Power plant life management and performance improvement
© Woodhead Publishing Limited, 2011
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© Woodhead Publishing Limited, 2011
Woodhead Publishing Series in Energy: Number 23
Power plant life management and performance improvement Edited by John E. Oakey
Oxford
Cambridge
Philadelphia
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© Woodhead Publishing Limited, 2011
Published by Woodhead Publishing Limited, 80 High Street, Sawston Cambridge CB22 3HJ, UK www.woodheadpublishing.com Woodhead Publishing, 1518 Walnut Street, Suite 1100, Philadelphia, PA 19102–3406, USA Woodhead Publishing India Private Limited, G-2, Vardaan House, 7/28 Ansari Road, Daryaganj, New Delhi – 110002, India www.woodheadpublishingindia.com First published 2011, Woodhead Publishing Limited # Woodhead Publishing Limited, 2011; Chapter 14 # Alstom Power Limited, 2011 (which owns all intellectual property rights in Chapter 14, including but not limited to copyright) 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 Control Number: 2011935505 ISBN 978-1-84569-726-6 (print) ISBN 978-0-85709-380-6 (online) ISSN 2044-9364 Woodhead Publishing Series in Energy (print) ISSN 2044-9372 Woodhead Publishing Series in Energy (online) The publisher’s 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 elemental chlorine-free practices. Furthermore, the publisher ensures that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Data Standards Ltd, Frome, Somerset, UK Printed by TJI Digital, Padstow, Cornwall, UK
© Woodhead Publishing Limited, 2011
Contents
Contributor contact details Woodhead Publishing Series in Energy Foreword Part I
1
xii xv xix
Power plant fuel flexibility, condition monitoring and performance assessment
1
Solid fuel composition and power plant fuel-flexibility
3
N. J. SIMMS, Cranfield University, UK 1.1 1.2 1.3 1.4 1.5 1.6 2
Introduction Fuel chemistry and characterisation Use of alternative fuels in combustion power plants and application of technology to improve fuel flexibility Future trends Sources of further information and advice References
3 5 30 33 34 35
Condition monitoring and assessment of power plant components
38
C. DE MICHELIS, Independent Consultant, previously CESI, Italy, C. RINALDI, RSE, Italy, C. SAMPIETRI, EXOVA, Italy and R. VARIO, CESI, Italy 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8
Introduction Monitoring boiler and heat recovery steam generator Steam turbines and generators Condition monitoring of gas turbines In situ assessment of gas turbine hot parts by non-destructive techniques Remote monitoring solutions Future trends Sources of further information and advice
© Woodhead Publishing Limited, 2011
38 41 61 74 84 95 98 101
vi
Contents
2.9
References
3
Availability analysis of integrated gasification combined cycle (IGCC) power plants 110
103
A. LAUGWITZ, M. GRA¨BNER and B. MEYER, TU Bergakademie Freiberg, Germany 3.1 3.2 3.3 3.4 3.5 3.6 3.7 3.8 3.9 3.10
Introduction Basic structure of integrated gasification Availability issues of the ASU Availability issues of the gasification unit Availability issues of acid gas removal (AGR) and sulfur recovery Availability issues of the combined cycle Summary of existing plants Forecast based on RAM modeling Future trends References
110 111 116 120 125 127 129 132 135 138
Part II Coal boiler plant: materials degradation, plant life management and performance improvement
143
4
145
Environmental degradation of boiler components N. J. SIMMS, Cranfield University, UK
4.1 4.2 4.3 4.4 4.5 4.6 4.7
Introduction Component operating environments Degradation mechanisms and modelling Quantification of damage and protective measures Future trends Sources of further information and advice References
145 147 152 169 172 175 176
5
Creep in boiler materials: mechanisms, measurement and modelling
180
V. SKLENICˇKA and L. KLOC, Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Czech Republic 5.1 5.2 5.3 5.4 5.5 5.6
Introduction Creep deformation and damage mechanisms in boiler materials Measurement methods Effect of operating environment Predictive modelling Sources of further information and advice
© Woodhead Publishing Limited, 2011
180 181 191 197 211 216
Contents
vii
5.7
References
216
6
Microstructural degradation in boiler steels: materials developments, properties and assessment
222
J. DOBRZAN´SKI, Institute for Ferrous Metallurgy, Poland and A. HERNAS and G. MOSKAL, Silesian University of Technology, Poland 6.1 6.2 6.3
6.4 6.5 6.6 6.7 6.8 7
Introduction The development of steel for power engineering Methodology for assessing the state of a material and determining the residual durability of the operational elements under creep conditions Characteristics of microstructure and property degradation processes Preparation of a classification system for material after operation Modeling degradation processes and their use Conclusion References Boiler steels, damage mechanisms, inspection and life assessment
222 223
230 248 260 261 266 266 272
A. SHIBLI, European Technology Development, UK 7.1 7.2 7.3 7.4 7.5 7.6
Introduction Boiler materials, metallurgy and microstructure Damage mechanisms and component failure Inspection and monitoring of damage and integrity/life assessment issues in high chromium martensitic steels Sources of further information and advice References
Part III Gas turbine plant: materials degradation, plant life management and performance improvement 8
Creep, fatigue and microstructural degradation in gas turbine superalloys
272 274 283 291 301 302 305
307
P. AUERKARI, VTT Technical Research Centre of Finland, Finland 8.1 8.2 8.3 8.4
Introduction Creep Fatigue Combined creep and fatigue
© Woodhead Publishing Limited, 2011
307 308 314 319
viii
Contents
8.5 8.6 8.7 8.8
Microstructural degradation Future trends Conclusion References
322 327 328 328
9
Gas turbine materials selection, life management and performance improvement T. A´LVAREZ TEJEDOR, Endesa Generacio´n, Spain
330
9.1 9.2 9.3 9.4 9.5 9.6 9.7 9.8 9.9 9.10 9.11
Introduction Superalloys Protective coatings Material applications Advanced materials and coatings Life management and diagnostic Future trends Sources of further information and advice References Appendix 1: nomenclature Appendix 2: key definitions
330 332 368 377 399 403 409 411 414 417 418
10
Gas turbine maintenance, refurbishment and repair
420
A. D. WILLIAMS, Wood Group GTS, UK 10.1 10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9 10.10 10.11
Introduction Field service overhaul and maintenance Parts refurbishment: incoming inspection Parts repair Coating and finishing technology Final repair operations Quality control and first article inspection Part life extension and optimisation Future trends Conclusion Further reading
Part IV Steam boiler and turbine plant: materials degradation, plant life management and performance improvement 11
Steam oxidation in steam boiler and turbine environments
420 421 426 429 436 445 445 446 448 449 449 451
453
G. R. HOLCOMB, National Energy Technology Laboratory, U.S. Department of Energy, USA 11.1
Introduction
453
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Contents
ix
11.2 11.3 11.4 11.5 11.6 11.7 11.8 11.9 11.10
Steam boiler and turbine environments Oxidation thermodynamics and kinetics Scale morphology and spallation Steam oxidation management Future trends Conclusion Sources of further information and advice References Appendix: nominal alloy composition for alloys of interest
454 460 467 476 478 481 481 482 487
12
Steam boiler component loading, monitoring and life assessment
490
J. TALER and P. DUDA, Cracow University of Technology, Poland 12.1 12.2
12.3 12.4 12.5 13
Introduction Analysis of different ways of conducting start-up and shut-down operations and their influence on thermal and total stress loads in critical pressure components Monitoring of remnant lifetime of pressure components Conclusions References
490
491 498 514 516
Steam turbine materials selection, life management and performance improvement
518
R. W. VANSTONE and S. OSGERBY, Alstom Power, UK 13.1 13.2 13.3 13.4 13.5
Introduction High temperature cylinders Low temperature cylinders Conclusion References
518 519 524 533 533
14
Steam turbine upgrades for power plant life management and performance improvement
535
F. C. MUND, Alstom Power, UK 14.1 14.2 14.3 14.4 14.5 14.6 14.7 14.8
Introduction Drivers Product selection and specification Performance improvement Mechanical design Installation Conclusion References
© Woodhead Publishing Limited, 2011
535 537 539 547 557 563 569 571
x
Contents
14.9
Appendix: glossary
Part V Heat exchangers and power plant welds: materials management and performance improvement 15
572 573
High-temperature heat exchangers in indirectly fired combined cycle (IFCC) systems: materials management and performance improvement 575 J. P. HURLEY, University of North Dakota Energy & Environmental Research Center, USA
15.1 15.2 15.3 15.4 15.5 15.6
Introduction High-temperature heat exchanger (HTHX) construction Pilot-scale HTHX testing Conclusions Acknowledgments References
575 582 597 602 603 603
16
Heat recovery steam generators: performance management and improvement
606
V. GANAPATHY, Boiler & HRSG Consultant, India 16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8 16.9 17
Introduction Gas turbine heat recovery steam generators (HRSGs) How pinch and approach points affect HRSG size and steam generation HRSG simulation Improving HRSG efficiency Conclusion Further reading References Appendix: nomenclature Power plant welds and joints: materials management and performance improvement
606 611 618 621 628 633 633 633 633 635
D. J. ABSON, TWI, UK 17.1 17.2 17.3 17.4 17.5 17.6 17.7 17.8
Introduction Materials selection and development Weld/joint degradation Application of degradation-protection technologies Impact on power plant performance/life management Dissimilar joints Inspection and hardness testing Repair
© Woodhead Publishing Limited, 2011
635 636 639 646 647 648 651 653
Contents 17.9 17.10 17.11 17.12
xi
Future trends Sources of further information and advice Acknowledgements References
655 658 658 658
Index
666
© Woodhead Publishing Limited, 2011
Contributor contact details
(* = main contact)
Chapter 2
Editor
C. Rinaldi RSE via Rubattino 54 20134 Milan Italy E-mail:
[email protected]
J. E. Oakey Energy Technology Centre Sustainable Systems Department Cranfield University Bedfordshire MK43 0AL UK E-mail:
[email protected]
Chapters 1 and 4 N. J. Simms Centre for Energy and Resource Technology Department of Environmental Science and Technology School of Applied Sciences Cranfield University Cranfield Bedfordshire MK43 0AL UK E-mail:
[email protected]
Chapter 3 A. Laugwitz IEC – Department of Energy Process Engineering and Chemical Engineering TU Bergakademie Freiberg Fuchsmu¨hlenweg 9 09596 Freiberg Germany E-mail: Alexander.Laugwitz@iec. tu-freiberg.de
Chapter 5 V. Sklenicˇka* and L. Kloc Institute of Physics of Materials Academy of Sciences of the Czech Republic Zˇizˇkova 22 CZ-616 62 Brno Czech Republic E-mail:
[email protected];
[email protected]
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Contributor contact details
xiii
Chapter 6
Chapter 9
J. Dobrzan´ski Institute for Ferrous Metallurgy 44-100, K. Miarki str. 12-14 Gliwice Poland E-mail:
[email protected]
T. A´lvarez Tejedor Endesa Generacio´n C) Ribera del Loira 28042 Madrid Spain E-mail:
[email protected]
A. Hernas and G. Moskal* Silesian University of Technology Department of Materials Science 40-019, Krasin´skiego str. Katowice Poland E-mail:
[email protected];
[email protected]
Chapter 10
Chapter 7 A. Shibli European Technology Development Ltd 6 Axis Centre Cleeve Road Leatherhead KT22 7RD Surrey UK E-mail:
[email protected]
Chapter 8 P. Auerkari VTT Technical Research Centre of Finland POB 1000 FI-02044 VTT Finland E-mail:
[email protected]
A. D. Williams Wood Group GTS Site 4 Piper Street Baldovie Industrial Estate Dundee Scotland DD4 0NT UK E-mail: andy.williams@woodgroup. com
Chapter 11 G. R. Holcomb U.S. Department of Energy National Energy Technology Laboratory 1450 Queen Ave SW Albany OR 97321 USA E-mail:
[email protected]. gov
Chapter 12 J. Taler* and P. Duda Cracow University of Technology Al. Jana Pawøa II 37 31-864 Krako´w Poland E-mail:
