM, Philadelphia, 1987, 83. R.L. Klueh and D.J. Alexander, J. Nucl. Mater., 1992, 187, 60. R.L. Klueh and D.J. Alexander, J. Nucl. Mater., 1992, 191–194, 896. B.A. Chin and R.C. Wilcox, Topical Conference on Ferritic Steels for Use in Nuclear Technologies, (eds), J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA, 1984, 347. H.-C. Schneider, B. Dafferner and J. Akata, J. Nucl. Mater, 2001, 295, 26. R.L. Klueh, M.A. Sokolv, K. Shiba, Y. Miwa, J.P. Robertson, J. Nucl. Mater., 2000, 283–287, 478. R.L. Klueh and D.R. Harries, High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, American Society for Testing and Materials, Philadelphia, 2001, 135. R.L. Klueh and J.M. Vitek, ‘The resistance of 9Cr–1MoVNb and 12Cr–1MoVW steels to helium embrittlement’, J. Nucl. Mater., 1983, 117, 295–302. U. Stamm and H. Schroeder, ‘The influence of helium on the high-temperature mechanical properties of DIN 1.4914 martensitic steel’, J. Nucl. Mater., 1988, 155– 157, 1059–1063. Hasegawa and H. Shiraishi, ‘Helium implantation effects on low activation 9Cr martensitic steels’, J. Nucl. Mater., 1992, 191–194, 910–914. K.J. Harrison, S.H. Rou and R.C. Wilcox, J. Nucl. Mater., 1986, 141–143, 508. D.J. Alexander, P.J. Maziasz and C.R. Brinkman, in Microstructures and Mechanical Properties of Ageing Materials, P.K. Liaw, R. Viswanathan, K.L. Murthy, E.P. Simonen and D. Frear (eds), The Minerals, Metals and Materials Society, Warendale, PA, 1993, 343. F.W. Noble, B.A. Senior and B.L. Eyre, Acta. Met., 35(5), 709. C.A. Hippsley and N.P. Howarth, Mater. Sci. Tech., 1988, 4, 791. S. Glasstone and A. Sesonske, Nuclear Reactor Engineering, Volume 2, CBS Publishers & Distributors, New Delhi, 2001, 820. S.L. Mannan, S.C. Chetal, Baldev Raj and S.B. Bhoje, in Materials R&D for PFBR, S.L. Mannan and M.D. Mathew (eds), IGCAR, Kalpakkam, 2003, 9. K. Herschbach, W. Schneider and K. Ehrlich, J. Nucl. Mater., 1993, 203, 233–245. M.I. Hamilton, C.D. Johnson, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yang and N. Abraham, ‘The effect of phosphorus and boron on the behaviour of titanium stabilised austenitic stainless steel developed for fast breeder reactor service’ Report No. PNL-SA\15303 1987. M. Fujiwara, H. Uchida, S. Ohta, S. Yuhaara, S. Tani and Y. Sato, in Radiation Induced Changes in Microstructure 13th International Symposium Part 1, ASTM STP 955, F.A. Garner, N.H. Packan and A. S. Kumar (eds), ASTM, Pa, 1987, 127. M. Fujiwara, H. Uchida and S. Ohta, J. Mater. Sci., 1994, 13, 908. M.L. Hamilton, G.D. Johnson, R.J. Puigh, F.A. Garner, P.J. Maziasz, W.J.S. Yang and N. Abraham, ‘The effect of phosphorus and boron on the behaviour of titaniumstabilized austenitic stainless steel developed for fast breeder reactor service’ Report No. PNL-SA-15303 1987. Schirra M, Creep and Fracture of Engineering Materials, Parker, J.D. (ed.), The Institution of Materials, London, 612. Brinkman C.R. ‘Elevated temperature mechanical properties of an advanced type 316 stainless steel’, Report No. ORNL/CP-101053, National Technology Information Science (NTIS), USA 1999.
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55 A.J. Lovell, D.R. Wilson, D.F. Leibnitz and W.H. Sutherland, Topical Conference on Ferritic Steels for Use in Nuclear Technologies, eds. J.W. Davis and D.J. Michel, AIME, Warrendale, PA, USA, 1984, 135. 56 C. Brown, V. Levy, J.L. Serin, K. Ehrlich, R.J.C. Rojer and H. Bergmann, Proceedings Fast Reactors and Related Fuel Cycles, FR-91, Vol. 1, Atomic Energy Society, Tokyo, 1991, 7.5. 57 C. Brown, R.J. Lilley and G.C. Crittenden, Nucl. Eng., 1994, 35, 122. 58 V.S. Khabarov, A.M. Dvorisahin and S.I Porollo, Proceedings of Technical Committee Meeting on Influence of High Dose Irradiation on Advanced Reactor Core Structural and Fuel Materials, IAEA-TECH DOC 1039, IAEA, Vienna, 1998, 139. 59 V.M. Poplavsky and L.M. Zabudko, Proceedings of Technical Committee Meeting on Influence of High Dose Irradiation on Advanced Reactor Core Structural and Fuel Materials, IAEA-TECH DOC 1039, IAEA, Vienna, 1998, 7. 60 A. Kohyama, A. Hishinuma, D.S. Gelles, R.L. Klueh, W. Diets and K. Ehrlich, J. Nucl. Mater., 1996, 231–237, 138. 61 S. Ukai, T. Okuda, M. Fujiwara, T. Kohoyashi, S. Mizuta and H. Nakashima: J. Nucl. Sci. Tech., 2002, 39, 872. 62 M.L. Hamilton, D.S. Gelles, R.J. Lobsinger, M.M. Paxton and W.F. Brown Fabrication Technology for ODS Alloy MA957, Report No PNL-13165 Pacific Northwest Laboratory Washington, 2000. 63 H. Tsai, A.B. Cohen, M.C. Billone and L.A. Niemark, ‘Irradiation performance of U–Pu–Zr metal fuels for liquid metal-cooled reactors’, Report ANL/ET/CP-82776, ANL, Illinois, NTIS No DE 95013679, 1994. 64 R.G. Pahl, D.L. Porter, C.E. Lahm and G.L. Hoffmam. Metall. Mater. Trans. A, 1990, 21A, 1863. 65 D.D. Keiser Jr., ‘Review of compatibility of FR fuel and austenitic stainless steel’, Report ANLIFE/PP-83153, Illinois, NIIS No DE 97000 549, 1994. 66 M. Kurata, T. Inoue and C. Sari, J. Nucl. Mater., 1994, 208, 144. 67 Advanced Metallic Fuels for Fast Reactors in Development – Status of Metallic, Dispersion and non-oxide advanced and alternative fuels for power and research reactors, IAEA TECDOC 1374, Ch. 3, Sept. 2003. 68 R.L. Klueh and D.R. Harries, High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, American Society for Testing and Materials, Philadelphia, 2001, 5. 69 D.R. Harries, G.J. Butterworth, A. Hishnuma and F.W. Wiffers, J. Nucl. Mater., 1992, 191–194, 92. 70 R.A. Forrest and J. Kopecky, The European Activation System (EASY), IEA Advisory group meeting on FENDL 2, Vienna, November 1991. 71 C.A.B. Forty, G.J. Butterworth and J-Ch Sublet, J. Nucl. Mater., 1994, 212–215, 640. 72 R.A. Forest, M.G. Sowerby and D.A.J. Endacott, in Fusion Technology 1990, B.E. Keen, M. Huguet and R. Hemsworth (eds.), Volume 1, North Holland, Amsterdam, 1991, 797. 73 P. Rocco and M. Zucchetti, J. Fusion. Energy, 1993, 12, 201. 74 P. Rocco and M. Zucchetti, Fusion Eng. Design, 1992, 15, 235. 75 P. Rocco and M. Zucchetti, J. Nucl. Mater., 1994, 212–215, 649. 76 V.K. Sikka, C.T. Ward and K.C. Thomas in ‘Ferritic steels for high temperature applications’, International Conference at Warren, PA, 1981, 65. 77 V. Shankar, T.P.S. Gill, S.L. Mannan and S. Sundaresan, Sci. Tech. Welding Joining, 2000, 5, 91.
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78 V. Shankar, T.P.S. Gill, A.L.E. Terrance, S.L. Mannan and S. Sundaresan, Met. Mater. Trans. A, 2000, 31 A, 3109. 79 V. Shankar, T.P.S. Gill, S.L. Mannan and S. Sundaresan, Mater. Sci. Eng., 2003, 343, 170. 80 Lundin C.D., Weld. J., 1982, 61, 305. 81 A.K. Bhaduri, I. Gowrishankar, V. Seetharaman, S. Venkadesan and P. Rodriguez, Mater. Sci. Tech., 1988, 4 (11), 1020. 82 A.K. Bhaduri, S. Venkadesan, P. Rodriguez and P.G. Mukunda, Int. J. Pressure Vessels Piping, 1994, 58 (3), 251. 83 Sireesha, M., Albert, S.K., Shankar, V. and S. Sundaresan, J. Nucl. Mater., 2000, 279 (1), 65–76. 84 Mopati Sireesha, Shaju K. Albert and Subramania Sundaresan, Steel Res., 2002, 73 (1), 26–30. 85 M. Sireesha, V. Shankar, S.K. Albert and S. Sundaresan, Mater. Sci. Eng., 292 (1), 74–82. 86 M. Sireesha, S.K. Albert, S. Sundaresan Met. Mater. Trans., 2005, 36A, 1495–1506. 87 J. N. Soo, in Rupture Ductility of Creep Resistant Steel A. Strang (ed.), Book No. 522, The Institute of Materials, London, 1991, 282. 88 J.N. Soo, in Rupture Ductility of Creep Resistant Steel, A. Strang (ed.), Book No. 522 The Institute of Materials, London, 1991, 294. 89 J. Petersheim, in 1st Bodycote International HIP Conference, Vasteras. Sweden, 1995. 90 R.L. Klueh and D.R. Harries, High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, American Society for Testing and Materials, Philadelphia, 2001, 208.
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23 Creep damage – industry needs and future research and development R . V I S WA N A T H A N and R . T I L L E Y Electric Power Research Institute, USA
23.1
Introduction
The phenomena of creep and creep fatigue have assumed great industrial significance worldwide in recent years. There are a variety of reasons for this. First and foremost, industrial plants operating at high temperatures have been subject to aging and the fleet of these aging plants which have been operating well past their design life of 30 to 40 years is growing. Reducing the cost of production is paramount for staying competitive. Reducing capital costs by deferring replacement of expensive components, and reducing operating and maintenance costs by optimizing operation and maintenance procedures will both be the key strategic objectives of plant owners. In addition, operation at higher than design temperatures and cycling caused by the de-regulated energy market have exacerbated the failure process by interaction of creep and fatigue. A lot of research has therefore focused on quantifying damage occurring owing to creep and creep fatigue mechanisms. This chapter will address industry needs, present selected results from studies focusing on these needs and identify future needs in research and development (R&D). Creep damage in components can take many forms. Excessive deformation caused by creep can lead to failure by wall thinning, loss of clearance and other dimensional changes. A second form of damage can consist of the formation of cavities at the grain boundaries, which can then link up to form large cracks at high temperatures. In another failure scenario creep damage may be initiated at high temperatures, but final fracture may occur in a brittle mode at lower temperatures owing to increased thermal stresses arising from transient conditions during shut-down or start-up. While dimensional changes can be readily seen and measured, in the other two failure modes, however, the failure is insidious and damage initiation can be observed only using specialized techniques. While most studies have focused on initiation of damage, more recently the crack growth regime has also received attention. Literature pertaining to crack initiation by progressive damage and subsequent 637 WPNL2204
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growth of cracks is too vast to be described here. Any review of this literature has therefore to be selective. The emphasis here is on industry perspective on creep damage assessment and future R&D needs. Since much of the industry interest centers on characterizing and quantifying creep damage and relating the results to life consumption, this paper will deal selectively with issues relating to three categories of damage measurement techniques, that is, calculational, non-destructive and destructive analysis.
