Creep Resistant Steels

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Creep-resistant steels

i WPNL2204

ii

Related titles: High-nitrogen steels and stainless steels (ISBN 978-1-84265-129-2) This authoritative book covers manufacturing methods for and the key properties of these important steels and their alloys. Electrochemistry in light water reactors (ISBN 978-1-84569-240-7) There has long been a need for effective methods of measuring corrosion within light water nuclear reactors. This important volume discusses key issues surrounding the development of high-temperature reference electrodes and other electrochemical techniques. Corrosion by carbon and nitrogen (ISBN 978-1-84569-232-2) Metal dusting is a form of corrosion involving the disintegration of metals and alloys into a dust of graphite and metal particles when exposed to a carburising atmosphere. This important book reviews the factors affecting metal dusting and how it can be prevented in sectors such as the chemical and petrochemical industries and in the direct reduction of iron ores. It also considers the related corrosion phenomena of carburisation and nitridation. Details of these and other Woodhead Publishing materials books, as well as materials books from Maney Publishing, can be obtained by: • visiting our web site at www.woodheadpublishing.com • contacting Customer Services (e-mail: [email protected]; fax: +44 (0) 1223 893694; tel.: +44 (0) 1223 891358 ext.130; address: Woodhead Publishing Ltd, Abington Hall, Abington, Cambridge CB21 6AH, England) If you would like to receive information on forthcoming titles, please send your address details to: Francis Dodds (address, tel. and fax as above; e-mail: [email protected]). Please confirm which subject areas you are interested in. Maney currently publishes 16 peer-reviewed materials science and engineering journals. For further information visit www.maney.co.uk/journals

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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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Contents

7

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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446 447 449 449 456 459

x

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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519 520 525 530 532 533

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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637 638

xii

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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(* = 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]


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


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]


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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



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



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


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]


Chapter 14 Stuart Holdsworth EMPA – Materials Science & Technology Überlandstrasse 129 CH-8600 Dübendorf Switzerland Email: [email protected]

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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



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


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



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

1 WPNL2204

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WPNL2204

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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Introduction

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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6


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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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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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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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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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


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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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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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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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


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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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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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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Creep-resistant steels 100 000 h creep rupture strength (MPa)

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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500

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 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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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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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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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

Creep Resistant Steels

Creep Resistant Steels

Creep Resistant Steels