[email protected];
[email protected]
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xiv
Contributor contact details
Chapter 13
Chapter 16
R. W. Vanstone and S. Osgerby* Alstom Power Newbold Road Rugby CV21 2NH UK E-mail: rod.vanstone@power. alstom.com; steve.osgerby@ power.alstom.com
V. Ganapathy Boiler & HRSG Consultant Flat B4 8-9 Nana Street T. Nagar Chennai 600017 India E-mail:
[email protected]
Chapter 17 Chapter 14 F. C. Mund Alstom Power Newbold Road Rugby CV21 2NH UK E-mail: friederike.mund@power. alstom.com
Chapter 15 J. P. Hurley University of North Dakota Energy & Environmental Research Center 15 North 23rd Street Stop 9018 Grand Forks ND 58202-9018 USA E-mail:
[email protected]
D. J. Abson TWI Granta Park Great Abington Cambridgeshire CB21 6AL UK E-mail:
[email protected]
Woodhead Publishing Series in Energy
1
Generating power at high efficiency: Combined cycle technology for sustainable energy production Eric Jeffs
2
Advanced separation techniques for nuclear fuel reprocessing and radioactive waste treatment Edited by Kenneth L. Nash and Gregg J. Lumetta
3
Bioalcohol production: Biochemical conversion of lignocellulosic biomass Edited by K.W. Waldron
4
Understanding and mitigating ageing in nuclear power plants: Materials and operational aspects of plant life management (PLiM) Edited by Philip G. Tipping
5
Advanced power plant materials, design and technology Edited by Dermot Roddy
6
Stand-alone and hybrid wind energy systems: Technology, energy storage and applications Edited by J.K. Kaldellis
7
Biodiesel science and technology: From soil to oil Jan C.J. Bart, Natale Palmeri and Stefano Cavallaro
8
Developments and innovation in carbon dioxide (CO2) capture and storage technology Volume 1: Carbon dioxide (CO2) capture, transport and industrial applications Edited by M. Mercedes Maroto-Valer
9
Geological repository systems for safe disposal of spent nuclear fuels and radioactive waste Edited by Joonhong Ahn and Michael J. Apted
10 Wind energy systems: Optimising design and construction for safe and reliable operation Edited by John D. Sørensen and Jens N. Sørensen
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Woodhead Publishing Series in Energy
11 Solid oxide fuel cell technology: Principles, performance and operations Kevin Huang and John Bannister Goodenough 12 Handbook of advanced radioactive waste conditioning technologies Edited by Michael I. Ojovan 13 Nuclear safety systems Edited by Dan Gabriel Cacuci 14 Materials for energy efficiency and thermal comfort in buildings Edited by Matthew R. Hall 15 Handbook of biofuels production: Processes and technologies Edited by Rafael Luque, Juan Campelo and James Clark 16 Developments and innovation in carbon dioxide (CO2) capture and storage technology Volume 2: Carbon dioxide (CO2) storage and utilisation Edited by M. Mercedes Maroto-Valer 17 Oxy-fuel combustion for power generation and carbon dioxide (CO2) capture Edited by Ligang Zheng 18 Small and micro combined heat and power (CHP) systems: Advanced design, performance, materials and applications Edited by Robert Beith 19 Advances in clean hydrocarbon fuel processing: Science and technology Edited by M. Rashid Khan 20 Modern gas turbine systems: High efficiency, low emission, fuel flexible power generation Edited by Peter Jansohn 21 Concentrating solar power (CSP) technology: Developments and applications Edited by Keith Lovegrove and Wes Stein 22 Nuclear corrosion science and engineering Edited by Damien Fe´ron 23 Power plant life management and performance improvement Edited by John E. Oakey 24 Direct-drive renewable energy systems Edited by Markus Mueller and Henk Polinder 25 Advanced membrane science and technology for sustainable energy and environmental applications Edited by Angelo Basile and Suzana Pereira Nunes
© Woodhead Publishing Limited, 2011
Woodhead Publishing Series in Energy
xvii
26 Irradiation embrittlement of reactor pressure vessels (RPVs) in nuclear power plants Edited by Naoki Soneda 27 High temperature superconductors (HTS) for energy applications Edited by Ziad Melhem 28 Infrastructure and methodologies for the justification of nuclear power programmes Edited by Agustı´n Alonso Santos 29 Waste to energy (WtE) conversion technology Edited by Marco Castaldi 30 Polymer electrolyte membrane and direct methanol fuel cell technology Volume 1: Fundamentals and performance Edited by Christoph Hartnig and Christina Roth 31 Polymer electrolyte membrane and direct methanol fuel cell technology Volume 2: In situ characterisation techniques Edited by Christoph Hartnig and Christina Roth 32 Combined cycle systems for near-zero emission power generation Edited by Ashok Rao 33 Modern earth buildings: Materials, engineering, construction and applications Edited by Matthew R. Hall, Rick Lindsay and Meror Krayenhoff 34 Handbook of metropolitan sustainability: Understanding and improving the urban environment Edited by Frank Zeman 35 Functional materials for energy applications Edited by John Kilner, Stephen Skinner, Stuart Irvine and Peter Edwards 36 Nuclear decommissioning: Planning, execution and international experience Edited by Michele Laraia 37 Nuclear fuel cycle science and engineering Edited by Ian Crossland 38 Electricity transmission, distribution and storage systems Edited by Ziad Melhem 39 Advances in biodiesel preparation: Second generation processes and technologies Edited by Rafael Luque and Juan Antonio Melero
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xviii
Woodhead Publishing Series in Energy
40 Biomass combustion science, technology and engineering Edited by Lasse Rosendahl 41 Ultra-supercritical coal power plant: Materials, technologies and optimisation Edited by Dongke Zhang 42 Radionuclide behaviour in the natural environment: Science, impacts and lessons for the nuclear industry Edited by Horst Geckeis and Christophe Poinssot 43 Calcium and chemical looping technology for power generation and carbon dioxide (CO2) capture: Solid oxygen- and CO2-carriers P. Fennell and E.J. Anthony 44 Materials ageing and degradation in light water reactors: Mechanisms, modelling and mitigation Edited by K.L. Murty 45 Structural alloys for power plants: Operational challenges and hightemperature materials Edited by Amir Shirzadi, Rob Wallach and Susan Jackson 46 Biolubricants: Science and technology Jan C.J. Bart, Emanuele Gucciardi and Stefano Cavallaro 47 Wind turbine blade design and materials: Improving reliability, cost and performance Edited by Povl Brøndsted and Rogier Nijssen 48 Radioactive waste management and contaminated site clean-up: Processes, technologies and international experience Edited by William E. Lee, Michael I. Ojovan and Carol M. Jantzen 49 Probabilistic methods of strength reliability and their application for optimum nuclear power plant life management (PLiM) Gennadij V. Arkadov, Alexander F. Getman and Anderi N. Rodionov 50 Coal utilization in industry: Towards cleaner production Edited by D.G. Osborne 51 Coal power plant materials and life assessment: Developments and applications Edited by Ahmed Shibli
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Foreword
A sustainable energy future will depend on making appropriate use of all possible resources while ensuring a stable and affordable supply that complies with current and future regulations for protecting the environment. Our future energy portfolio will aim to maximize the use of so-called renewable energy sources such as wind, solar (photovoltaic and solarthermal), geothermal, hydro, wave-power, and biomass, but it must involve realistic expectations for these sources. Further, transitioning to the maximized use of renewable energy must be coordinated with an orderly reduction of reliance on traditional energy sources currently used for baseline production of electric power, chiefly coal, natural gas, and nuclear fission. Some scenarios envisage continued reliance on nuclear power to maintain a baseline capacity. However, the ability to adjust the output of nuclear plants to cope with a need for large amounts of generating capacity at short notice, when renewable sources are unable to generate, for instance, is rather limited. Further, the size of this required ‘reserve’ capacity is considerable, as recent experience in the UK during the exceptionally cold and becalmed month of December 2010 attests, when many old and semiretired fossil power plants had to be brought on line to deputize for the absence of wind power. The only viable options for this reserve capacity at present are natural gas-fired gas turbines and coal-fired steam boilers, so that for the foreseeable future it appears that some degree of the overall generating capacity must continue to involve the traditional, non-renewable sources. In such duty, fossil-fired plants not only will be required to operate at maximum efficiency to be compatible with reduced emission of environmental pollutants (including carbon dioxide), but also must be capable of rapid load cycling. Maximizing efficiency of fossil-fueled systems based on the Brayton and Rankin cycles typically involves operation at the highest possible temperatures, pushing critical structural components to operate at their maximum temperature limits. In addition, normal commercial operation of these plants increasingly will require them to accept a range of fuels, such as
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xx
Foreword
gaseous fuels with varying H/C ratios or derived from coal or biomass, and coal from different sources worldwide. The imposition of severe thermal cycling on the requirements for extreme high-temperature strength and improved environmental resistance further complicate the task of maintaining the desired reliable, long-term operating characteristics of these plants. Successful developments will require a detailed knowledge of how such changes in the operating environment influence the degradation and failure modes of key structural components, and the development of mechanistically-based predictive methods to allow the most appropriate scheduling of maintenance and avoidance of unexpected failures. The content of this volume provides excellent insight into the issues that need to be addressed if fossil-fueled power plants are to attain the desired higher levels of operating efficiency while maintaining the necessary levels of reliability. In particular, discussion of approaches being developed for life prediction of key components and the research to provide the underpinning base of understanding emphasize the importance of the linkages among these activities, and their importance to the reliable operation of these power systems. These issues also emphasize that once the mix of generating sources in our future energy portfolio is decided, appropriate investment will be needed to ensure that all approaches, renewable and traditional, can perform as intended and are ready in the desired timeframe. This goal requires not only establishing the infrastructure and maintenance procedures for the new, renewable, systems but also committing sufficient R&D resources to ensure that fossil- and biomass-based generation are as reliable and efficient as possible. Ian G. Wright Denver, Colorado, USA
© Woodhead Publishing Limited, 2011
1 Solid fuel composition and power plant fuel flexibility N . J . S I M M S , Cranfield University, UK
Abstract: A fundamental part of any power plant system is the fuel that it uses. For solid-fuel-fired systems, the composition of the fuels affects every aspect of the plants, from fuel handling and storage, through the fuel conversion process (e.g. combustion or gasification) and heat recovery/energy generation to the gas cleaning systems and environmental emissions (by gas, liquid or solid streams). Traditionally, coals have been used in large scale power plants, but now biomass and waste fuels are being introduced to reduce net CO2 emissions and to increase the consumption of more sustainable fuels. This chapter outlines the main types of solid fuels, analysis methods and their compositions, before highlighting the composition related issues that need to be considered in assessing their role in fuel switching and co-firing. Key words: coal, biomass, waste, fuel composition, fuel analyses, cofiring.