23.2
Calculational methods for estimating damage
The very first step taken in assessing creep damage is using calculational methods. Although it is relatively easy to quantify damage in laboratory creep tests conducted at constant temperature and stress (load), components in service hardly ever operate under constant conditions. Start–stop cycles, reduced power operation, thermal gradients and other factors result in variations in stresses and temperatures. Procedures are needed that will permit estimation of the cumulative damage under changing exposure conditions.
23.2.1 Damage rules for creep The most common approach to calculation of cumulative creep damage is to compute the amount of life expended by using time or strain fractions as measures of damage. When the fractional damages add up to unity, then failure is postulated to occur. The most prominent rules are as follows: Life-fraction rule (LFR)1: t
Σ ti = 1
[23.1]
ri
Strain-fraction rule2: ε
Σ εi = 1 ri
[23.2]
where ti and ε˙ i are the time spent and strain accrued at condition i, and tri and ε˙ ri are the rupture life and rupture strain under the same conditions. Several authors have shown that the LFR is valid for temperature changes but not for stress changes. It has also been shown that the ductility and failure mode are also important. In an industrial context the LFR is often used to calculate the cumulative damage leading up to crack initiation. To evaluate total life of the component, crack growth analysis is needed. To calculate cumulative damage, the plant records and the time–temperature history of the component are reviewed. The creep or creep–fatigue life fraction consumed is calculated using the history, material properties and damage
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rules. This procedure is usually inaccurate owing to errors in assumed history, in the material properties and in the damage rules. Temperature history information may be refined by supplemental non-destructive or destructive examinations such as microstructural studies, hardness measurements and oxide-scale measurements. The uncertainties in material properties can be reduced by building a database, particularly with respect to service-retrieved components, taking into account environmental effects and heat-to-heat variations
23.2.2 Linear damage rule for creep–fatigue A variety of rules have been enunciated for calculating total damage when both creep and fatigue damage are present in a component. The most popular among these is the linear damage rule, in which the life fractions consumed in creep (t/tr) and fatigue (N/Nf) are added as follows:
Σ NN + Σ tt = D f
[23.3]
r
Differing approaches to calculating N/Nf, t/tr and the value of D at failure have led to alternate rules which have been discussed by Viswanathan.3 A critique of damage rules as they apply to fossil plant components has been published by Viswanathan.4,5 A detailed review of literature shows that there are divergent opinions regarding which damage approach provides the best basis for life prediction. It is quite clear that a number of variables, such as test temperature, strain range, frequency, time and type of hold, waveform, ductility of the material and damage characteristics, affect the creep–fatigue life. The conclusions drawn in any investigation may therefore apply only to the envelope of material and test conditions used in that study. The validity of any damage approach has to be examined with reference to the material and service conditions relevant to a specific application. Broad generalizations based on laboratory tests, which often may have no relevance to actual component conditions, do not appear to be productive. Thus, one should use a tailored, case-specific approach for any given situation.
23.2.3 Ductility exhaustion Another approach that is widely used is the ductility–exhaustion approach. The ductility–exhaustion approach is simply a strain-based life-fraction rule in which the fatigue damage and creep damage are summed up in terms of the fractional strain damage for each category, as follows:6
1 = ∆ε p + ∆ε c Nf Dp Dc
[23.4]
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where ∆ερ is the plastic strain-range component at half life, Dρ is the fatigue ductility obtained from pure fatigue tests, ∆εc is the true tensile creep-strain component and Dc is the lower-boundary creep-rupture ductility of the material. The first term in Equation [23.4] denotes the fatigue-damage component and the second term denotes the creep-damage component. Although the first term is fairly easy to understand and obtain from test data, the second term, especially the definitions of ∆εc and Dc, needs to be clarified. The problem arises from two issues. First, not all creep strain is viewed as damaging and only that strain which accumulates below a critical strain rate necessary to cause constrained cavity growth is viewed as damaging.7 Second, the rupture ductility of a material is not a constant but decreases with decreasing strain rate. Hence, in defining a failure criterion, an appropriate lower boundary value has to be defined for Dc.
23.2.4 Lacunae in calculational methods One of the major problems in evaluating the applicability of calculational methods is that in many cases it is necessary to use all the available data in deriving the damage rule and thus it is possible to examine only the accuracy with which a given method describes the data. There also is a scarcity of instances in which service experience has been compared with results from specific life-prediction methods. In general, the available methods are utilized only to predict the lives of samples tested under laboratory conditions. Validation against test data in the laboratory and in-service data on actual equipment would lead to more confidence in the use of the various rules. Results from most studies show that even the best of the available methods can predict life only to within a factor of 2 to 3. Some of the cited reasons for these inaccuracies have already been discussed. Some additional reasons are: failure of the methods of modeling changing stress relaxation and creep characteristics caused by strain softening or hardening, use of monotonic creep data instead of cycle creep data, and lack of sufficiently extended duration test data. None of the damage rules available today is entirely based on sound mechanistic principles. They are all phenomenological in nature, involving empirical constants that are material dependent and difficult to evaluate. Extrapolation of the rules to materials and conditions outside the envelope covered by the specific investigation often results in unsuccessful life predictions. Material behavior under isothermal LCF conditions in the laboratory often turns out to be totally irrelevant to material behavior under thermomechanical fatigue cycles involving in-phase or out-of-phase thermal cycles in the field. For application to service components, the stress–strain variation for each type of transient and its time dependence must be known with accuracy. Such calculations are difficult and expensive to perform. Because of these limitations and the simplicity of the linear damage summation
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using the life-fraction rule, the latter approach continues to enjoy popularity in engineering applications. Despite the inaccuracy of calculation procedures to estimate cumulative damage, they will continue in use because of their simplicity, inexpensiveness and non-invasive character. They lend themselves to on-line monitoring since damage can be calculated in real time. Improved accuracy can be achieved by keeping better plant records, better definition of material properties by storing and aging samples from the original heat of material and treatment of the various uncertainties by probabilistic techniques. Furthermore, the calculation procedure will remain a mainstay of on-line monitoring techniques.
23.2.5 Non-relevance of isothermal low cycle data to thermomechanical fatigue It is extremely important that in calculating creep–fatigue damage under cyclic conditions, the appropriate data obtained using a strain cycle simulating the actual service strain cycle is used. In many instances of fatigue, the temperature varies along with the strain, giving rise to what is known as thermomechanical fatigue (TMF). Two simple waveforms in TMF testing are shown and compared with the LCF test in Fig. 23.1. If maximum temperature corresponds to peak compression, as in the center diagram in the figure, it is known as the out-of-phase cycle (OP). If the maximum tensile stress occurs at the peak temperature, it is known as the in-phase (IP) cycle. Depending further on when the hold time is superimposed, various cycle shapes are possible. In the past, thermal fatigue traditionally has been treated as being synonymous with isothermal LCF at the maximum temperature of the thermal cycle. Consequently, life prediction techniques have evolved from iso-thermal LCF literature. More recently, advances in finite element analysis and in servohydraulic test systems have made it possible to analyze complex thermal cycles and to conduct TMF tests under controlled conditions that simulate these complex cycles. The assumed equivalence of isothermal LCF tests and TMF tests has been brought into question as a result of number of studies and the pertinent issues have been reviewed by Viswanathan and Bernstein.8 High tensile strains at high temperature (IP) would favor creep, whereas high tensile strains at low temperature (OP) would favor cracking of the oxide scale at the low temperature followed by environmentally induced accelerated creep damage during subsequent high-temperature exposure. Hence, Kuwabara et al. rationalized that in case of materials where damage is driven by creep, IP cycles would be more damaging than OP cycles, for a given strain range.9 In other materials, where the environmental contribution is significant, OP cycles may be more damaging than IP or isothermal LCF cycles. In addition to environmental effects, differences also arise between
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Out-of-phase
In-phase
0
– + t 0 –
0
σ
∆ε
t
– + t
0
0 –
Temp. T
Tmin
∆ε
TmaxTmTmin
Time t
σ
+ 0
Tm
0
Time t
σ
Strain. ε
t
∆ε
Stress. σ
Stress. σ
Strain. ε
Time t ∆ε
TmaxTmTmin Tmin
0
+ 0
Tm
Strain. ε
T = Tmax
Tmax
ε
Stress. σ
Temp. T
Temp. T
Tmax
σ
+ t 0
∆ε
– + t
∆ε
ε
0 –
23.1 Schematic diagrams showing waveforms of temperature, strain, and stress in thermal and isothermal fatigue tests.
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Thermal fatigue Isothermal fatigue
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cycles in terms of the relaxed mean stresses. The relative severity of the different cycles can also change with material ductility, maximum temperature and hold time. Consequently, a simple classification of material behavior is not possible. A case in point is the cracking encountered in the ligaments between tube penetrations in CrMo steel piping in power plants illustrated in Fig. 23.2(a). The figure shows a deep crack, filled with oxides, indicating a very strong environmental contribution. The cracking mode has been identified as creep fatigue. The creep–fatigue damage summation approach was found to be inconsistent with the early initiation of cracks observed in the pipes. The actual failure involved crack initiation by repeated cracking of oxide scales at low temperatures during shutdown transients and subsequent creep damage at the high operating temperature. The metallography of the cracked region showed numerous oxide spikes (see Fig. 23.2) confirming that the oxide cracking is a crack initiation mechanism. This example clearly illustrates the need to use appropriate thermomechanical fatigue data simulating of actual component cycles in predicting the crack initiation life of components.9
23.3
Non-destructive evaluation methods
Conventional non-destructive evaluation (NDE) methods such as ultrasonic testing (UT), dye penetrant examination and eddy current examination reveal large flaws and crack-like defects. More advanced techniques are needed to detect fine-scale creep damage prior to formation of cracks and to be able to relate it to the remaining life. These techniques have specific limitations and have been reviewed in detail elsewhere. High-resolution NDE techniques such as positron annihilation, ultrasonic velocity measurements, phased and focused beam ultrasonic techniques and mechanistic property change measurements have been demonstrated to be capable of detecting creep cavities in steels and are continuing to evolve as field deployable, inspection options. In particular, linear phased array UT has gained significant application experience in creep damage detection and characterization in steam piping weld areas. The availability of portable equipment supporting 16-channel transducers is a key enabling development for this application. The key drivers for these advancements include the following: the ability to sweep the ultrasonic beam through multiple angles in fine increments; the ability to focus ultrasonic energy into a very localized region and the ability to capture inspection data and perform more elaborate signal processing using computer-based techniques to display and characterize the signal features better and to correlate such features to the damage present. Results are, however affected by microstrucual and grain size variations owing to variations in manufacturing and specific correlations are therefore needed.
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(a)
(b)
23.2 Ligament cracking in boiler headers, (b) oxide spike in a ligament crack.
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23.3.1 Phased array ultrasonic testing The use of phased array technology to focus the ultrasonic beam has been successfully applied to detecting pre-crack damage.10 In this approach, a specially designed probe with an annular array of elements produces a focused beam in which the depth of focus can be varied by changing the time delays for pulsing and receiving. The enhanced signal-to-noise performance of this technique provides the basis for refined detection and characterization. The inspection process with focusing is, however, inherently slow and generally requires a very localized application. Accordingly, other techniques, such as conventional ultrasonics or analytical modeling, are used to establish the areas for detailed inspection. The advanced transducer and conventional UT technologies benefit strongly by the opportunity for refined computer processing to the acquired ultrasonic response. Techniques involving waveform analysis, spectrum processing, noise analysis and others offer opportunities to improve the correlation between inspection results and the actual level of damage present in the component. In general, these areas are evolving through an extensive process of data acquisition for plant components followed by component removal and destructive verification of the actual level of damage. To date, these approaches have been offered as a means of improving defect characterization and avoiding false calls. Continued improvement can be expected in this area which will greatly benefit all forms of non-destructive testing.