1.1
Introduction
One of the fundamental parts of any power generation system is the fuel that it uses. Many factors depend directly on the fuel used, for example: . . .
the type of power generation system, e.g. pulverised fuel, fluidised bed, moving grate; the scale of the system; for coal-fired boilers typically in the range 500– 800 MWe and for biomass-fired boilers usually less than 30 MWe; system efficiency; with coal-fired power stations having much higher
3 © Woodhead Publishing Limited, 2011
4
Power plant life management and performance improvement
. .
efficiencies (typically ~35–47 %) than waste-fired systems (up to ~25 % efficiency); gas cleaning requirements and emissions; ash and other waste disposal, including solid and liquid products from the gas cleaning systems; economic viability.
.
Many of these topics are dealt with in detail in other chapters of this book, with this chapter focusing on solid fuel compositions. Potential solid fuels for generating power include: .
.
.
Coals. Traditionally large scale solid-fuel-fired power generation systems have been based on the use of pulverised coal, with the coal being sourced from local mines to minimise transport costs. However, variations in mining costs around the world, combined with national energy policies and national/international environmental regulations, have prompted an increase in world-traded coal. For example, UK government statistics show that in 2008, 48.3 m tonnes of coal were used for power generation (out of a total coal consumption of 58.2 million tonnes), with ~70 % of the coal used in the UK being imported (UK DECC, 2009). Biomass. Biomass, such as wood or crop residues, has traditionally been used to generate heat on a small scale but now biomass is being used in increasing quantities for the generation of both heat and electrical power. The growth of specific crops for use in energy generation systems is being actively encouraged and biomass fuel supply chains developed (EUBIONET2, 2007). A notable development in the use of biomass in recent years has been co-firing it with coal in traditionally pulverised coal-fired power stations; this has been actively encouraged as one route of minimising net CO2 emissions (UK DTI, 2007). Waste products. Within Europe, one route for the disposal of solid wastes has been via combustion processes to generate heat (e.g. in cement kilns) or to generate heat and/or power in waste to energy plants. Changes in the approach to dealing with waste, with increased emphasis on recycling and a reduction in the quantities going to landfill (Council of the European Union, 1999), have encouraged the consideration of wastes as potential fuels for power generation. In addition, significant fractions of some waste streams can be regarded as being biomass derived and so viewed as sustainable and renewable fuels (IEA Bioenergy, 2003).
© Woodhead Publishing Limited, 2011
Solid fuel composition and power plant fuel flexibility
5
In the drive towards reducing the impact of coal used for power generation in terms of CO2 emissions, approaches include: . . .
.
the use of more sustainable fuels, for example: ○ replacing coal with another fuel (biomass or waste products), and ○ co-firing coal with biomass or waste products; using natural gas as an alternative fuel; increasing the efficiency of electricity generation from coal, e.g. by building new more efficient advanced power stations, or using alternative generation processes such as fuel gasification, or by using the waste heat generated during these processes; using carbon capture and storage systems on existing and new solidfuel-fired power systems (post-combustion, pre-combustion and oxyfiring options).
This chapter focuses on the solid fuels (coals, biomass and wastes) that are currently available, the compositions of these fuels, methods of assessing them and issues that need to be considered for fuel flexibility.
1.2
Fuel chemistry and characterisation
1.2.1 Fuel specifications and analysis methods The analysis of fuels has developed in parallel to their use, and numerous standards have been developed for specific fuels in different countries and by different standards organisations (e.g. ASTM International, British Standards Institute (BSI), European Committee for Standardisation (CEN), International Organisation for Standardisation (ISO)). These various analytical standards have been developed to cover the requirements of different fuel users (e.g. coke production and iron ore smelting, as well as for power generation) in addition to the needs of regulatory bodies for use in applying environmental regulations, determining fuel subsidies, etc. From the point of view of using a fuel for power generation purposes, the key types of analyses can be grouped into those related to: . . .
chemical analyses, to show the chemical make-up of the fuel in different ways; physical analyses, to determine the properties that influence fuel preparation; other general properties needed for fuels, such as energy content, ash fusion temperatures, etc.
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Power plant life management and performance improvement
Chemical analyses During the last century various types of analyses have been developed to show the chemical make-up of coals, with many being aligned to meet the needs of specific industrial uses (Francis and Peters, 1980; Raask, 1985; Speight, 1994, 2005). Some of these analytical methods are now used to characterise coals for use in power generation systems and variations on them have been developed to enable biomass and waste fuels to be analysed to produce equivalent data (sections 1.2.2 to 1.2.4). .
Proximate analysis. This produces an analysis of the fuels in terms of four specific parameters (BS 1016, 2010): ○ ‘Moisture’ is a measure of the amount of water present in the fuel. It is measured by drying a fuel under standard conditions. Due to the different ways in which water can be present in a fuel (e.g. on the fuel surface, hydroscopically, associated with organic and inorganic compounds and/or in cell structures) a number of different methods have been developed to determine it, but they all measure the change in weight of the fuel as a result of a drying process. It is important to note that fuel moisture content can change between sampling and analysis, as well as during storage. Fuel moisture content is expressed as a percentage of the weight of water compared to the original fuel (i.e. on an as-received or wet basis). ○ ‘Ash’ is a measure of the non-combustible residue left after a fuel is burnt. This property is determined by thoroughly burning the fuel and then weighing the residue and expressing this as a percentage of the original as-received (i.e. wet) fuel, or the fuel after drying. These two measures of the ash content are related as follows: ashwt%dry ¼ ashwt%ar 6100=ð100 water contentðwt%ÞÞ
○
○
where ashwt% dry is the weight % of ash in the dry fuel and ashwt%ar is the weight % of ash in the as-received fuel. ‘Volatile matter’ is a measure of the components within a fuel (but excluding moisture) that are vapourised in the absence of air at high temperature. These components typically include a mixture of hydrocarbons (short and long chained, and aromatic compounds) as well as some sulphur and chlorine compounds. Due to the variable nature of the volatile components in fuels, this fuel parameter has to be determined under strictly controlled standard conditions. Volatile matter can be expressed in terms of the % by weight in a dry fuel, as-received fuel, dry and ash free (daf) fuel, or dry and mineral matter free (dmmf) fuel. ‘Fixed carbon’ is the carbon content of the fuel left after the volatile
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Solid fuel composition and power plant fuel flexibility Table 1.1 basis
7
Relationship between proximate analysis parameters and analysis
Analysis basis
Method of calculation of fixed carbon in fuel
Dry Dry ash free (daf) As-received (ar)
Fixed C = 100 ash (dry) volatiles (dry) Fixed C = 100 volatiles (daf) Fixed C = 100 ash (ar) water content volatiles (ar)
Table 1.2 basis
Relationship between ultimate analyses parameters and analysis
Analysis basis
Definition
Dry Dry ash free (daf) As-received (ar)
C + H + O + N + S + Cl + ash = 100 C + H + O + N + S + Cl = 100 C + H + O + N + S + Cl + ash + water content = 100
.
.
.
matter is driven off. (It does not correspond to the carbon content of the fuel as carbon can be a significant part of the volatile matter content of the fuel.) This parameter was originally intended to give an indication of the suitability of a coal for coking purposes. It can be calculated in terms of % by weight on the basis of dry, dry ash free or as-received analysis, as indicated in Table 1.1. Ultimate analysis. This group of analyses produces measurements of the quantities of specific elements in fuels. The content of carbon (C), hydrogen (H), oxygen (O), nitrogen (N), sulphur (S) and chlorine (Cl) in fuels is commonly measured, with each being determined using different specific standard analysis techniques. The results of these analyses are expressed in terms of the % by weight of the element in the fuel on the basis of dry, dry ash free or as-received analysis, as indicated in Table 1.2. The oxygen content is often not measured but calculated as the difference between 100 and the % values of the measured components. When the oxygen content is measured the total sum may not equal 100 due to experimental errors in the various analyses. Ash analysis. Following the convention established in coal analysis standards (Stringer, 1995; ASTM-D3174-11, 2011), the results of ash analyses are often expressed in terms of the % by weight of the elements in the dry ash in terms of their highest oxides. These oxides are not representative of the actual chemical compounds in the ashes that contain these elements (section 1.2.2), and so care is required in interpreting such data (Stringer, 1995). Minor and trace metal analysis. The content of the metal elements present in fuels can also be determined by a wide range of analytical techniques, including emission spectroscopy, flame photometry, ICP-
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Power plant life management and performance improvement
.
MS, etc., according to specific standards for the different fuels (sections 1.2.2 to 1.2.4) (Francis and Peters, 1980; Raask, 1985; Clarke and Sloss, 1992; Carpenter and Skorupska, 1993; Speight, 1994, 2005; van Loo and Koppeian, 2007). In these cases, the results are usually expressed in terms of the mg of element per kg of dry fuel (i.e. ppm). Biochemical composition. For some fuels it is appropriate to determine their biochemical composition (van Loo and Koppeian, 2007). This is usually expressed in terms of the % by weight in the dry material: e.g. for cellulose, hemi-cellulose, lignin, fats, protein, pectin, starch, extractives, C5 and C6 sugars, total non-structural matter, carbohydrates, etc.
Physical analyses . .
.
.
‘Bulk density’ of a fuel, in terms of kg/m3; this can vary widely between different types of fuel. ‘Grindability’ is an important parameter for coals that need to be processed by grinding before being used in a pulverised-fuel fired power station. One empirical method of measuring this is the Hardgrove grindability index (BS 1016 – Part 20, 1981). ‘Abrasion’ is a process by which one material is worn away by another material passing over it. This is important for fuel handling systems in which fuels may contain a wide variety of constituents of varying hardnesses, some of which may cause damage to the handling systems. For coals this includes abrasion by quartz and pyrites. One simple test for abrasion involves putting the fuel into a rapidly rotating mill (1500 revolutions per minute) with four steel blades of known mass and then measuring the mass loss of the blades after 8 minutes of operation (Foster et al., 2004); the data generated give an abrasion index. Particle size distributions: these measurements are needed to ensure that the appropriate fuel preparation and handling equipment is used for the different fuels; e.g. to assess the performance of coal grinding mills on a specific coal, or coal–biomass mixture, or the effectiveness of biomass hammer mills in the preparation of biomass. Oversized fuel particles do not burn out completely and so reduce the efficiency of a combustion system, as well as increase the carbon content of ash residues. (For coalfired systems, this leads on to problems with the sale and use of pulverised fuel ash residues as cement replacement materials or in lightweight aggregates.) Particle sizes can be measured by standardised sieving, photometry and optical techniques.
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Solid fuel composition and power plant fuel flexibility
9
Other general fuel properties .
Energy content. A critical property for any potential fuel is its energy content; this is often referred to as the calorific value of a fuel. Analytically, the energy content of a material can be measured using a bomb calorimeter and specific standard methods (sections 1.2.2 to 1.2.4). The values determined are usually the higher heating values, but following established conventions these results can be expressed in terms of either higher or lower heating values, with the difference being the heat of evaporation of the water formed from the hydrogen in the fuel as well as the moisture present in the fuel. Higher heating values (HHVs) are also referred to as gross heating values, calorific values and heats of combustion. Lower heating values (LHVs) are also referred to as net heating values (Phyllis, 2010).
As a large quantity of data on fuel energy contents and compositions has been generated over a long period of time, it has been found that it is possible to calculate HHVs from the fuel elemental composition. One example of this is the Milne formula (Phyllis, 2010): HHVMilne ¼ 0:341 C þ 1:322 H 0:12 O 0:12 N þ 0:0686S 0:0153 ash where C, H, etc., are the mass and the ash fractions in wt% of dry material and HHV the heating value in MJ/kg. By using the hydrogen and ash fractions (wt% dry) and moisture fraction w (wt% ar) the different HHVs and LHVs can be calculated, for example: HHVar ¼ HHVdry ð1 w=100Þ HHVdry ¼ HHVdaf ð1 ash=100Þ LHVar ¼ HHVar 2:442fð8:936H=100Þð1 w=100Þ þ w=100g .
Ash fusion. One of the frequently encountered challenges with fuels is that, under some operating conditions, the combustion residue (ash) can form a hard glassy slag on some heat exchanger surfaces. Most fuel conversion systems perform better with the ash in a powdery form (though systems do exist that handle ashes in molten form). Ash fusion tests are carried out by viewing moulded samples of fuel ashes (in the form of cones, pyramids or cubes) through windows in high temperature furnaces. A series of ash fusion temperatures have been defined (Raask, 1985): ○ initial deformation temperature, when the corners of the mould first become rounded;
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Power plant life management and performance improvement ○ ○ ○
softening (sphere) temperature, when the top of the mould takes on a spherical shape; hemisphere temperature, when the entire mould takes on a hemisphere shape; flow (fluid) temperature, when the molten ash collapses to a flattened pool.
The ash fusion temperatures can be measured in both oxidising and reducing conditions, and are highly dependent on the detailed test methods used. Therefore, standards exist for specific fuels to enable comparable results to be obtained (sections 1.2.2 to 1.2.4). .
Coal petrology. This is a method of characterising coal by examining polished cross-sections by optical microscopy and determining the reflectance of three phases (liptinite, vitrinite and inertinite) within the coal structures. The vitrinite phase in particular can be used in the classification of coals (Speight, 2005).
1.2.2 Coal fuels Coals are fossil fuels derived from plant matter that has been saved by water and/or mud from oxidation and biodegradation and then subjected to high pressures and temperatures for prolonged periods; this process is described in detail elsewhere (e.g. Raask, 1985; Speight, 1994). Thus coals can be classed as sedimentary organic rocks. There are many variables that influence this process, including initial plant matter, pressure history, temperature history and time. As a result of differences in such variables during their formation, a wide range of coals have developed, but these can
1.1 Breakdown of coal constituents (adapted from Jones, 2005).
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Solid fuel composition and power plant fuel flexibility
11
be thought of as being built up from a number of different types of materials. As illustrated in Fig. 1.1, the primary split is between coal matter (C, H, O, N, etc.) and inert material (moisture and various minerals). The inert material can be further broken down into its location within the coal structure. It should be noted that several elements can occur by more than one route in this view of the make-up of coal. As coal is a well established fuel around the world, numerous analytical methods have been developed and adopted by the various standard bodies (e.g. ASTM International, 2010; BSI, 2010; European Committee for Standardisation (CEN), 2010; ISO, 2010). Table 1.3 lists some of the coal related standards available from ASTM (2010), but similar standards exist from the other standard bodies (e.g. the multi-part BS 1016 from BSI (2010) contains another selection of standards relating to coal analysis). During the study of peats and coals, a number of different methods of classifying them have been developed. One commonly used method divides peats and coals up by into five broad types: . . . . .
‘peat’, material at an early stage in coal formation; ‘lignite’ (or ‘brown coal’), with a high moisture content; ‘sub-bituminous coal’; ‘bituminous coal’, a dense, usually black coal, frequently with a banded structure; ‘anthracite’, a glossy, hard, black coal with a high carbon content and low in volatile matter.
Figure 1.2 provides an illustration of this, together with the relative abundances of these types of coals and their common usages. In addition, Fig. 1.2 illustrates that there are clear trends in composition in passing along this series of fuels; the moisture content decreases, but the carbon and energy contents both increase. Table 1.4 gives a more detailed breakdown of this approach to coal classification and illustrates some of the differences between the standards used in different countries. In this case the trend from peat towards anthracite is characterised by the progressively lower water and volatile matter and higher energy contents. An alternative coal classification system is shown in Table 1.5. This coal ranking system defines a coal class based on the volatile matter content of a coal and its dilation on heating. Other methods can use coal petrology to determine the vitrinite reflection and so assist in the ranking of coals (Table 1.4).