23.3.2 Magnetic testing The magnetic property response by materials has long been identified as affected by early stage microstructural change (such as produced by creep evolution). Accordingly, considerable research has focused on the use of this effect as a technique to detect creep damage. In general, the significant challenge for these types of NDE tools for creep damage has been the need to separate microstructural changes caused by to creep damage from those caused by material processing. As noted previously, the creep damage process produces changes in material microstructure as cavitation evolves along grain boundaries and develops into cracking. Non-destructive testing based upon measuring changes in magnetic properties such as hysteresis, remanence and coercivity by applying a magnetic field to the component of interest would be sensitive to the development of creep damage. Specifically, it is postulated that the presence of voids along grain boundaries has an impact on domain wall activity and demagnetization response. Research has validated these qualitative relationships as shown in Fig. 23.3.12
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60 50 40
Se am Se am
w el d
30
w el
w
d
Reduction in EMF from base metal plate (%)
646
w
ith
µ
ith
µ
HA Z,
HA Z
,µ
W el d
20 10
1
µ
W el
bu
d
w
tn
ith
oc
c re
ep
ree p
2 Current (A)
da m
da m 3
ag e
ag e
4
23.3 Magnetic response variation owing to the presence of creep damage. µHAZ and µweld are the magnetic permeability of the heat affected zone of the weld, respectively.
23.3.3 Magnetic acoustic emission In addition to the measurement of the magnetic property response, it is possible to use an applied magnetic field to induce acoustic emissions in the material. This process, called magnetoacoustic emission (MAE), occurs generally when domain walls, in a changing magnetic field, move in discrete jumps from one pinned position to another, producing acoustic waves, which can then be detected by a transducer. Research, in modeling and experiment, indicates a proportionate reduction in MAE owing to microstructural creep damage.11 It is noted that this process relies on the high sensitivity of acoustic emission testing but is an active rather than passive approach to generating the emissions. Further, the domain wall excitation produced by the applied magnetic field is not identical to the stress variations during plant operations which induce defect areas to emit acoustically, as detected under the acoustic emission monitoring approach discussed previously. Research to refine the basic approach for magnetic based NDE is on-going. A recent effort has addressed the issue of creep damage development in austenitic materials. The interest here is that the microstructural changes associated with aging produce an increased magnetic response over time. It has been found that long-term creep of certain steels results in a decrease in the electrical resistivity. Hardness decreases have been correlated with expended life for softening-type ductile creep damage in rotor-type steels in Japan, although lack of knowledge of initial hardness makes direct correlation inaccurate. Density decreases have been correlated with the degree of cavitation
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damage. Use of microstructural catalogs, interparticle spacing measurements and carbide/ferrite chemistry changes have been investigated in steels and superalloys but the results have not been consistent in providing quantitative answers regarding life consumed. Microstructural features are used primarily as supplementary tools to estimate the local temperature. Strain measurements using capacitance-type strain gauges have been shown to be capable of monitoring long-term, localized strain accumulation and are likely to find widespread use.
23.3.4 Strain monitoring Strain measurements are often employed to detect and measure creep damage. Gross changes such as component swelling and other dimensional changes have been monitored in the past. Owing to unknown variations in the original dimensions, changes in dimensions cannot be determined with confidence. Dimensional measurements fail to provide indications of highly damaging and localized creep strains such as those in the heat-affected zones of welds and regions of stress concentrations in the base metal. Cracking can frequently occur without manifest overall strain. Furthermore, the critical strain accumulation preceding fracture can vary widely with a variety of operational and material parameters, and with stress state. To enable measurement of localized strains, an ‘off-line’ condition surveillance system has been developed. The system uses the replication principle to evaluate localized strains and life consumption.13 A surface grid is scribed at the region of interest and preserved by means of an oxidation-resistant coating. A hard replica of high stability is used to duplicate the grid. Biaxial strain assessment is then made by high-resolution measurement of the replicas taken at successive plant inspection shutdowns. A predictive strain-rate lifetime model approach is used to establish ‘fitness for service’. No field experience of this technique has been reported. More recently capacitance strain gauges have come into wide use. They have been found to be capable of monitoring long term, localized strain accumulation and are likely to find widespread use.
23.3.5 Acoustic emissions monitoring Defects in a material structure are known to produce detectable acoustic emissions under conditions of applied stress. In this fashion, acoustic emission (AE) has potential use as an on-line monitoring technology for creep and other damage. The localized sources of acoustic emission can originate from a number of causes. The resultant transient sound waves propagate radially throughout the structure, attenuating with distance from the source as a function of frequency. The wave direction and wave mode types can change as they are reflected and refracted at the boundaries of the structure. The
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waves are detected by a piezoelectric sensor, which transforms their mechanical displacement into an electrical signal. These signals are very small at the sensor (microvolts to millivolts) and must be filtered and amplified near the sensor before transmission via cable to electronic data acquisition and processing equipment. In AE inspection, the sound emitted by the material during changes in stress states is detected and correlated to its source. Recent advancements in both equipment and signal interpretation are providing a basis for moving the detection of damage from the cracking stage to the stage of pre-crack cavitation. In this case, the measurement is indirect since the acoustic emissions are from groups of cavities and are detected without absolute location within the thickness of the material. In recent years, specific application of AE monitoring has been made to creep damage detection in high energy steam piping.14 The key benefit of an acoustic emission monitoring approach is the economic leverage gained by inspection during actual plant operation and without the full removal of the piping insulation that would be required for conventional inspection techniques. Additionally, piping systems typically undergo variations in temperature and pressure that provide a variable stress environment for activating damage when emitting acoustically. Key improvements have been made in order to differentiate signals of no interest from those of critical interest via a range of pre- and post-data acquisition filtering of signal characteristics. On-going work is now examining the potential for quantitative (extent of damage) determinations from the AE processes.
23.3.6 Evolution of creep cavitation One of the most widely used non-destructive techniques is surface replication. In this technique, the damage surface features are completely replicated on an acetate tape that can then be examined under a microscope at high magnification to reveal the extent of creep cavitation. Portable microscopes have also been used for field examination of the components. Alternatively, extremely small slices of the component have also been removed and used for metallographic observation. These approaches were popularized when Neubauer and Wedel15 characterized cavity evolution in steels at four stages – isolated cavities, oriented cavities, linked cavities (microcracks) and macrocrack and formulated recommendations corresponding to the four stages of cavitation, as shown in Fig. 23.4. To provide a theoretical and quantitative basis for cavity evolution, Cane and Shammas16 used a constrained cavity growth model and proposed a relationship between the number fraction of cavitated boundaries (A parameter) and the life fraction consumed, t/tr, using heat-specific constants. Values of these constants either had to be assumed or determined experimentally for
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Fracture
III
Creep strain
D C B A II
Exposure time
23.4 Creep life assessment based on cavity classification.
each heat, thus diminishing the usefulness of the model. Based on interrupted creep tests on simulated heat-affected zone in 1Cr–0.5 steels, it was subsequently concluded that the data contained too much scatter to verify the life prediction model proposed by Cane and Shammas. The data could nevertheless be used empirically, by plotting all the data together in the form of a scatterband whose lower limits are defined by the equation: A = 0.517(t /tr) – 0.816
[23.5]
Direct correlations between cavity classification and expended life fraction have also been established for 1.25Cr–0.5Mo steels, as shown in Fig. 23.5. The cavity classification system has been further refined and enhanced to include additional subcategories. While these models apply to cavitation in the coarse-grained heat-affected zone (CGHAZ) of welds, cavitation in the fine-grained heat-affected zone (FGHAZ) known as type IV cracking is yet to be quantified. The literature pertaining to life prediction of type IV damaged welds has been reviewed Ellis and Viswanathan.17 Replication is also not useful when creep cavitation occurs below the surface. Catastrophic failures of longitudinal seam welded piping have been reported in the USA, where cracks started at the midwall region owing to stress concentration effects at weld cusps and resulted in catastrophic failure of a hot reheat pipe at the Mojave Boiler Station as shown in Fig. 23.6.18 The field replication data forming the basis of the Wedel–Neubauer recommendations have shown no consistent trends in cavitation evolution with operating time or with calculated creep life fraction consumed. More field data are needed before clear correlation can be established between replica results and cumulative creep damage. In using the A parameter method, the specific procedure used to measure A is crucial. The A parameter is defined as the number fraction of cavitating grain boundaries encountered in a line parallel to the direction of maximum
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Damage classification
D C B 1 A B C D
A 1
Ratings Undamaged Isolated cavities Oriented cavities Linked cavities Macrocracks
0 0
0.2
0.4 0.6 Expended life fraction, (t / tr)
0.8
1.0
23.5 Correlation between damage classification and expended creep life fraction for 1.25Cr–0.5Mo steels. (1) Undamaged, (A) isolated damage, (B) oriented cavities, (C) linked cavities, (D) macrocracks.
23.6 Rupture in Monroe No. 1 north hot reheat line.
principal stress. To measure A reproducibly, the procedure needs to be standardized. The density of cavitation has also been used as a measure of creep damage. There is a need for standardized measurement of the ‘A parameter as well as the cavity/density since there is no way of comparing/correlating
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results from different studies. There are a number of other limitations of this metallographic technique. Quantitative correlations vary with steel composition, steel ductility and other properties. Consequently, a large database is needed to define the degree of scatter in the quantitative relations. More field validation of the replication technique is also required.
23.3.7 Analysis of carbides Several investigators have attempted to correlate microstructural change with the extent of creep damage. Carbide evolution in 2.25Cr–1Mo steel has been investigated by Stevens and co-workers19,20 and Munson.21 In both studies, the amount of M6C carbides as a percentage of the total weight percent of carbides increased with time and temperature. Plotted in terms of a Larson– Miller time–temperature parameter, the results (Fig. 23.7) show the family of curves from the Stevens and Flewitt study to be shifted laterally by a significant amount compared to that of Munson. These studies suggest that although the evolution of M6C may be used as a qualitative index of service temperature, wide variations may occur owing to differences in initial microstructure and composition. Data in Fig. 23.7 also show that the kinetics of M6C formation are considerably accelerated by phosphrous. Improved procedures for carbide extraction, as well as a larger database on samples with various initial compositions and heat treatments are needed before the microstructure technique can be used assessment of plant life. 0.04 %P 0.02 %P 0.005 %P Aged under stress Munson
60
50
Stevens and Flewitt
M6C (%)
40
30
20
10
0 34
35
36 37 38 39 LMP = (T + 460) (20 + logt)
40
41
23.7 Evolution of M6C in 2.25 Cr–1Mo steel as a function of aging time and temperature.
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Stevens and Flewitt also found that the evolution of the M6C phase was unaffected by stress. Nakatani et al. have developed correlations between the percent M6C and rupture-life reduction so that the percent M6C in a service-exposed component could be used to deduce the creep-life fraction consumed.22
23.3.8 Hardness-based techniques for creep damage assessment The changes in hardness in low-alloy steels as a function of time and temperature have been extensively quantified, so that hardness changes can be used to estimate the operating temperature.23,24 Correlations have also been established between tensile (hence hardness) and the Larson–Miller rupture relationships for low-alloy steels. These correlations enable selection of the appropriate Larson–Miller plot corresponding to a given hardness level, which can be used to calculate the remaining life,25 see Fig. 23.8. Goto26 and Kadoya et al.27 have proposed using hardness as a stress indicator in the remaining life assessment of CrMoV rotors26,27 used in steam turbines. They have quantified the effect of stress on the aging process so 2.2
60–70 ksi 70–75 ksi 75–80 ksi 80–90 ksi 90–100 ksi 100–110 ksi 110–115 ksi 115–120 ksi 120–130 ksi > 130 ksi
2.0 1.8 1.6
Log stress (ksi)
1.4 1.2 1.0 0.8 0.6
60 ksi 100 ksi 140 ksi
0.4 0.2 0 25 000
LMP = 40 975 + 57 (UTS) –5225 log σ–2450 (log σ)2 30 000
35 000 40 000 LMP = T (20 + log tr)
45 000
50 000
23.8 2.25 Cr–1Mo rupture data showing UTS dependence. ksi is 1000 psi (pounds per square inch).