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Table 1.3
Examples of coal analysis standards
Designation
Title
D388-05 D1412-07
Standard Classification of Coals by Rank Standard Test Method for Equilibrium Moisture of Coal at 96 to 97 Percent Relative Humidity and 30 8C Standard Test Method for Determination as Carbon Dioxide of Carbonate Carbon in Coal Standard Test Method for Fusibility of Coal and Coke Ash Standard Test Method for Forms of Sulfur in Coal
D1756-02(2007) D1857-04 D2492-02(2007) D2961-02(2007)
Standard Test Method for Single-Stage Total Moisture Less than 15 % in Coal Reduced to 2.36-mm (No. 8 Sieve) Topsize
D3172-07a
Standard Practice for Proximate Analysis of Coal and Coke
D3173-03(2008)
Standard Test Method for Moisture in the Analysis Sample of Coal and Coke
D3174-04
Standard Test Method for Ash in the Analysis Sample of Coal and Coke from Coal
D3175-07
Standard Test Method for Volatile Matter in the Analysis Sample of Coal and Coke
D3176-09
Standard Practice for Ultimate Analysis of Coal and Coke
D3177-02(2007)
Standard Test Methods for Total Sulfur in the Analysis Sample of Coal and Coke
D3180-07
Standard Practice for Calculating Coal and Coke Analyses from As-Determined to Different Bases
D3302/D3302M-09 Standard Test Method for Total Moisture in Coal D4208-02(2007)
Standard Test Method for Total Chlorine in Coal by the Oxygen Bomb Combustion/Ion Selective Electrode Method
D4239-08
Standard Test Methods for Sulfur in the Analysis Sample of Coal and Coke Using High-Temperature Tube Furnace Combustion Methods
D5142-09
Standard Test Methods for Proximate Analysis of the Analysis Sample of Coal and Coke by Instrumental Procedures
D5373-08
Standard Test Methods for Instrumental Determination of Carbon, Hydrogen, and Nitrogen in Laboratory Samples of Coal
D5865-07a
Standard Test Method for Gross Calorific Value of Coal and Coke
D7582-09
Standard Test Methods for Proximate Analysis of Coal and Coke by Macro Thermogravimetric Analysis
D3682-01(2006)
Standard Test Method for Major and Minor Elements in Combustion Residues from Coal Utilization Processes
D3683–04
Standard Test Method for Trace Elements in Coal and Coke Ash by Atomic Absorption
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Solid fuel composition and power plant fuel flexibility
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Table 1.3 (cont.)
Designation
Title
D3684-01(2006)
Standard Test Method for Total Mercury in Coal by the Oxygen Bomb Combustion/Atomic Absorption Method
D3761-96(2002)
Standard Test Method for Total Fluorine in Coal by the Oxygen Bomb Combustion/Ion Selective Electrode Method
D4326–04
Standard Test Method for Major and Minor Elements in Coal and Coke Ash by X-Ray Fluorescence
D4606-03(2007)
Standard Test Method for Determination of Arsenic and Selenium in Coal by the Hydride Generation/Atomic Absorption Method
D5016-08e1
Standard Test Method for Total Sulfur in Coal and Coke Combustion Residues Using a High-Temperature Tube Furnace Combustion Method with Infrared Absorption
D5987-96(2007)
Standard Test Method for Total Fluorine in Coal and Coke by Pyrohydrolytic Extraction and Ion Selective Electrode or Ion Chromatograph Methods
D6316-09b
Standard Test Method for Determination of Total, Combustible and Carbonate Carbon in Solid Residues from Coal and Coke
D6349–09
Standard Test Method for Determination of Major and Minor Elements in Coal, Coke, and Solid Residues from Combustion of Coal and Coke by Inductively Coupled Plasma – Atomic Emission Spectrometry
D6357-04
Test Methods for Determination of Trace Elements in Coal, Coke, and Combustion Residues from Coal Utilization Processes by Inductively Coupled Plasma Atomic Emission, Inductively Coupled Plasma Mass, and Graphite Furnace Atomic Absorption Spectrometry
D6414-01(2006)
Standard Test Methods for Total Mercury in Coal and Coal Combustion Residues by Acid Extraction or Wet Oxidation/ Cold Vapor Atomic Absorption
D6721-01(2006)
Standard Test Method for Determination of Chlorine in Coal by Oxidative Hydrolysis Microcoulometry
D6722-01(2006)
Standard Test Method for Total Mercury in Coal and Coal Combustion Residues by Direct Combustion Analysis
D7348-08
Standard Test Methods for Loss on Ignition (LOI) of Solid Combustion Residues
Source: compiled from ASTM (2010).
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1.2
Relationship between coal types and coal carbon and moisture contents (adapted from World Coal, 2010).
14 Power plant life management and performance improvement
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Solid fuel composition and power plant fuel flexibility
15
Table 1.4 Coal classification UN-ECE
Peat Ortho-lignite
Coal types and peat USA Germany (DIN) (ASTM) Peat
Total water content (%)
Torf
Energy content Volatiles Vitrinite af * maf ** reflection in oil (%) (MJ/kg)
75
6.7
35
16.5
0.3
25
19
0.45
10
25
Weichbraunkohle
Lignite
Meta-lignite
Mattbraunkohle
SubSub-bituminous bituminous Glanzbraunkohle coal coal
45
0.65
40
0.75
35
1.0
28
1.2
19
1.6
14
1.9
10
2.2
Bituminous coal
High volatile Gasflammkohle bituminous coal Gaskohle Med. vol. bitum. coal Fettkohle
Steinkohle Hartkohle
Flammkohle
36
Low volatile bitum. coal Esskohle
Anthracite
Semianthracite
Magerkohle
Anthracite
Anthrazit
3
36
*
Ash-free. Moisture and ash-free. Source: adapted from Euracoal (2010). **
From the point of view of using coals in power generating systems, the properties of the coals progressively change through the classification systems, with the anthracite coals having the highest calorific values and the lowest H/C ratios, in contrast to the lignites having the lowest calorific values and the highest H/C ratios. One way of showing this progression in terms of fuel composition is using a van Krevelen diagram in which the H/C ratios of fuels are plotted as a function of their O/C ratios (van Krevelen, 1950). Figure 1.3 includes the various coal types in such a diagram and also shows the position of biomass and peats (the data points on this diagram represent coals and biomass used in power generating systems). This type of diagram is useful in showing some of the trends in bulk composition from biomass through peat and the different coal types through to anthracite. More detailed coal compositions are produced using the analytical methods outlined in section 1.2.1 above. Examples of coal analyses are given in Table 1.6 to indicate typical values for three power station coals sourced
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Table 1.5 Class
Coal rank
Volatile matter* General description (weight %)
101 102 201 202 203 204 206
< 6.1 3.1–9.0 9.1–13.5 13.6–15.0 15.1–17.0 17.1–19.5 19.1–19.5
301 305 306 401 402 501 502 601 602 701 702 801 802 901 902
19.6–32.0 19.6–32.0 19.6–32.0 32.1–36.0 > 36.0 32.1–36.0 > 36.0 32.1–36.0 > 36.0 32.1 > 36.0 32.1–36.0 > 36.0 32.1–36.0 > 36.0
Anthracites Dry steam coals Coking steam coals Heat altered low volatile steam coals Prime cooking coals
Low volatile steam coals
Medium volatile coals
Mainly heat altered coals Very strongly coking coals Strongly coking coals Medium coking coals High volatile coals Weakly coking coals Very weakly coking coals Non-coking coals
*
Volatile matter–dry mineral matter free basis. In coal, those products, exclusive of moisture, are given off as gas and vapour determined analytically.
1.3 Relationship between H/C and O/C ratios for coals and biomass.
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Solid fuel composition and power plant fuel flexibility Table 1.6
Typical coal analyses
Parameter
Unit
Moisture
wt% ar
UK (Thorseby) 4.8
South American (El Cerrejon) 7.0
South African (Koornfontein) 3.8
Ash
wt% dry
11.8
9.0
13.9
CV (gross)
MJ/kg daf
34.1
33.1
34.1
32.9
31.8
32.9
wt% daf
84.3
79.9
84.5
CV (net) C
17
H
4.6
5.3
5.2
O
7.9
12.21
8.8
N
1.8
1.7
2.1
S
2.13
0.73
0.6
Cl
0.67
0.03
0.1
Ash analysis (% on ash) SiO2
54.4
58.72
43.7
Al2O3
24.5
21.30
34.0
Fe2O3
10.7
7.19
3.0 7.2
CaO
2.36
2.20
MgO
1.62
2.81
2.2
K2O
3.13
2.24
<0.5
Na2O
1.88
1.03
0.4
TiO2
1.07
0.89
1.7
BaO
0.11
0.11
—
Mn3O4
0.05
0.06
—
P2O5
0.15
0.21
1.0
SO3
3.65
3.92
6.3
in the UK, South America and South Africa. It should be emphasised that a wide range of compositions can be found for coals mined in diverse locations around the world, as a result of differences in their formation and the local geology. There are less significant variations between coals mined in smaller geographic areas (though they are still influenced by the local geology), but there are even some differences between coals produced from different seams in the same coal mine. While the form and quantities of the major elements present in coal are critical in combustion processes (as well as gasification and pyrolysis processes), it is the minor and trace elements that cause many of the operational problems for practical power generating systems (Raask, 1985), e.g. emissions/gas cleaning system requirements for SOX, NOX and other
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Power plant life management and performance improvement
1.4
Sulphur and chlorine contents of selected biomass and coals.
species; fouling and slagging on the various heat exchanger surfaces; fireside and dewpoint corrosion. Figure 1.4 illustrates the variation in S and Cl contents for a wide range of example coals (and biomass). For coals in general, S contents can range from approximately zero up to ~4 weight % (dry ash free basis), while the Cl content can range from approximately zero up to ~0.7 weight % (dry ash free basis). Higher S content coals require more gas cleaning to remove and control SOX emission. Higher S and Cl levels both play a role in the formation of slagging and fouling deposits, as well as the various corrosion processes (Raask, 1985). The ash in the fuel analysis mostly arises from the mineral impurities that are found in coals (Fig. 1.1) (Francis and Peters, 1980; Raask, 1985; Speight, 1994). However, the results of standard ash analyses on coals do not give a good representation of the form of these elements (Stringer, 1995). Table 1.7 lists the main mineral types found in coals and provides examples of minerals that are frequently found in coals (Stringer, 1995). It is the decomposition and interaction reactions of these minerals during combustion (or gasification/pyroloysis) that produces most of the ash (bottom ash and fly ash), as well as the slagging and fouling deposits (Raask, 1985).
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Solid fuel composition and power plant fuel flexibility Table 1.7
19
Common minerals found in coal
Mineral group
Mineral name
Chemical formula
Clay minerals
Montmorillonite
Al2Si4O10(OH)2.H2O
Illite
KAl2(AlSi3O10)(OH)2
Kaolinite
Al4Si4O10(OH)8
Sulphide minerals
Pyrite
FeS2
Marcasite Sulphate minerals
Gypsum
FeS2 CaSO4.2H2O
Carbonate minerals
Chloride minerals Silicate minerals
Oxide minerals
Anhydrite
CaSO4
Jarosite
(Na,K)Fe3(SO4)2(OH)6
Calcite
CaCO3
Dolomite
(Ca,Mg)CO3
Siderite
FeCO3
Ankerite
(Ca,Fe,Mg)CO3
Halite
NaCl
Sylvite
KCl
Quartz
SiO2
Albite
NaAlSi3O8
Orthoclase
KAlSi3O8
Fayalite
Fe2SiO4
Haematite
Fe2O3
Magnetite
Fe3O4
Rutile
TiO2
Source: Stringer (1995).