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that by comparing the kinetics of hardness change in the rotor with that of thermally aged samples, the local stress can be determined. This value of stress and the known value of service temperature are used in conjunction with the Larson–Miller rupture plot to estimate tr for the material in its current damage state. In contrast, McGuire and Gooch show that the magnitude of the stress is unimportant28 once a ‘threshold stress level’ has been exceeded. For all stress levels in the range 70–240 MPa, the hardness change in the stressed condition was found to be 21% higher compared to the change in the unstressed condition. A creep model incorporating structural degradation as monitored by hardness changes has been proposed by Cane and Bissall.13 By equating the kinetics of hardness change to the kinetics of interparticle spacing changes, the decrease in the threshold back stress for creep was modeled. Substitution of the threshold back stress in the Norton creep law yields an expression for secondary creep rate ε˙ in terms of hardness changes. By integrating the ε˙ between t = 0 and t = tr, where tr is the time to failure defined in terms of the time to reach an arbitrarily chosen critical strain, the remaining rupture life is predicted. The model is currently based on limited data and involves numerous assumptions that can only be justified by further research.
23.3.9 Hardness and low-cycle fatigue life Considerable work has been carried out (26-23) in applying hardness to calculation of fatigue-life consumption in the groove regions of a CrMoV steel rotor used in steam turbines. It has been observed that low-cycle fatigue damage results in strain softening and can be measured as a hardness decrease. The premise, therefore, is that if the fatigue curve corresponding to the current hardness (in service) could be defined, the fatigue-life fraction consumed could be calculated by entering the appropriate total strain range ∆ε˙ t versus number of cycles to crack initiation (Nf) curve, as shown in Fig. 23.9.26 These relationships have been quantified.26
23.4
Accelerated destructive tests
One of the techniques widely used for life assessment of components involves removal of samples from a service component and conducting creep tests or stress rupture tests under accelerated conditions. Acceleration is achieved by increasing the stress or temperature or both. The data are extrapolated to service conditions using various parametric techniques reviewed in detail by Viswanathan.3 An example of commonly used procedures is briefly described here.
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Hv = 252 Hv = 252
Hv = 203 102
103 104 Number of cycles to crack initiation
105
23.9 Estimation of low-cycle fatigue properties by hardness for a Cr–Mo–V rotor forging at 500°C. Hv = Vickers hardness number.
23.4.1 Parametric extrapolation techniques Larson and Miller first introduced the concept of a time–temperature grouping in the form T(K + log t), where generally t is the time to rupture for a given material. A plot of stress versus the above parameter resulted in a single plot, within the limits of scatter, for any combination of time and temperature. A value of K = 20 is commonly used, although the constant is now known to have a range of material-specific values. For the purposes of the remaining life evaluation, specimens from the aged component are tested at higher stresses and temperature and the results are extrapolated to service conditions by plotting the stress versus the Larson–Miller parametric combination of temperature and time to rupture. Monkman and Grant29 found that, for many alloy systems, the relationship between the minimum creep rate, ε˙ , and time to rupture, tr, can be expressed by: log tr + m log ε˙ = constant
[23.6]
where m is a constant. For most materials the evaluated m has values approaching unity, so that Equation [23.6] can be rewritten as:
ε˙ tr = constant
[23.7]
This relationship enables an estimation of tr if the minimum creep rate can be determined either from accelerated tests or by creep or stress relaxation tests under service conditions. In general, the relationship between the minimum creep rate and stress is given by: ε˙ = Aσ n
[23.8]
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where A and n are stress-independent constants. Because creep is a thermally activated process, its temperature sensitivity would be expected to obey an Arrhenius-type expression, with a characteristic activation energy, Q, for the rate-controlling mechanism. Equation [23.8] can therefore be rewritten as:
ε˙ = A0 σ n exp – Q/ RT
[23.9]
Since ε˙ and tr are correlated through Equation [23.8], Equation [23.9] is usually applicable to express tr instead of ε˙ with the signs reversed for n and Q. This is the basis for the parameters defined by Orr et al.30 and is often used to extrapolate ε˙ or tr from one set of stress–temperature conditions to another. A very general description of the creep curve under constant stress conditions is given by the θ projection concept put forward by Evans et al.,31 in which creep strain, ε˙ , is considered to be the sum of two competing processes using the equation:
ε˙ = θ1[1 – exp(–θ2t)] + θ3[exp(θ4t) – 1]
[23.10]
where, θ1, θ2, θ3, and θ4 are all experimentally determined constants which are functions of stress and temperature: θ1 and θ2 define the primary or decaying strain rate component, and θ3 and θ4 describe the tertiary or accelerating strain rate component. The absence of a steady second-stage creep rate is implied by the model. A wide range of creep curve shapes can be modeled with various combinations of the constants. Analysis of extensive creep data on ferritic steels has shown that the log θ values vary systematically and linearly with stress. Hence, for a given material, if the θ functions can be defined on the basis of short-time tests at high stresses, then the values at lower stresses (longer times) can be obtained by extrapolation and the longtime creep curves under low-stress conditions can be readily predicted by substituting the θ values in Equation [23.10]. From the shape of the creep curve, the remaining life is estimated.
23.4.2 Isostress rupture tests This technique, widely used in Europe and the USA, is based on the conclusions stated earlier that the LFR is valid for temperature-based extrapolations and not for stress-based extrapolations. Rupture tests are conducted at elevated temperatures at a constant stress close to the service stress and the tr versus T data are extrapolated linearly to the service temperature to estimate the remaining life. Refinements of this technique include the use of miniature samples, consideration of oxidation and specimen size effects, verifications of the LFR, the applicability of uniaxial results to predict component behavior under multiaxial stresses and the effects of cycling. The use of miniature specimens 2 mm in diameter and 10 mm long instead of the conventional
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large-size specimens has been demonstrated. Based on the fact that large specimens (base metal) last two to three times as long as laboratory-size creep specimens tested in air, correction factors have been determined for applying laboratory data to predict the life of heavy section components. The applicability of the LFR under isotress test conditions has been verified. It has been shown that uniaxial test data can be used to make conservative estimates of the remaining life of tubes operating under biaxial loading. Selection of time–temperature combinations must be carried out judiciously. It has also been observed that the isotress extrapolation procedure using tr versus T yields the most conservative estimate of remaining life compared with tr versus 1/T and the Larson–Miller-type extrapolations, although contrary results have also been reported in the literature.
23.4.3 Small punch testing The small punch or disk bend test has particular value in life prediction of operating equipment since the test requires very small amounts of material (a typical specimen disk is 0.25 inches (6.35 mm) diameter × 0.020 inches (0.5 mm thick) and often the required volume of material can be acquired by operating equipment in a virtually non-destructive manner (see for example Foulds and Viswanathan).32 The application of the small punch (SP) test for creep has gained significant interest in the last decade, primarily as a result of research in Europe (see for example, a 2003 summary of the European round-robin testing).33 Most recently, the CEN (one of three European standardization organizations recognized by the EC) has been working to develop a code of practice for the small punch test. The code of practice, intended to achieve a practical level of uniformity in implementation of the test method, includes material creep testing in addition to the more mature application of the test for tensile and toughness properties. The practice documents are being developed as a workshop agreement and are nearing completion.34 The focus on development of the SP test for creep has been on the use of the small punch test to predict the time to rupture of a uniaxially loaded test specimen under specific conditions of applied load and temperature. This is because the creep life prediction for operating equipment has been based on extrapolation of conventional, uniaxially loaded test specimen rupture times to the field application in question (see for example Foulds and Viswanathan).35 In essence then, the current challenge of the use and interpretation of small punch test data boils down to answering the question of what is the load that should be applied to the small punch test specimen to produce a rupture time equal to that which would be produced in a uniaxial test specimen under the load or initial applied stress condition of interest. For example, to estimate the remaining life of a pressure vessel operating at a particular stress and
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temperature, we need to small punch test the as-removed vessel sample material at some load that would give us rupture times equal to those that a uniaxial test specimen would give us when subjected to the same stress as the pressure vessel. However, since the stress state in the SP test (biaxial) differs from that in a uniaxial test (uniaxial) and since the stress field in the small punch test varies through the test, interpretation is not straightforward. Research over the past decade has led to an encouraging semi-empirical interpretation of the test that could lead to field use. A summary of available correlations between the small punch test load and the ‘equivalent’32 uniaxial test specimen stress has been summarized,36 all of which ignore bending in the punch specimen, assuming the rupture time is predominantly driven by membrane stresses. The correlations have been reviewed and the European round-robin data regressed over a simplified general form to suggest the following usable equation.33
F / σ = 3.33 k SP R –0.2 rt1.2
[23.11]
where F is small punch test load, σ is the ‘equivalent’ uniaxial specimen stress, R is the small punch receiving die opening radius, r is the punch radius and t is the punch specimen thickness, and kSP is an empirical correlation factor that should be determined for given material and test specimen geometry/ dimensions. Application of the SP test requires initial testing to determine this correlation factor. The basic approach to using the small punch test for creep life prediction can follow the temperature-accelerated method popular for uniaxial testing (e.g., Foulds and Viswanathan)32 that is, following establishment of kSP and with knowledge of the stress and temperature of the operating equipment in question, a series of small punch tests at temperatures elevated above the equipment operating temperature may be run, each at a constant load determined from the above equation. The small punch rupture times may then be extrapolated to the operating temperature of interest on a temperature–log (rupture time) basis (common with low alloy steels) or the 1/temperature–log (rupture time) basis (sometimes used with austenitic stainless steels). Alternative acceleration methods using a combination of elevated stress and temperature may also be used wherein extrapolation is performed using the Larson–Miller parameter, although the temperature acceleration method is considered more reliable. As the SP test sees more field use and as the specimen and test configurations achieve better uniformity, we can expect that its application to creep life prediction will increase.
23.4.4 Stress relaxation testing The best short term (less than one week of testing) test is the stress relaxation test (SRT), with a constant displacement rate (CDR) test providing an
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independent measure of embrittlement.37 Three or four 20 h test runs can provide a plot of creep rate versus a time–temperature parameter. Using the Monkman–Grant relation, the rupture life at the desired stress and temperature can be estimated. Alternatively, the data may be used to compute the stress for a creep strain of 1% as a function of temperature. This may be converted to rupture life under the desired test conditions using the Gill–Goldhoff correlation. This approach is the one advocated generally, and is described in detail elsewhere.37
23.5
High temperature crack growth
All of the techniques described so far relate to life prediction from a crack initiation point of view. For heavy wall components, the initiation criteria have to be combined with crack growth data to perform a fracture mechanics analysis of the remaining life. Fatigue crack growth analysis procedures are well established. For creep crack growth and creep–fatigue crack growth, however, methodologies and data needed for analysis have been gathered only during the last few years.
23.5.1 Creep crack growth Extensive creep crack growth data pertaining to CrMo piping steels and CrMoV rotor steels have been collected, analyzed and consolidated.38,39 It has been observed that a crack-tip driving force parameter termed Ct, which takes time-dependent creep deformation into account, correlates much better with crack growth rates (da/dt) than the traditionally used elastic stress intensity factor K. The relation between da/dt and Ct can be expressed as:38–40 da = bC m t dt
[23.12]
Ct = σ ε˙ (A, n) aH (geometry, n)
[23.13]
and where σ is the stress far from the crack tip, obtained by stress analysis, ε˙ is the strain rate far from the crack tip, which is a function of the constants A and n in the Norton relation, a is the crack depth obtained from NDE measurements and H is a tabulated function of geometry and the creep exponent n. The values of A and n are either assumed from prior data or generated by creep testing of samples. By assembling all the constants needed, the value of Ct can be calculated. Once Ct is known, it can be correlated to the crack growth rate through the constants b and m in Equation 23.12. Combining Equations 23.12 and 23.13
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provides a first order differential equation for crack depth (a) as a function of time (t). Theoretically, this equation can be solved by separating variables and integrating. However, the procedure is complicated by the time dependency of Ct and the crack size dependence of the term H. To circumvent this, crack growth calculations are performed with the current values of da/dt (or a), to determine the time increment required for incrementing the crack size by a small amount ∆a (i.e. ∆t = ∆a/a). This provides new values of a, t and Ct, and the process is then repeated. When the value of a reaches the critical size ac as defined by KIC, JIC, wall thickness, remaining ligament thickness, or any other appropriate failure parameter, failure is deemed to have occurred. A number of variables affect the crack growth rate by modifying b, Ct or m and a large database is needed in order to apply the methodology. Among the variables to be considered are service degradation, presence of inclusions, impurities, ductility, test temperature, crack tip constraint and primary creep.
23.5.2 Creep–fatigue crack growth Major advances have been made in developing the methodologies and data needed to treat crack growth under the combined effects of creep and fatigue at elevated temperatures. The loading conditions in elevated temperature power generation components can often be simply represented by a trapezoidal wave shape consisting of a loading period, a hold time and an unloading period. For creep–fatigue crack growth, the total crack growth rate is then given by:
da da da dN total = dN fatigue + dN hold
[23.14]
da n q dN total = C ( ∆K ) + C1 [( C t ) ave ] t h
[23.15]
where the first term denotes the Paris law for the fatigue crack growth component and the second term combines the crack growth owing to creep, including that owing to stress relaxation. The average da/dt and Ct are obtained as follows: da = 1 da dt avg t h dN hold
[23.16]
and ( C t ) avg = 1 th
∫
th
[23.17]
C t dt
0
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The (da/dN) hold is the crack growth during the hold period and is obtained by subtracting the cycle-dependent crack growth rate from the total crack growth rate. Figure 23.10(b) shows a plot of (da/dt)avg against (Ct)avg for 1.25Cr– 0.5Mo steels at 538°C. The data include test results for a 90-s hold time, 10-min hold time and also the creep crack growth rate data. When plotted as a function of (Ct)avg, the time rates of crack growth for these very different conditions fall on the same trend curve. In contrast, when the crack growth per cycle is plotted as a function of ∆K (see Fig. 23.10(a)) different curves are obtained. The significance of the above trend with regard to predicting the hold-time effect in engineering components is obvious since creep crack growth data can be used to estimate creep–fatigue crack growth data and vice versa. In order to use this approach to predict crack growth during hold time it is necessary to estimate the magnitude of (Ct)avg for components. An equation has been proposed for estimating the (Ct) of any geometry for a material deforming by elastic–cyclic plasticity and power-law creep.40 The details of this equation are given by Yoon et al.40 and are outside the scope of this overview. The applicability of this approach to steels other than Cr– Mo steels is yet to be explored. The potential of time-dependent fracture mechanics (TDFM) in establishing the design life of new components or safe inspection intervals for components in service, or for performing risk assessments is obvious. The technology has come a long way in the past 20 years but much still remains to be done to develop total confidence in the approach. A majority of tests and analyses performed assume isothermal conditions in which the influence of environment is not explicitly included. More research into understanding creep–fatigue environment interactions is necessary for accurate life predictions. Considerable research is needed into analytical methods for treating crack growth under thermalmechanical loading and new test methods are needed that provide crack growth data under temperature gradients. The limitations of parameters such as ∆J under thermal gradients should be explored. An area that has not been explored much is that of load interactions during crack growth at elevated temperatures. In the presence of transient thermal stresses, it becomes quite important to treat the effects of overload on the crack growth rate during the subsequent hold time. There are significant opportunities for developing standard methods for creep–fatigue crack growth testing. These tests are highly specialized and require very precise controls and measurements. The data analysis is also complex so that forcing some uniformity into how data are treated will also help the overall goal of developing a well accepted life prediction methodology. Extension of these methods to directionally solidified alloys single crystal alloys and to intermetallics is needed. These materials can exhibit a range of behavior not seen in Cr–Mo ferritic and austenitic stainless steels. For example, depending on the loading conditions and orientation, the same alloy may
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40
10–2 1/10/1 (sec) 1/98/1 1/600/1 1/900/1 1/24h/1
10–3
1.25Cr–0.5 Mo 538°C (1000°F) 10–2
10–4
10–3
10–5
(b)
5 × 10–6 10
20 ∆K .(ksi inch) (a) 10–2
10–1
–2
10
KJ m–2 h–1 10–1 1
30
101
Trapezoidal waveshape 1/10/1 1/900/1 1/95/1 1/24/1 1/600/1 CCG
40
102 1
10–1
10–3
mm h–1
(da/dt)avg (inch h–1)
10–1
mm h–1
da/dN (inch cycle–1)
Trapezoidal waveshape
10–2 10–4
10–3 1.25Cr–0.5 Mo 538°C (1000°F)
10–5
10–4 (a) 10 –6 10–5
10–4 10–3 10–2 10–1 (Ct)avg (kips inch h–1) (b)
1
23.10 Comparison of creep–fatigue crack growth rates with (a) fatigue crack growth rate plotted against ∆K, (b) creep crack growth data against (Ct)avg.
behave as a creep-ductile or a creep-brittle alloy. Solving this problem will require good numerical simulations that are now well within the capability of the current technology. Monitoring of service experience is very important in determining which aspects of the problem deserve priority over others. Service experience is also important to validate the models after they are developed and implemented.
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23.5.3 Unique issues to weldments There are many issues unique to weldments that are often not taken into account in life prediction procedures using calculations or accelerated testing. High-temperature steam pipe seam weldment failures have occurred under operating conditions believed to be well within design, and at accumulated service hours far short of any lifetime predicted by the base metal rupture database (i.e. at very low expended life fractions). Premature failures could be attributed to inherently weaker weldment properties, strain concentrations arising from inhomogeneous creep properties, or to some combination of the two effects. Design databases for high-temperature components are generally based on data for homogeneous base metal tested as small specimens in uniaxial tests. None of these conditions apply to the behavior of heavy section weldments. The application of isostress rupture testing to a high-temperature welded structure typically involves consideration of weld configuration, time-dependent stress distribution, inhomogeneous creep deformation, creep ductility variation across the weldment, and so on. To add to the uncertainty of these effects, very little data are available on the stress rupture properties of weldments and limited data show that certain types of weldments can be inherently weaker compared to base metal owing to segregation of impurities and inclusions (see Viswanathan and Foulds).35 Typical weldments in operating components are subject to multiaxial stresses and this introduces yet another level of complexity to application of the method. Specific concerns with respect to welds are: • • • • • • • •
Welds are composite materials that contain many zones and pose many problems for the stress analyst. Failure can occur by crack growth at any of the interfaces, and it is hard to predict/simulate/accelerate the failures. Properties of welds vary widely with process. Multi-pass welds lead to through-thickness variability in properties. Sub-surface (type IV) cavities are hard to inspect. Strength mismatch, cusp angle and roof angle influence failure profoundly. Interfaces in the weldment are subject to carbide growth, depletion, embrittlement, and so on. It is hard to predict/model these phenomena. Crack growth data are limited and fracture mechanics methodology is underdeveloped.
23.6
Future trends
It is clear from the discussion in this chapter that there are many discrepancies in current technology for detecting, characterizing and quantifying creep and
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creep damage in components operating at high temperatures. The demands on creep damage assessment and its effect on the remaining life of components under creep will lead to increased use of on-line monitoring and decision making tools, in situ replication, microscopy, chemical analysis, and so on, which will be used on a more quantitative basis. The initiation and evolution of type IV cracking will be of great interest. Remote, rapid, wide coverage NDE techniques will be developed to perform quick and large scale surveys of damage. NDE techniques not requiring preparatory work (scaffolding, insulation removal, blade removal) will be developed. Improved signal processing and pattern recognition techniques will be developed better to characterize defects. Methods for non-destructive evaluation of in-service toughness, creep and fatigue damage will become more commonplace. Better tracking of operating history, improved databases on material properties and use of relevant TMF data will be utilized leading to improved calculational procedures and understanding of creep–fatigue. More realistic test procedures simulating service conditions need to be developed. Treatment of uncertainties in data will increasingly be treated by probabilistic methods. Miniature specimen (e.g. small punch test), stress relaxation and other improved destructive test techniques will be of greater interest. The effects of weld procedures, weld geometry of PWHT, optimized fillers, mismatch effects, and other variables in weld performance in creep will be more thoroughly investigated. Weld strength reduction factors will be defined and the results applied to design as well as to life prediction. Creep and creep–fatigue crack growth methodologies, especially with focus on weldment will be developed. It is difficult to prioritise among these many areas competing for R&D dollars. In general, going from calculational to NDE to destructive assessment involves increasing cost but also increasing accuracy. Which ones of this three-phase approach a utility company might choose to implement would depend on the age of the plant, budget considerations, operation and maintenance policies and many other factors. It is often the case that a company has only a given budget and a decision has to be made about which actions will provide maximum ‘bang for the buck’. A probabilistic approach, combined with a sensitivity analysis might help determine where the maximum value for money could be obtained. In conclusion, it is reasonable to say that substantial progress has been made during the last decade in condition assessment of in service components. Some more work remains to be done, as reviewed in this chapter.