Coals also contain many trace elements, i.e. elements present at levels below 1000 ppm. In fact most naturally occurring elements can be found in different coals (Clarke and Sloss, 1992). Table 1.8 lists trace elements typically found in coals at levels of > 10 ppm; more comprehensive lists are available elsewhere (e.g. Clarke and Sloss, 1992). In practice these trace elements can behave in a wide range of different ways during combustion processes (and differently during gasification and pyrolysis), with the result that they can be distributed (or partitioned) between bottom ash, fly ash, various deposits and the gas phase. For simplicity, this behaviour of elements can be grouped so that they are classed as having different volatilities during combustion and in combusted fuel gas streams (Fig. 1.5). Hg species are notable for being particularly volatile and are of concern in terms of atmospheric emissions, even though Hg is generally present at levels of <1 ppm in coals (Clarke and Sloss, 1992).
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Table 1.8
Trace elements in coals – with typical levels of > 10 ppm Coal
Element Arsenic Barium Boron Bromine Cerium Chromium Copper Fluorine Lanthanum Lead Lithium Manganese Neodymium Nickel Phosphorous Rubidium Strontium Titanium Vanadium Yttrium Zinc Zirconium
Typical (ppm)
Range (ppm)
10 200 50 20 20 20 15 150 10 40 20 70 10 20 150 15 200 600 40 15 50 50
0.5–80 20–1000 5–400 0.5–90 2–70 0.5–60 0.5–50 20–500 1–40 2–80 1–80 5–300 3–30 0.5–50 10–3000 2–50 15–500 10–2000 2–100 2–50 5–300 5–200
Note. Emissions from some other race elements are also of concern, but their occurrence is typically at levels of <10 ppm in the fuel (e.g. mercury with a typical level of 0.1 ppm and range of 0.02–1 ppm). Source: adapted from Clarke and Sloss (1992).
Detailed reports are available from the International Energy Agency Clean Coal Centre on many specific aspects of coal compositions and their impact on coal usage; e.g. general impurity removal (Couch, 1995), trace metals (Clarke and Sloss, 1992), halogens (Sloss, 1992), and S and N species (Nalbandian, 2004). Finally it should be noted that, in practice, coals are often only partially analysed on a routine basis to determine the parameters that are needed for pricing, quality control, plant operation, emissions and regulatory purposes (e.g. CV, moisture, ash, sulphur content).
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1.5 Classification of volatility of trace elements in combustion systems. Reproduced with permission from IEA CCC from Clarke L B and Sloss L L (1992), Trace Elements – Emissions from Coal Combustion and Gasification, IEA Coal Research, IEACR/49.
1.2.3 Biomass fuels There is a wide range of types of biomass that could potentially be used in power generation systems. However, the growth of biomass depends on local environments and soil conditions, so that the types that are available in practical terms vary between geographic regions. Biomass can be classed in different ways (White and Plaskett, 1981; Simms et al., 2007a; van Loo and Koppeian, 2007; Livingston, 2009), but one method uses the biomass production route as the basis: . . . . .
Energy crops: ○ woods, e.g. coppiced willow, poplar, cottonwood; ○ grasses, e.g. miscanthus, reed canary grass, switch grass. Agricultural and forestry residues: ○ straws, e.g. from wheat, barley, oats, rice, maize, oil seed rape; ○ forest residues. Processing residues from: ○ olives, almonds, palm nuts, sugarcane, rice; ○ sawdust, bark, wood off-cuts. Seaweeds, both naturally occurring and cultivated. Animal wastes and sewage sludges. Some types of biomass can be supplied in a processed form as pellets, such
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Power plant life management and performance improvement
as cereal co-products (CCP) or wood by-products, which increases their energy density and makes them easier to handle, transport and store. As a result of the low energy density of many types of biomass, many are only available to local markets, but processed biomass and those with higher energy densities (e.g. some forms of wood and pelletised products) are available on a world traded basis (Colechin, 2005; EUBIONET2, 2007; Simms et al., 2007a). The compositions of biomass can be determined and presented in several distinct ways depending on the intended use of the biomass and the data. For example, different types of biomass can be used as fuels for thermal processes, such as combustion, gasification or pyrolysis, or anaerobic digestion systems, and these processes require different types of fuel data. In addition, other data are also required for the optimisation of the production of the biomass and the development of improved plant varieties. Thus, potential biomass can be analysed in terms of their contents of lignin, hemicellulose, cellulose, starch, fats, proteins, etc., for some purposes (e.g. White and Plaskett, 1981; van Loos and Koppeian, 2007). However, for thermally based power generation applications, and comparisons with other solid fuels, it is more useful to produce analyses in terms of the traditional fuel parameters (section 1.2.1). The differences in the structures of biomass and coals means that in order to generate these fuel parameters for biomass, it has been necessary to develop a new set of standard analytical methods. As for coal analyses, alternatives are available from different standards bodies (e.g. ASTM International, 2010; BSI, 2010; CEN, 2010; ISO, 2010); these biomass standards are all relatively new, having been developed during the last decade in response to the desire to use biomass much more widely as a fuel for power generation. Table 1.9 gives examples of the main biomass analysis standards available from the European Committee for Standardisation (CEN, 2010). Using such standards, a wide range of potential biomass fuels have been analysed to produce consistent and comparable fuel composition data; these are now becoming available through many publications and a wide range of databases that are accessible through the internet (section 1.5). Table 1.10 gives examples of fuel analyses of selected biomasses. Fuel moisture has been omitted from this table as it is extremely variable and readily changes with fuel storage; for land-based harvested biomass, levels of 40–60 % moisture may be found, but these can be reduced during storage to 15–20 %, and further with fuel processing (EUBIONET2, 2007). As such, the initial moisture contents are much higher than for traditional power station coals. The ash contents of most biomass are generally lower than for coals, though some types of residual biomass can have relatively high ash contents (e.g. palm kernels in Table 1.10). For other traditional fuel parameters (e.g. CV, C, H and O contents), when expressed on a dry ash free basis they fall within
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Solid fuel composition and power plant fuel flexibility Table 1.9
23
Examples of biomass analysis standards
Standard reference
Title
CEN/TS 14588:2004 Solid Biofuels – Terminology, Definitions and Descriptions CEN/TS 14774-1:2004 Solid Biofuels – Methods for Determination of Moisture Content – Oven Dry Method – Part 1: Total Moisture – Reference Method CEN/TS 14774-2:2004 Solid Biofuels – Methods for the Determination of Moisture Content – Oven Dry Method – Part 2: Total Moisture – Simplified Method CEN/TS 14774-3:2004 Solid Biofuels – Methods for the Determination of Moisture Content – Oven Dry Method – Part 3: Moisture in General analysis sample CEN/TS 14775:2004 Solid Biofuels – Method for the Determination of Ash Content CEN/TS 14778-1:2005 Solid Biofuels – Sampling – Part 1: Methods for Sampling CEN/TS 14778-2:2005 Solid Biofuels – Sampling – Part 2: Methods for Sampling Particulate Material Transported in Lorries CEN/TS 14779:2005 Solid Biofuels – Sampling – Methods for Preparing Sampling Plans and Sampling Certificates CEN/TS 14780:2005 Solid Biofuels – Methods for Sample Preparation CEN/TS 14918:2005 Solid Biofuels – Method for the Determination of Calorific Value CEN/TS 14961:2005 Solid Biofuels – Fuel Specifications and Classes CEN/TS 15103:2005 Solid Biofuels – Methods for the Determination of Bulk Density CEN/TS 15104:2005 Solid Biofuels – Determination of Total Content of Carbon, Hydrogen and Nitrogen – Instrumental Methods CEN/TS 15105:2005 Solid Biofuels – Methods for Determination of the Water Soluble Content of Chloride, Sodium and Potassium CEN/TS 15148:2005 Solid Biofuels – Method for the Determination of the Content of Volatile Matter CEN/TS 15149-1:2006 Solid Biofuels – Methods for the Determination of Particle Size Distribution – Part 1: Oscillating Screen Method Using Sieve Apertures of 3.15 mm and Above CEN/TS 15149-2:2006 Solid Biofuels – Methods for the Determination of Particle Size Distribution – Part 2: Vibrating Screen Method Using Sieve Apertures of 3.15 mm and Below CEN/TS 15149-3:2006 Solid Biofuels – Methods for the Determination of Particle Size Distribution – Part 3: Rotary Screen Method CEN/TS 15150:2005 Solid Biofuels – Methods for the Determination of Particle Density CEN/TS 15210-1:2005 Solid Biofuels – Methods for the Determination of Mechanical Durability of Pellets and Briquettes – Part 1: Pellets CEN/TS 15210-2:2005 Solid biofuels – Methods for the Determination of Mechanical Durability of Pellets and Briquettes – Part 2: Briquettes CEN/TS 15234:2006 Solid Biofuels – Fuel Quality Assurance CEN/TS 15289:2006 Solid Biofuels – Determination of Total Content of Sulphur and Chlorine CEN/TS 15290:2006 Solid Biofuels – Determination of Major Elements CEN/TS 15296:2006 Solid Biofuels – Calculation of Analyses to Different Bases CEN/TS 15297:2006 Solid Biofuels – Determination of Minor Elements CEN/TS 15370-1:2006 Solid Biofuels – Method for the Determination of Ash Melting Behaviour – Part 1: Characteristic Temperatures Method
Source: compiled from CEN (2010).