23.7
References
1 Robinson E L, ‘Effect of temperature variation on the creep strength of steels’, Trans. ASME, 1938, 160, 253–259.
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2 Lieberman Y, ‘Relaxation, tensile strength and failure of E1 510 and Kh1 F-L steels’, Metalloyed Term Obrabodke Metal, 1962, 4, 6–13. 3 Viswanathan R, Damage Mechanisms and Life Assessment of High Temperature Components, ASM International, Metals Park, 1989. 4 Viswanathan R, ‘Low cycle fatigue life prediction in LCF 3’, Low Cycle Fatigue and Elastic Behavior of Materials, K T Rie, Gounling H W, Konig G, Neumann P, Nowak H, Schwalbe K H and Seeger T, Elsevier, UK, 1992, 695–721. 5 Viswanathan R, ‘Creep fatigue life prediction of fossil plant components in creep’, Fatigue, Flaw Evaluation and Leak Before Break Assessment, Y S Garud (ed.), ASME, PVP, New York, 1993, volume 266, 33–51. 6 Priest R H, Beauchamp D J and Ellison E G, ‘Damage during creep-fatigue’, in Advances in Life Prediction Methods, ASME Conference, Albany, American Society of Mechanical Engineers, 1983, 115–222. 7 Miller, D A, Priest R H and Ellison E G, ‘A review of material response and life prediction techniques under fatigue–creep loading conditions’, High Temp. Mater. Proc., 1984, 6 (3 and 4), 115–194. 8 Viswanathan R and Bernstein H, J. ‘Some issues in creep–fatigue life prediction of fossil power plant components’, Trans. Indian Inst. Metals, 2000, 59 (3), 185– 202. 9 Kuwabara K, Nitta A and Kitamura T, Advances in Life Prediction, Ed., D A Woodford and Whitehead R (eds), ASME, New York, 1985, 141–152. 10 Bisbee H and Nottingham, L, December ‘Longitudinal seam weld characterization by focused ultrasonics’, SPIE Proceedings, 1996, 2947, 88–99. 11 Sablik M J and Augustyniak B, 2000, ‘Modeling the magnetic field dependence of magnetoacoustic emission and its dependence on creep damage, Review of Progress in Quantitative Non-destructive Evaluation, Thompson D O and Chimenti D E (eds), American Institute of Physics, New York, Volume 19B, 1557–1564. 12 Govindaraju M R, Kaminski D A, Devine M K, Biner S B and Jiles D C, ‘Nondestructive evaluation of creep damage in power plant steam generators and piping by magnetic measurements’, NDT & E Inte., 1997, 30, 11–17. 13 Cane B J and Bissall A M, ‘Predictive assessment of damage in elevated temperature weldment’, paper presented at the EPRI Plant Maintenance Technology Conference, Houston, November 14–16, 1986. 14 EPRI, Acoustic Emission Monitoring of High-Energy Steam Piping, Volume 1: Acoustic Emission Guidelines for Hot Reheat Piping, EPRI Report TR-105265-V1, November 1995. 15 Neubauer B and Wedel V, ‘Restlife Estimation of Creeping Component By Means of Replication’, D A Woodford and R Whitehead (eds), Advances in Life Prediction, ASME, New York, 1983, 307–314. 16 Cane B J and Shamma M, A Method for Remanent Life Estimation By Quantitative Assessment of Creep Cavitation on Plant, Report TPRD/L2645/N84,?, UK. As cited by Viswanathan et al. 1994, ‘Life Assessment Of Superheater/Reheater Tubes In Fossil Boilers’, J. Pressure Vessel Technol., 1994, 1, 59–75. 17 Ellis F and Viswanathan R, ‘Review of Type IV Cracking’, Fitness for Service Evaluations, Volume 380, ASME PVP, 1998, 59–75. 18 Wells C and Viswanathan R, ‘Life Assessment of High Energy Piping’, in Technology for the 90s, Au Yang M K (ed.) ASME Pressure Vessel and Piping Division, New York, 179–216. 19 Stevens R A and Lonsdale D, Isolation and Quantification of Various Carbide Phases
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21 22
23
24
25
26
27
28
29
30 31
32 33
34 35
36 37
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in 2.25Cr–1Mo Steel, SER/SSD/-84-0046/N, Central Electricity Generating Board, England, June 1984. Stevens R A and Flewitt P E J, The Effect of Phosphorus on the Microstructure and Creep Strength of 2.25Cr–1Mo Steel, SER/SSD/84-0020/R, Central Electricity Generating Board, England, March 1985. Munson R, Radian Corporation, Austin, TX, 1990, private communication. Nakatani H, et al., ‘Metallurgical Damage Detection and Life Evaluation System for Boiler Pressure Parts’, Paper presentated at the EPRI Conference on Predictive Maintenance of Fossil Plant Components, Boston, MA, October 1990. Askins M C, Remaining Life Estimation of Boiler Pressure Parts, Vol. 4, Metallographic Models for Weld Heat Affected Zones, Report CS-5588, Electric Power Research Institute, Palo Alto, CA, November 1989. Ellis F V, Robert B W and Henry J F, Remaining Life Estimation of Boiler Pressure Parts, Vol. 4, Metallographic Models for Weld Heat Affected Zones, Report CS5588, Electric Power Research Institute, Palo Alto, CA, November 1989. Grunloh H and Ryder R H, Life Assessment of Boiler Pressure Parts, Vol. 7, Superheater, Reheater Tubes, Report TR-103377, Vol. 7, Electric Power Research Institute, Palo Alto, CA, November 1993. Goto T, Study on Residual Creep Life Estimation Using Nondestructive Material Property Tests, ‘Mitsubishi Technical Bulletin, No. 169, Mitsubishi Heavy Industries, Takasago, April 1985. Kadoya Y, Goto T, Uaeke M and Fujii H, 1985 ‘Material Characteristics NDE System from High Temperature Rotors’, Paper No. 85-JPGC, PWR-10 presented at the ASME/IEEE Joint Power Generation Conference. McGuire J and Gooch D J, ‘Metallographic Techniques for Residual Life Assessment of 1CrMoV Rotor Forgings’, Proceedings of the International Conference on Life Assessment and Extension, Nederlands Instituut Voor Lastichniek, The Hague, 1989, Volume 2, p 116. Monkman F C and Grant N J, ‘An Empirical Relation Between Rupture Life and Minimum Creep Rate in Creep Rupture Tests’, Proc. ASTM 56, 1938, 1956, 593– 620. Orr R L, Sherby O D and Dorn J E, Trans. ASM, 1954, 46, 113. Evans R W, Parker J D and Wilshire B, ‘An extrapolative procedure for long-term creep strain and creep life prediction’, In Recent Advances in Creep and Fracture of Engineering Materials and Structures, Pineridge Press, Swangear, 135–184. Foulds J and Viswanathan R, ‘Nondisruptive Material Sampling and Mechanical Testing’, J. Nondestructive Evaluation, 1996, 15 (3), 151–162. Bicego V, Di Persio F, Hurst R and Ranta J H, ‘Small Punch Creep Test Method: Results from A Round Robin carried out within EPERC TTF5’, 29th MPA Seminar, Stuttgart, 9–10 October 2003. CEN Workshop Agreement WS21, Small Punch Test Method for Metallic Materials, CEN, Brussels Belgium, 2004. Foulds J and Viswanathan R, ‘Accelerated Stress Rupture Testing for Creep Life Prediction – Its Value and Limitations’, J. Pressure Vessel Technol., May 1998, 120, 105–115. Li Y and Sturm R, ‘Small Punch Tests for Welded Heat Affected Zones’, International Conference on Welds 2005, Geesthacht, September 2005. Woodford D.A. ‘Creep strength evaluation of serviced and rejuvenated T91 using the stress relaxation method’, in Advances in Materials Technology For Fossil Power
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Plants, Chicago, American Society For Metals ASM, Metals Park OH, 2004, 1101– 1115. 38 Saxena A, Han J and Banerji K, Creep Crack Growth in Boiler and Steam Pipe Steels, Report CS-5583, Electric Power Research Institute, Palo Alto, CA, January 1988. 39 Saxena A, Creep Crack Growth in CrMoV Rotor Steels, EPRI RP2481-5 Report, Palo Alto, CA. 40 Yoon K B and Saxena A ‘Characterization of Creep-Fatigue Crack Growth Behavior Under Trapezoidal Waveshape Using Ct Parameter’, Int. J. Fracture, 1993, 59, 95– 102.
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A-USC plants 565–70 accelerated destructive tests 653–8 acceleration of cyclic creep 375, 384 acoustic emissions (AE) monitoring 647–8 activation energies 241–2 advanced turbines 174 age hardening 158 ageing embrittlement 611–13 alloy design 539–70 austenitic steels 45–8, 554–70 Fe-Ni based austenitic alloys 565–70 for header pipes 541 for heavy-wall thickness piping 546–54 high strength low-Cr steels 541–6 for main steam pipes 541 martensitic high-Cr steels 546–54 and oxidation 527–8 for power plant components 539 for reheaters 540–1, 554–64 for superheaters 540–1, 554–64 for water walls 539–40 analytical models of crack formation 512 anelastic deformation 371, 372 anomalous temperature dependence 528 arc welding 476 argon-oxygen decarburization (AOD) 66 ASME Code 158, 169 ASME P91 steel 30, 539 athermal yield stress 265, 270–5, 277, 359–61 atomic bonding 224–5, 226–7 austenitic steels 23, 42–64, 68–70 15%Cr-15%Ni 51–2 18%Cr-8%Ni 43, 44, 51
20-25% Cr 44, 52–3 age hardening 158 alloy design 45–8, 554–70 austenitising temperature 23 boiler tube applications 48–57 chemical plant applications 62–4 cold working 286 constant stress creep 368–9 cooling from the annealing temperature 158 Discaloy 60, 61 electrical conductivity/resistivity 236–7 electro slag remelting (ESR) process 61, 66 grain boundaries 341–5 heat exchanger applications 48–57 low initial dislocation density 379–80 microstructure 366 nickel alloyed 42–3 in nuclear reactors 615–21 and radiation swelling 599–600 strengthening mechanisms 292–5 Tempaloy A-1 steel 44 thermal conductivity 231 for thick-section pipe 57–62 TP 316 steel 57, 59 for turbine components 57–62, 591 under-stabilising technique 47 void swelling 599–600 Young’s moduli 224 austenitising temperature 23 back-stress concept 322–4 bainitic low Cr steels 287–9 boiler tube applications 48–57
667 WPNL2204
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Index
bolting materials 142 Boltzmann constant 345, 346 bond energy 226–7 bond length 227 bond strength 227 boron 38, 329–30, 336–7, 342, 345 Brite-Euram-Projects 80 brittle failure 25 calculational methods of damage assessment 638–43 cumulative damage 638–9 ductility exhaustion 639–40 lacunae 640–1 life-fraction rule 638, 641–3 linear damage rule 639 strain-fraction rule 638 carbide analysis 651–2 carbon steels 19 casting process 174, 175–82 12Cr steels 199–200 electric arc furnaces (EAF) 175, 178–80 electroslag remelting (ESR) 61, 66, 175, 180–1 hot topping process 181–2 ladle refining furnaces (LRF) 66, 175, 178–80 open hearth furnaces 64, 66, 175 vacuum arc remelting (VAR) 182 vacuum degassing 66, 175 vacuum induction melting (VIM) 175, 182 cavity growth 352–5, 648–51 cavity surface diffusion 354–5 CB8 alloy 306–12, 315–20, 324–5 CCT (continuous cooling transformation) diagrams 188–9 CEN standards 81, 95–150 bolting materials 142 component creep tests 147 design standards 148 general principles 95 material specifications 95 relaxation testing 147–8 residual life assessment 149–50 testing programmes 147 testing techniques 144–6 user specifications 149
welding consumables 142 welding procedure qualifications 142, 144 chemical composition 155 chemical plant applications 62–4 chromium 10-12% Cr steels 32–6 11% Cr steels 30, 36–8 12Cr steels 199–200, 201 15%Cr-15%Ni system steels 51–2 18%Cr-8%Ni austenitic steels 43, 44, 51 20-25% Cr steels 44, 52–3 9-12% Cr steels 26–32, 45, 67–8, 290–2, 295–301 and grain boundaries 330 CrMoV steels 21, 25, 26–7, 192–5 and oxidation 527 and turbine components 192–204 classical nucleation theory (CNT) 313–14 coatings for oxidation-resistant steels 533 cobalt diffusion 250 Coble creep 10, 275, 276 cold rolling 284–5, 286 component creep tests 147 composite model 388–9 COMTES projects 211 conductivity electrical 234–8 thermal 230–4 constant stress creep 368–70 constituent atoms 244–5 constitutive equations 403–17 data distribution 412 defect assessment 416 end-of-life criteria 414–15 future trends 416 grain boundaries 345–6 and material characteristics 412 model fitting effectiveness 405–11 model selection 412 multi-axial stress rupture 415 service life predictions 412–16 continuous cooling transformation (CCT) diagrams 188–9 cooling from the annealing temperature 155, 158 core components 615–17, 621–4 COST programmes 79–80, 306–12, 584
WPNL2204