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Table 1.10
Typical biomass analyses
Parameter Unit
SRC willow
Ash wt% dry 2 CV (gross) MJ/kg 20.3 daf CV (net) 18.8 C wt% daf 49 H 6.2 O 44 N 0.5 S 0.05 Cl 0.03
Miscanthus
Straw (wheat, barley, rye)
Coniferous Palm Olive wood kernel residue (logging residues)
4 19.8 18.4 49 6.4 44 0.7 0.2 0.2
5 19.8 18.5 49 6.3 43 0.5 0.1 0.4
7.5 19.8 18.3 50.2 6.6 40 3.2 0.2 0.2
4.5 21.4 18.3 49 6.0 40 2.24 0.1 0.1
50–700 2000–7000 100–500 2000–26000 400–1300
750 3000 2500 3000 3000
1500 6000 900 23000 2000
Elemental analysis (mg/kg dry basis) Al 3–1000 40–600 Ca 2000–9000 900–3000 Fe 30–600 40–400 K 1700–4600 1000–11000 Mg 200–800 300–900 Mn 80–160 200–500 Na 10–450 400–1200 P 500–1300 2000–10000 Si 2–7200 10–50
500–3000 200 300–2900 7000 1000–20000 3000
100 1500 5000
2 21 20 52 6.1 41 0.5 0.04 0.01
2000–8000 1000–4000 400–2000 250 75–300 500 200–10000
Source: compiled from CENTS 14961:2005.
a relatively narrow range of values, as illustrated in Table 1.10. Compared to traditional power station coals (e.g. Table 1.6), biomass have less C, but more H and O. This difference in basic fuel composition can also be seen in the trend from biomass through peat and coals to anthracite, which was presented in Fig. 1.3. The levels of other elements (including N, S, Cl, P and various metallic elements) have been found to vary widely between biomass types and the various parts of plants (e.g. heart wood compared to bark and leaves from the same tree), as well as with growing conditions, the timing of harvesting, biomass storage, etc. (White and Plaskett, 1981; EUBIONET2, 2007; van Loo and Koppeian, 2007; Simms et al., 2007a; Doran, 2009; Livingston, 2009). These are illustrated in Table 1.10 with an indication of the ranges found for some elements, in particular biomass. However, some useful trends in biomass compositions can be identified: . .
Biomass have low S levels compared to most coals (Fig. 1.4). The Cl contents of biomass (up to 2.5 wt% daf) span a wider range than for coals. Some biomass can have higher Cl contents than coals, but many do not; the faster growing biomass (e.g. cereal crops) tend to have the higher Cl contents.
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1.6 Selected minor and trace elements present in a range of different biomass (Simms et al., 2007b).
.
.
In biomass, the alkali metals are mostly present in a different chemical form than for coals (Livingston, 2009), and so are much more readily released during combustion. There is also a different balance between the alkali metals, with much more K than Na; as with Cl, higher K levels are found in faster growing biomass (Simms et al., 2007a). Other elements can be found at high levels in selected biomass (e.g. N, P, Ca), which can lead to difficulties in using these potential fuels in some power systems (due to emissions, the formation of fouling deposits and/ or corrosion issues). Figure 1.6 illustrates the relative abundance of selected elements in biomass compared with a typical bituminous power station coal (Simms et al., 2007a).
The compositions of many types of biomass indicate that they have the potential to be used as fuels, but their differences compared to coal highlight the need for care in choosing specific biomass for use in power generation systems (especially those that have been developed and designed for specific coals). The differences in the minor elements present in biomass can result in numerous issues related to emissions, deposition (fouling/slagging) and corrosion, causing both operational and maintenance issues, as well as restricting the efficiencies of the biomass-fired power generation systems (compared to coal-fired systems) (Simms et al., 2007a; Livingston, 2009). Other important practical issues relate to the preparation of biomass fuels into forms suitable for use in power generating systems. It is necessary to process the biomass, using cutting, shredding, milling, drying, etc., to
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produce the fuel in the form needed for its use (Doran, 2009). Thus, fuel storage and handling issues are critical for biomass use, and different to those encountered with coals due to their different physical properties (Francis and Peters, 1980; White and Plaskett, 1981; EUBIONET2, 2007; van Loo and Koppeian, 2007; Simms et al., 2007a; Doran, 2009; Livingston, 2009).
1.2.4 Waste derived fuels Wastes are usually classified in terms of their origins, but the classification used varies around the world. As part of its activities to assess the emissions of greenhouse gases, the UN International Panel on Climate Change carried out an extensive assessment of the production and use of wastes around the world (Pipatti et al., 2006). It used the following categories for waste streams: . . . .
municipal solid wastes (MSW); industrial wastes; sludges; other wastes.
It was noted that many other waste classifications could be used and that some waste streams could be allocated to different categories on a national or regional basis; for example, commercial and demolition wood could be classified as industrial wastes, municipal solid wastes or put into its own category depending on the location of the waste (Pipatti et al., 2006). Table 1.11 summarises results for municipal solid wastes (MSW) for different regions worldwide from this report. For the purposes of this study, MSW was taken to include the following categories (Pipatti et al., 2006): . . . . . . . . . . .
food waste; garden (yard) and park waste; paper and cardboard; wood; textiles; nappies (disposable diapers); rubber and leather; plastics; metal; glass (and pottery and china); other (e.g. ash, dirt, dust, soil, electronic waste).
As a result of the different classifications used around the world, there are notable gaps in the data produced (Table 1.11). In contrast, Table 1.12 gives an alternative breakdown of the composition
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26.2 40.3 43.5 41.1
53.9 43.4 51.1 23 40.4
30.1 23.8 36.9 24.2
36.0 67.5
33.9 43.8 44.9 46.9
Asia Eastern Asia South-central Asia South-eastern Asia Western Asia and Middle East
Africa Eastern Africa Middle Africa Northern Africa Southern Africa Western Africa
Europe Eastern Europe Northern Europe Southern Europe Western Europe
Oceania Australia and New Zealand Rest of Oceania
America North America Central America South America Caribbean 23.2 13.7 17.1 17.0
30.0 6.0
21.8 30.6 17.0 27.5
7.7 16.8 16.5 25 9.8
18.8 11.3 12.9 18.0
Paper/cardboard
6.2 13.5 4.7 2.4
24.0 2.5
7.5 10 10.6 11
7.0 6.5 2 15 4.4
3.5 7.9 9.9 9.8
Wood
3.9 2.6 2.6 5.1
4.7 2.0
1.0
1.7 2.5 2.5
3.5 2.5 2.7 2.9
Textiles
1.4 1.8 0.7 1.9
1.4
1.1
1.0 0.8 0.9 0.6
Rubber/leather
8.5 6.7 10.8 9.9
6.2 13.0
3.0
5.5 4.5 4.5
14.3 6.4 7.2 6.3
Plastic
4.6 2.6 2.9 5.0
3.6 7.0
1.0
1.8 3.5 3.5
2.7 3.8 3.3 1.3
Metal
6.5 3.7 3.3 5.7
10.0 8.0
2.3 2.0 2
3.1 3.5 4.0 2.2
Glass
9.8 12.3 13.0 3.5
14.6
11.6 1.5 1.5
7.4 21.9 16.3 5.4
Other
Note 1. Data are based on weight of wet waste of MSW without industrial waste at generation around year 2000. Note 2. The region-specific values are calculated from national, partly incomplete composition data. The percentages given may therefore not add up to 100 %. Some regions may not have data for some waste types – blanks in the table represent missing data. Source: Pipatti et al. (2006).
Food waste
MSW composition data in terms of fuel origins
Region
Table 1.11
Solid fuel composition and power plant fuel flexibility 27
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Table 1.12
Example of more detailed breakdown of MSW from UK survey
Category
Weight percentage
Paper and card Kitchen and garden waste Glass Textiles Plastics Ferrous metals Non-ferrous metals Miscellaneous combustible Disposable nappies Miscellaneous non-combustible Fines
33.2 20.2 9.3 2.1 10.2 5.7 1.6 8.1 3.9 1.8 6.8
Source: Burnley (2007).
of municipal solid waste (MSW) generated in the UK from a detailed survey of regional waste generation (Burnley, 2007). Such surveys emphasise the extreme variability that can be found in waste streams. However, not all waste streams are destined to be considered as potential fuels, with the ‘waste hierarchy’ favouring waste reduction, re-use and re-cycling over energy recovery. The materials that are left for use in energy recovery systems can be used as raw fuels, or further processed to generate refuse derived fuels (RDFs) or solid recovered fuels (SRFs). In recent years, to encourage the use of such fuels, there has been a move towards the standardisation of such fuels by the European Committee for Standardisation (CEN). As a result, Simms et al. (2007a) indicate that this has produced a system for classifying solid recovered fuels (SRFs) in terms of their: . . .
mean net CV; mean value of the chlorine content; median and 80th percentile values of the mercury content (on an asreceived basis).
A series of standards for SRFs has now been developed by the CEN, and examples of these are listed in Table 1.13. These cover analytical methods to generate measurements of all the solid fuel properties required for comparison with coal and biomass fuels. However, as a result of the variable feedstocks, the fuels produced have variable properties, as illustrated in Table 1.14. Thus, the performance of such fuels cannot be generalised and has to be assessed on a site basis when the characteristics of the local fuel(s) are known. However, the availability of the data required to do this in a standardised form enables the methods developed for use with coal and biomass fuel assessment to be translated (but applied with care).