Index CB8 alloy 306–12, 315–20, 324–5 costs to industry 637 crack formation 459–68, 484–95, 504, 643 analytical models 512 creep crack growth 658–9 creep-fatigue crack growth 659–61 detectability 494–5 in dissimilar welds 488–90 growth rate equation 509 high temperature crack growth 658–62 microscopic features 509–13 reheat cracking 486–8 role of constraint 493–4 type I cracks 484, 485 type II cracks 484, 485 type III cracks 484–5, 486–8 type IV cracks 485, 491–5, 513–17, 649 V-type cracks 509–11 W-type cracks 509–11 in welded joints 484–95, 513–17 see also fracture mechanics creep crack growth 658–9 creep curve shape variations 429–30 creep deformation 637–63 athermal yield stress 265, 270–5, 277 calculational methods of assessment 638–43 cavity evolution 352–5, 648–51 Coble creep 10, 275, 276 costs to industry 637 damage evolution 5 definition 3 deformation mechanism map 9–11, 275–7 fracture mechanism map 11–13, 350–64 homologous temperature 3–4 logarithmic creep 6 microstructure evolution 5 minimum creep rate 5, 243 modes of deformation 265, 637 Nabarro-Herring creep 10, 275, 276 necking of specimens 4–5 NIMS Creep Data Sheets 9 non-destructive evaluation 643–53, 663
669
NRIM Creep Data Sheets 9 and oxidation 7, 519–34 rate curves 6–7, 265 rupture strength 7–9, 16–17, 279 stages of 3–4 stress-strain responses of materials 265–7 temperature and strain rate dependence 267–9 threshold stress 78 in welded joints 483–4 yield stress 265, 267–9, 270–5, 277 creep ductility 298, 351–2, 431–3, 639–40 creep strain curves 403 creep-fatigue behaviour 446–69 crack behaviour 459–68, 659–61 experimental procedures 447–9 fatigue stresses 217–18 life estimation 449–56 multiaxial behaviour 456–9 Nelder-Mead method 459 stress-strain behaviour 449 creep-induced strain 422–7 CrMoV steel 192–5 cross gliding 248 crystallographic slip 365 cumulative damage calculation 638–9 cyclic creep 374–5, 382 acceleration 375 deceleration 374–5 cyclic variation of stress 384 damage assessment calculations 638–43 damage evolution 5 damage rules 638–9 damage tolerance values 430–1 data distribution in constitutive equations 412 data rationalization 425–6 DBTT (ductile to brittle transition temperature) 337–8, 605–10 deceleration of cyclic creep 374–5, 384 defect assessment 416 defect development and oxidation 523–5 deformation mechanism map 9–11, 275–7 degradation of particle hardening 383–4 delta ferrite 200, 478
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Index
design standards 148 detectability of crack formation 494–5 deuterium 625 development of creep-resistant steels 15–70 austenitic steels 23, 42–64, 68–70 ferritic steels 19–42 purity of heat-resistant steels 64–7 requirements for heat-resistant steels 18–19 steel melting 64–7 differential heat treatment 205, 207 diffusion behaviour 241–63 activation energies 241–2 constituent atoms 244–5 cross gliding 248 data application and searches 260–3 dislocation climb 248 dislocation glide 249 and failure 250 glide velocity 249 grain boundary diffusion 252–3 interdiffusion reactions 245 interdiffusion values 258 in iron and iron-base alloys 255–60 lattice diffusion 245–6 magnetic transformation 250–2 matrix deformation 248–9 microstructure change 249–50 oxygen diffusion 250 segregation of minor elements 253–4 short-circuit diffusion 243, 246–8 solute atoms 249 stacking-fault energy 254–5 time-dependent deformation 242–3 tracer diffusion values 257 vacancies 243–4, 260 DIN standards 79 Discaloy 60, 61 dislocation climb 248 glide 249 hardening 284–6, 541 models 385–9 precipitate-dislocation interaction 320–1 strengthening 541 dispersion hardening 281–4 dissolution of fine carbonitrides 296–8
ductile-to-brittle transition temperature (DBTT) 337–8, 605–10 ductility exhaustion 639–40 ductility of materials 351–2 durability strength 15–16, 17 DVM creep rate limit test 15–16, 19 dynamic recrystallisation 330, 334–6 ECCC see European Creep Collaborative Committee (ECCC) 18%Cr-8%Ni austenitic steels 43, 44, 51 elastic behaviour 219–25 modulus of elasticity 221–5 stress and strain 219–21 elastic deformation 371 electric arc furnaces (EAF) 175, 178–80 electrical conductivity/resistivity 234–8 electro slag remelting (ESR) process 61, 66, 175, 180–1 11% Cr steel 30, 36–8 embrittlement caused by ageing 611–13 end-of-life criteria 414–15 engineering grain boundaries 329 enhanced steam oxidation 520–5 EPERC see European Pressure Equipment Research Council (EPERC) estimation of damage 638–43 Eurofer ODS 339 Euronorms standards 80 European Carbon and Steel Collaboration 79 European Commission 81–2 European Creep Collaborative Committee (ECCC) 79, 80, 85–92 contribution to standardisation 90–2 future of 92 Memorandum of Understanding (MoU) 85 motivation and history 85 organisation 90 Working Groups 90, 91–2 European Pressure Equipment Research Council (EPERC) 80, 92–5 contribution to standardisation 94–5 history 92, 94 objectives 92, 94 organisation 94
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Index European specifications see specifications and standards (Europe) fabrication and joining 629–31 failure and diffusion behaviour 250 fast breeder reactors (FBRs) 598, 613–24 welded joints 630 fatigue behaviour see creep-fatigue behaviour FATT (fracture appearance transition temperature) 207 Fe-Ni based austenitic alloys 565–70 FEM (finite element method) analysis 185 ferrite effects on strengthening mechanisms 300–1 ferritic steels 19–42 10-12% Cr steels 32–6 11% Cr steel 30, 36–8 56T5 steel 27 9-12% Cr steels 26–32, 45, 67–8, 290–2, 295–301 boron additions 38, 329–30, 336–7, 342, 345 carbon steels 19 constant stress creep 369–70 cooling from the annealing temperature 155 CrMoV steel 21, 25, 26–7, 192–5 dislocation strengthening 541 FV448 steel 27 grain boundaries 330–40 H46 steel 27 HCM2S steel 25 HCM 12 steel 30–1, 32 high initial dislocation density 380–1 laths 331 low alloy steels 19–26 microstructure 366–7, 474–82 molybdenum steels 20 in nuclear reactors 621–4, 626–9, 631 oxide dispersion strengthened (ODS) 292, 338–40, 631 P91 steel 30, 539 rotor steels 30, 31 STX 21 research project 38–42 TAF steel 27 thermal conductivity 231
671
for thick-section boiler components 38–42 TMK 1 and 2 steels 31 ferromagnetic state 250 ferromagnetic transformation 252 15%Cr-15%Ni system steels 51–2 56T5 steel 27 finite element method (FEM) analysis 185 fission reactions 597–8 forging process 174, 183–6 12Cr steels 200 fracture appearance transition temperature (FATT) 207 fracture mechanics 504–17 effect of mechanical constraint 507–9 microscopic 509–13 non-linear 504–7 see also crack formation fracture mechanism map 11–13, 350–64 athermal yield stress 359–61 cavity growth 352–5 ductility of materials 351–2 modes of fracture 351 multi region analysis 361–2 rupture strength changes 356–8, 359–61 and microstructural degradation 358–9 stress and temperature dependence 352–5 fusion reactors 625–7 fusion welding 474–5 FV448 steel 27 German specifications and standards 78–9 Gibbs equilibrium segregation 333–4 Gibbs free energy 312, 313 glide velocity 249 Gr.91 steel 433–6 Gr.92 steel 436–40 grain boundaries 329–47 austenitic steels 341–5 and boron 329–30, 336–7, 342, 345 cavity growth controlled by 352–4 and chromium behaviour 330 constitutive equations 345–6 diffusion 252–3
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Index
ductile to brittle transition temperature 337–8 dynamic recrystallisation 330, 334–6 engineering 329 ferritic steels 330–40 future trends 348–9 and hafnium 338, 342, 345, 347 intergranular cracking 330, 353 molecular dynamics (MD) modelling 347 precipitate-grain boundary interaction 322 precipitate-subgrain boundary interaction 321–2 precipitation 329, 331–3, 341 properties 345–6 segregation effects 329–30, 333–4, 341–2 subgrain boundary migration 393 triple junction mobility 347 vacancy concentration 354 and zirconium 342, 345 grain growth zone 479 grain refined zone 479 grain refining behaviour 188 H46 steel 27 hafnium 338, 342, 345, 347 Hall-Petch relationship 330 hardening see strengthening mechanisms hardness-based testing techniques 652–3 HCM2S steel 25 HCM 12 steel 30–1, 32 header pipes 541 heat affected zone (HAZ) 472, 478–80, 497–8, 513–17 HAZ simulation 480–2 heat exchanger applications 48–57 heat indication test (HIT) 192 heat transfer mechanisms 231 heat treatment process 174, 186–90 12Cr steels 201 continuous cooling transformation (CCT) diagrams 188–9 differential heat treatment 205, 207 grain refining behaviour 188 normalizing heat treatments 187–8 pearlite transformation 188 preliminary heat treatments 187–8
quality heat treatment 188–90 quenching 188, 189–90, 201 stress relief treatment 190 tempering 190 heavy-wall thickness piping 546–54 helium embrittlement 610–11 high initial dislocation density 380–1 high pressure-low pressure combination (HLP) rotors 204–7 high strength low-Cr steels 541–6 high temperature crack growth 658–62 history of creep-resistant steels see development of creep-resistant steels homologous temperature 3–4 Hooke’s Law 221 hot topping process 181–2 Hull-Rimmer equation 345 in situ TEM observations 389–93 ingot making process 174, 182–3 interdiffusion reactions 245 interdiffusion values 258 intergranular cracking 330, 353 iron and iron-base alloys diffusion behaviour 255–60 irradiation creep 602–5 irradiation embrittlement 605–10 isostress rupture tests 655–6 Italian specifications and standards 79 Japanese specifications see specifications and standards (Japan) JIS Code 158–69 JSME Code 158 kinetics of dislocation glide 385–7 kinetics model 312–13 Kirkendall effect 245 lacunae in calculational methods 640–1 ladle refining furnaces (LRF) 66, 175, 178–80 Larson-Miller parameters 361, 567 laths 331 lattice diffusion 245–6 life assessment see service life life extrapolation in Gr.91 steel 433–6 life extrapolation in Gr.92 steel 436–40
WPNL2204
Index life-fraction rule 638, 641–3 linear damage rule 639 lithium 625 logarithmic creep 6, 405 longitudinal seam welds 472 Lorenz number 236 low alloy steels 19–26 low initial dislocation density 379–80 low-cycle fatigue life 653 low-temperature tempering 300 machining 190 magnetic acoustic emissions 646–7 magnetic testing 645 magnetic transformation 250–2 main steam pipes 541 martensitic steels 155, 224 high-Cr 546–54 tempered 9-12Cr steels 290–2, 381 for turbine components 591 material characteristics 412 material specifications 95 matrix deformation 248–9 mechanical constraint 507–9 mechanical properties 155, 169, 173 mechanical tests 190–2 mechanisms of creep 365–402 austenitic steel microstructure 366 composite model 388–9 constant stress creep 368–70 crystallographic slip 365 cyclic creep 374–5, 382 cyclic variation of stress 384 degradation of particle hardening 383–4 dislocation models 385–9 equations 375–6 evolution of dislocation structure model 387 ferritic steel microstructure 366–7 in situ TEM observations 389–93 kinetics of dislocation glide 385–7 microstructural interpretation of creep rate 375–84 microstructural model 401–2 particle hardening 387–8, 393–4 primary creep and loading strain 379–81 solid solution hardening 376–8
673
stress change responses 370–3, 381–2 subgrain coarsening 382–3 tertiary creep 382–3 transmission electron microscope (TEM) observations 389–93 velocity of glide 394 Metallographic Atlas 9 metallurgical tests 190–2 METI Code 158–69 microscopic fracture mechanics 509–13 microstructural model 401–2 microstructure of austenitic steels 366 changes 249–50 degradation 358–9 evolution 5 of ferritic steels 366–7, 474–82 interpretation of creep rate 375–84 and precipitation hardening 306–12, 320–2 Microstructure Data Sheets 9 Miner rule 451 minimum creep rate 5, 243 modelling bolt relaxation testing 148 in complex systems 312–15 composite model 388–9 crack formation 512 dislocation 385–9 evolution of dislocation structure model 387 fitting effectiveness 405–11 microstructural model 401–2 molecular dynamics (MD) modelling 347 precipitation kinetics model 312–13 selection 412 modes of deformation 265, 637 modes of fracture 351 modulus of elasticity 221–5 molecular dynamics (MD) modelling 347 molybdenum steels 20 Monkman-Grant relationship 6, 346 monotonic creep 382 multi region analysis 361–2 multiaxial behaviour 456–9 multiaxial stress rupture 415 multilayer welding 480
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Index
MX carbonitrides 286, 553 Nabarro-Herring creep 10, 275, 276 national standardisation bodies 81, 84–5 necking of specimens 4–5 Nelder-Mead method 459 neutron absorption cross-section 598 nickel steel 42–3, 232 NIMS Creep Data Sheets 9 9-12% Cr steels 26–32, 45, 67–8, 290–2, 295–301 non-destructive evaluation 643–53, 663 non-equilibrium segregation 334 non-linear fracture mechanics 504–7 normalizing heat treatments 187–8 Norton’s law 6 notch weakening 17, 581 NRIM Creep Data Sheets 9 nuclear reactors 597–632 and austenitic steels 615–21 core components 615–17, 621–4 embrittlement caused by ageing 611–13 fabrication and joining 629–31 fast breeder reactors (FBRs) 598, 613–24 and ferritic steels 621–4, 626–9, 631 fission reactions 597–8 fusion reactors 625–7 helium embrittlement 610–11 irradiation creep 602–5 irradiation embrittlement 605–10 neutron absorption cross-section 598 radiation damage 598–611 radiation swelling 599–602 specifications and standards 79 steam generators 627–9 thermal nuclear reactors 598 turbines 627–9 types of reactors 613 nucleation of precipitates 313–14 ODS steels 292, 338–40, 631 oil quenching 190, 192 open hearth furnaces 64, 66, 175 Orowan stress 272, 281–4, 388 Orr-Sherby-Dorn (OSD) method 361–2 Ostwald ripening 383–4 over-tempered region 480
overlay welding 201–2 oxidation 7, 519–34 acceptable rates of oxidation 528–9 and alloying additions 527–8 anomalous temperature dependence 528 and chromium content 527 defect development 523–5 hydrogen production 520 mechanisms of enhanced steam oxidation 520–5 rates of steam oxidation 525–9 and service life 530–2 spalling of oxide scales 523–5, 530 stability of oxides 520–1 stages of steam oxidation 521–3 testing steam oxidation resistance 519–20 voids and gap formation 525 oxidation-resistant steels 532–3 coatings 533 composition 532–3 surface modifications 533 oxide dispersion strengthened (ODS) 292, 338–40, 631 oxygen diffusion 250 P91 steel 30, 539 parametric extrapolation techniques 654–5 and strain analysis 423–5 partially transformed zone 479 particle hardening 383–4, 387–8, 393–4 degradation 383–4 patterns of strain accumulation 427–33 pearlite transformation 188 PED (Pressure Equipment Directive) 80, 82–3 phased array ultrasonic testing (UT) 645 physical and elastic behaviour 217–23 elastic behaviour 219–25 electrical resistivity and conductivity 234–8 fatigue stresses 217–18 thermal properties 225–34 thermal stress parameter 217–18 piping and tubing steels 158 plastic deformation 265, 266, 267, 278, 371
WPNL2204
Index Poisson’s ratio 220 post weld heat treatment (PWHT) 543 power law behaviour 426–7 power plant components 539 see also nuclear reactors; turbines power-law creep 507 precipitation at the grain boundary 329, 331–3, 341 precipitation hardening 281–4, 305–26, 476–7 back-stress concept 322–4 effects of precipitates 305 kinetics model 312–13 loss of strengthening 324–5 microstructure analysis 306–12 microstructure-property relationships 320–2 modelling in complex systems 312–15 nucleation of precipitates 313–14 precipitate evolution 307–12, 315–20 precipitate-dislocation interaction 320–1 precipitate-grain boundary interaction 322 precipitate-subgrain boundary interaction 321–2 thermodynamic equilibrium analysis 315–16 see also Z-phase precipitation precipitation of new phases 296–8 preferential recovery of microstructure 298 preferred absorption glide (PAG) creep 603 preliminary heat treatments 187–8 Pressure Equipment Directive (PED) 80, 82–3 primary creep 4, 405 and loading strain 379–81 producer certifications 83 production of steels for turbines 174–214 12Cr steels 195–204 casting process 174, 175–82 CrMoV steel 192–5 forging process 174, 183–6 future trends 208–12 heat treatment process 174, 186–90 high pressure-low pressure combination (HLP) rotors 204–7
675
ingot making process 174, 182–3 machining 190 testing and non-destructive examination 190–2 progressive coarsening 298 purity of heat-resistant steels 64–7 quality heat treatment 188–90 quenching 188, 189–90, 201 radiation damage 598–611 helium embrittlement 610–11 irradiation creep 602–5 irradiation embrittlement 605–10 radiation swelling 599–602 rate of creep curves 6–7, 265 rates of steam oxidation 525–9 reheat cracking 486–8 reheaters 540–1, 554–64 relaxation testing 147–8 requirements for heat-resistant steels 18–19 residual life assessment 149–50 resistivity 234–8 role of constraint 493–4 rotor steels 30, 31 rotors 22, 578 welded design 590 rupture life 352–5 rupture strength 7–9, 16–17, 279, 356–61 and microstructural degradation 358–9 rupture tests 16–17, 145 secondary creep 4, 5, 405 segregation effects 329–30, 333–4, 341–2 segregation of minor elements 253–4 service life constitutive equations 412–16 estimation 449–56, 662 and oxidation 530–2 rupture life 352–5 of turbine components 583–91 shear stress 220–1 short-circuit diffusion 243, 246–8 small punch testing 656–7 solid solution hardening 279–81, 376–8 solute atoms 249 spalling of oxide scales 523–5, 530
WPNL2204
676
Index
specifications and standards (Europe) 78–151 Brite-Euram-Projects 80 CEN standards 81, 95–150 COST programmes 79–80, 306–12, 584 DIN standards 79 Euronorms standards 80 European Carbon and Steel Collaboration 79 European Commission 81–2 European Creep Collaborative Committee (ECCC) 79, 80, 85–92 European Pressure Equipment Research Council (EPERC) 80, 92–5 future trends 150–1 in Germany 78–9 in Italy 79 national standardisation bodies 81, 84–5 non PED creep applications 83–4 and nuclear power 79 Pressure Equipment Directive (PED) 80, 82–3 producer certifications 83 Thermie-Projects 80, 208, 211 and turbine manufacturers 84 in the UK 79 specifications and standards (Japan) 155–73 ASME Code 158, 169 JIS Code 158–69 JSME Code 158 METI Code 158–69 piping and tubing steels 158 steam turbine steels 169 super alloy steels 169 types of Japanese steels 155, 158 Srolovitz mechanism 283, 284 stability of oxides 520–1 stacking-fault energy 254–5 stages of creep deformation 3–4 stages of steam oxidation 521–3 steam generators 627–9 see also turbines steam oxidation see oxidation steel melting 64–7
strain analysis 421–42 appraisal of data analysis 440–1 creep curve shape variations 429–30 creep ductility 431–3 creep strain curves 403 creep-induced strain 422–7 damage tolerance values 430–1 data rationalization 425–6 future trends 441–2 life extrapolation in Gr.91 steel 433–6 life extrapolation in Gr.92 steel 436–40 parametric approaches 423–5 patterns of strain accumulation 427–33 power law behaviour 426–7 practical implications 433–41 see also stress-strain behaviour strain monitoring 647 strain-fraction rule 638 strengthening mechanisms 279–301 austenitic steels 292–5 bainitic low Cr steels 287–9 dislocation hardening 284–6, 541 dispersion hardening 281–4 dissolution of fine carbonitrides 296–8 ferrite effects 300–1 loss of creep ductility 298 loss of strengthening mechanisms 295–301, 324–5 low-temperature tempering 300 oxide dispersion strengthened (ODS) 338–40 precipitation hardening 281–4, 305–26, 476–7 precipitation of new phases 296–8 preferential recovery of microstructure 298 progressive coarsening 298 recovery of excess dislocations 300 solid solution hardening 279–81, 376–8 sub-boundary hardening 286–7 subgrain strengthening 382–3 tempered martensitic 9-12Cr steels 290–2, 301 stress relaxation testing 657–8 stress-strain behaviour 219–21, 449
WPNL2204
Index responses of materials 265–7, 370–3, 381–2 stress relief cracking 486–8 stress relief treatment 190 temperature dependence of rupture 352–5 see also strain analysis STX 21 research project 38–42 sub-boundary hardening 286–7 subgrain boundary migration 393 precipitate interaction 321–2 subgrain coarsening 382–3 subgrain strengthening 382–3 super alloy steels 169 superheaters 540–1, 554–64 surface modifications 533 surface replication 648–51 TAF steel 27 Tempaloy A-1 steel 44 temperature austenitising temperature 23 ductile-to-brittle transition 337–8, 605–10 fracture appearance transition 207 high temperature crack growth 658–62 low-temperature tempering 300 stability in testing 146 strain rate dependence 267–9 time-temperature parameter analysis 361–2 tempering 190, 290–2, 301, 381 10-12% Cr steels 32–6 tertiary creep 4, 382–3, 405 testing accelerated destructive tests 653–8 acoustic emissions (AE) monitoring 647–8 at constant load 3 at constant stress 7 carbide analysis 651–2 CEN standards 144–7 component creep tests 147 data assessment 146 deformation upon loading 269 DVM creep rate limit test 15–16, 19 hardness-based techniques 652–3
677
heat indication test (HIT) 192 isostress rupture tests 655–6 magnetic acoustic emissions 646–7 magnetic testing 645 mechanical tests 190–2 metallurgical tests 190–2 non-destructive evaluation 643–53, 663 parametric extrapolation techniques 654–5 relaxation testing 147–8 rupture tests 16–17, 145 small punch testing 656–7 steam oxidation resistance 519–20 strain monitoring 647 stress relaxation testing 657–8 stress and temperature ranges 265 surface replication 648–51 temperature stability 146 thermocouple calibrations 146 thermomechanical fatigue (TMF) testing 641 of turbine steel 190–2 ultrasonic testing (UT) 643, 645 thermal fatigue 217–19, 238 thermal nuclear reactors 598 thermal properties 225–34 conductivity 230–4 expansion 225–9 thermal stress parameter 217–18 THERMIE project 80, 208, 211 thermocouple calibrations 146 thermodynamic equilibrium analysis 315–16 thermomechanical fatigue (TMF) testing 641 thick-section components 38–42, 57–62 Thomas process 64, 66 threshold stress 78 time-dependent deformation 242–3, 350 time-temperature parameter (TTP) analysis 361–2 TMK 1 and 2 steels 31 TP 316 steel 57, 59 tracer diffusion values 257 transmission electron microscope (TEM) observations 389–93 triple junction mobility 347 tritium 625
WPNL2204
678
Index
WPNL2204
turbines 169, 573–93, 627–9 advanced turbines 174 and austenitic steels 57–62, 591 critical components 574 manufacturers specifications and standards 84 material properties 576 notch weakening 581 rotors 22, 578 welded design 590 service life of components 583–91 see also production of steels for turbines 12Cr steels 195–204 20-25% Cr steels 44, 52–3 type I cracks 484, 485 type II cracks 484, 485 type III cracks 484–5, 486–8 type IV cracks 485, 491–5, 513–17, 649 UK specifications and standards 79 ultimate tensile strength (UTS) 266 ultrasonic testing (UT) 643, 645 under-stabilising technique 47 uniaxial relaxation testing 148 user specifications 149 V-type cracks 509–11 vacancies 243–4, 260, 354 vacuum arc remelting (VAR) 66, 182 vacuum carbon deoxidation (VCD) 66, 175 vacuum degassing 66, 175 vacuum induction melting (VIM) 66, 175, 182 vacuum oxygen decarburization (VOD) 66 velocity of glide 394 void swelling 599–600 voids and gap formation 525 W-type cracks 509–11 water quenching 189–90 water spray quenching 189–90 water walls 539–40
Wedel-Neubauer recommendations 649 weld strength factor (WSF) 495–6 weld strength reduction factor (SRF) 495 welded joints 472–98 arc welding 476 CEN standards 142, 144 crack formation 484–95, 513–17 in dissimilar welds 488–90 creep behaviour 483–4 delta ferrite residuals 478 in fast breeder reactors (FBRs) 630 and ferritic steel microstructure 474–82 fusion welding 474–5 future trends 496–8 grain growth zone 479 grain refined zone 479 heat affected zone (HAZ) 472, 478–80, 497–8, 513–17 HAZ simulation 480–2 implications for industry 495–6 and life prediction procedures 662 longitudinal seam welds 472 metal development 482–3 multilayer welding 480 over-tempered region 480 partially transformed zone 479 post weld heat treatment (PWHT) 543 and precipitation strengthening 476–7 weld metal development 482–3 zone of unchanged base material 480 welding consumables 142 welding procedure qualifications 142, 144 Wiedemann-Franz Law 236 X22CrMoV steel 26–7 yield stress 265, 267–9, 270–5, 277 Young’s modulus 218, 221, 224, 238 Z-phase precipitation 68, 239, 320 Zener stress 394 zirconium 342, 345 zone of unchanged base material 480
WPNL2204