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Solid fuel composition and power plant fuel flexibility Table 1.13
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Examples of European waste analysis standards
Standard reference Title CEN/TS 15357 CEN/TS 15358
CEN/TS 15359 CEN/TS 15440 CEN/TR 14980 CEN/TR 15441 CEN/TS 15400 CEN/TS 15401 CEN/TS 15402 CEN/TS 15403 CEN/TS 15404 CEN/TS 15405 CEN/TS 15406 CEN/TS 15407 CEN/TS 15408
CEN/TS 15410 CEN/TS 15411
CEN/TS 15412 CEN/TS 15413 CEN/TS 15442 CEN/TS 15443 CEN/TS 15414-1
CEN/TS 15414-2
CEN/TS 15414-3
Solid Recovered Fuels – Terminology, Definitions and Descriptions Solid Recovered Fuels – Quality Management Systems – Particular Requirements for Their Application to the Production of Solid Recovered Fuels Solid Recovered Fuels – Specifications and Classes Solid Recovered Fuels – Method for the Determination of Biomass Content Solid Recovered Fuels – Report on Relative Difference Between Biodegradable and Biogenic Fractions of SRF (TR) Solid Recovered Fuels – Guidelines on Occupational Health Aspects (TR) Solid Recovered Fuels – Methods for the Determination of Calorific Value Solid Recovered Fuels – Methods for the Determination of Bulk Density Solid Recovered Fuels – Methods for the Determination of the Content of Volatile Matter Solid Recovered Fuels – Methods for the Determination of Ash Content Solid Recovered Fuels – Methods for the Determination of ASH MELTING BEHAVIOUR Solid Recovered Fuels – Methods for the Determination of the Density of Pellets and Briquettes Solid Recovered Fuels – Methods for the Determination of Bridging Properties of Particulate Solid Recovered Fuels Solid Recovered Fuels – Methods for the Determination of Carbon (C), Hydrogen (H) and Nitrogen (N) Content Solid Recovered Fuels – Methods for the Determination of Sulphur (S), Chlorine (Cl), Fluorine (F) and Bromine (Br) Content Solid Recovered Fuels – Method for the Determination of the Content of Major Elements (Al, Ca, Fe, K, Mg, Na, P, Si, Ti) Solid Recovered Fuels – Methods for the Determination of the Content of Trace Elements (As, Ba, Be, Cd, Co, Cr, Cu, Hg, Mo, Mn, Ni, Pb, Sb, Se, Tl, V and Zn) Solid Recovered Fuels – Methods for the Determination of Metallic Aluminium Solid Recovered Fuels – Methods for the Preparation of the Test Sample From the Laboratory Sample Solid Recovered Fuels – Methods for Sampling Solid Recovered Fuels – Methods for Laboratory Sample Preparation Solid Recovered Fuels – Determination of Moisture Content Using the Oven Dry Method – Part 1: Determination of Total Moisture by a Reference Method Solid Recovered Fuels – Determination of Moisture Content Using the Oven Dry Method – Part 2: Determination of Total Moisture by a Simplified Method Solid Recovered Fuels – Determination of Moisture Content Using the Oven Dry Method – Part 3: Moisture in General Analysis Sample (Continued)
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Table 1.13 (cont.) Standard reference Title CEN/TS 15415 CEN/TR 15508 CEN/TR 15716 CEN/TS 15590
CEN/TR 15591 CEN/TS 15639
Solid Recovered Fuels – Determination of Particle Size and Particle Size Distribution by Screen Method Key Properties of Solid Recovered Fuels to be Used for Establishing a Classification System (CEN/TR) Solid Recovered Fuels – Determination of Combustion Behaviour (TR) Solid Recovered Fuels – Determination of Potential Rate of Microbial Self Heating Using the Real Dynamic Respiration Index Solid Recovered Fuels – Determination of the Biomass Content Based on the 14C Method (TR) Solid Recovered Fuels – Methods for the Determination of Mechanical Durability of Pellets
Source: compiled from CEN (2010). Table 1.14
Typical waste fuel analyses
Component
Units
Minimum value
Mean value
Maximum value
Water content Volatiles Ash CV (gross) CV (net) C H O N S Cl
wt% wet wt% daf wt% dry kJ/kg daf kJ/kg daf wt% daf wt% daf wt% daf wt% daf wt% daf wt% daf
2.9 74.6 4.4 13130 12126 33.9 1.72 15.8 0.12 0.01 0.006
14.6 88.7 17 24597 22915 54.8 8.12 34 0.94 0.4 0.716
38.7 99.4 44.2 44029 40986 84.8 15.16 43.7 2.37 1.4 1.558
1.3
Use of alternative fuels in combustion power plants and application of technology to improve fuel flexibility
Traditionally pulverised coal combustion power plants have tended to be large scale, with individual boilers of 500–800 MWe grouped together (to give total power plant capacities of 2000–4000 MWe). Plants that were built in the 1960s–1970s often have steam systems that operate with maximum steam parameters of approximately 140–160 bar/540–560 8C and now operate with efficiencies of ~35–37 % (following various upgrades and environmental protection measures that have respectively had the effect of increasing and decreasing system efficiencies over the years). New coal
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systems use individual boilers of similar sizes, but with steam systems with maximum operating parameters of approximately 290 bar/620 8C giving efficiencies of 46–47 % (Farley, 2010). In contrast, biomass combustion plants tend to be much smaller, with newly built power plants of up to ~30 MWe and with efficiencies of ~30 % (Doran, 2009). Waste-to-energy plants are also much smaller, with electrical generating capacities of up to 30 MWe and efficiencies of ~25 % (Prewin, 2011). However, the smaller scale and potential locations of biomass and waste generating plants lend themselves to possible use in combined heat and power applications (if appropriate local heat loads can be found), which give increased efficiencies. The use of the alternative solid fuels either alone or by co-firing is influenced by many factors, including: . . . . . . .
fuel costs/subsidies; fuel availability; suitability of alternative fuel for existing process; fuel storage/handling/preparation; power plant efficiencies and scale; need for local use of heat/power; need for additional gas cleaning facilities to meet environmental regulations.
The compositions of the fuels particularly influence points 2 to 4 (e.g. Maciejewska et al., 2006). However, it has been found that certain combinations of fuels are not desirable when co-firing due to their tendencies to increase deposition in different locations along the combustion system hot gas path and also to increase corrosion damage to heat exchangers.
1.3.1 Fuel substitution The availability of biomass and waste fuels, as well as the significant differences in both their physical and chemical properties, has guided their use as single fuels towards dedicated power plants. Many new biomass and waste-to-energy power plants have been built during the last decade and the numbers of these plants are expected to increase significantly in the immediate future as increasing emphasis is placed on switching to renewable and more sustainable fuels. However, for old coal-fired power stations that are being decommissioned as part of current environmental initiatives, one interesting alternative idea to just scrapping them is to adapt them in such a way that biomass could be used to fire one of the boilers (500–600 MWe units) that make up the overall
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station. To make this viable several issues would need to be addressed, including: . . . . . . .
sourcing very large quantities of easily transportable biomass; using a power station suitably located to receive all this biomass; replacement of fuel handling, storage and preparation systems; new fuel burners; down-rating the boiler (to use lower final steam temperatures); installation of appropriate gas cleaning systems; long-term economic viability of such a scheme (given the unreliability of regulations and subsidies for sustained periods).
There is a need to improve the efficiencies of power plants firing biomass and waste fuels. One of the key limiting factors in restricting their efficiencies is deposit formation and corrosion on the final superheater. Deposit formation restricts the heat transfer to this heat exchanger and also provides a chemically aggressive environment that rapidly corrodes the heat exchanger materials (Chapter 4). New boilers are designed to try to partially counter these effects, but one new technology developed by Vattenfall is targeted at altering the environments generated in the boilers when using these fuels. The ‘ChlorOut’ process sprays a sulphur rich compound (ammonium sulphate) into the gas stream and controls this on the basis of minimising the alkali metal chlorides that are present in the gas stream (Vattenfall Research and Development, 2005). Initial trials have shown this to be effective in reducing both deposition and fireside corrosion with fuels rich in alkali chlorides (Vattenfall Research and Development, 2005).
1.3.2 Co-firing fuels Co-firing of biomass fuels in previously coal-fired power plants has proved to be a successful route to introduce significant quantities of biomass into the power generation market. In the UK, the levels of biomass co-firing have steadily increased during the last 10 years up to 10 % (on an energy basis) for some biomass–coal mixtures, with the use of still higher levels being actively investigated. The use of 10 % biomass in a 2000 MWe power plant represents 200 MWe of biomass derived power (and should be compared to the ~30 MWe output of a new biomass–plant fired operating at a lower efficiency). The coal–biomass mixes that can be used in such systems are limited by a number of factors: .
Fuel transport/handling/storage systems designed for coal. This includes ships, trains or lorries for fuel transport and external storage for coal or internal storage for biomass fuels.
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.
.
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Fuel preparation systems. Coal grinding mills can tolerate a few % biomass in a mixed fuel feed, but at higher levels separate dedicated biomass mills are required and then the two prepared fuel streams need to be blended. Combustion systems. Adding biomass to existing coal feeds and then burning the mixed fuels in existing burners (designed for coal use only) is one alternative; another is not to mix the fuel streams and then to use separate biomass and coal burners distributed evenly around the boiler. Slagging/fouling/corrosion. Differences in fuel compositions can cause increased rates of deposition and different deposit compositions with some biomass/coal mixtures. Such differences can result in reduced heat transfer/more frequent cleaning requirements, increased corrosion damage and ultimately shorter component lives coupled with reduced boiler reliability (Davis and Pinder, 2004; Livingston, 2010; Simms et al., 2007b).
In order to minimise the risks associated with the introduction of cofiring, technology developments have focused on all of these topics (Davis and Pinder, 2004; Overgaard et al., 2004; Simms et al., 2007b; Livingston, 2010; Waldron, 2010). For activities related to combustion and slagging/ fouling/corrosion these have included trials on both pilot plants and power station boilers, which have focused on specific coal/biomass combinations and included thorough monitoring of the power plants (e.g. Henderson et al., 2002). In addition, more fundamental supporting research has been carried out to gain a better understanding of the processes involved with multiple fuels, development of predictive models and discovery of approaches that can be used in controlling them (covered in Chapter 4). In particular there are opportunities to minimise the risks involved by the careful selection of combinations of coal and biomass fuels.
1.4
Future trends
Given the current views on the causes of global warming and the future availability of fossil fuels, it seems inevitable that increasing levels of biomass and wastes will be used to generate heat and power in the future. Government policies and regulations are now increasingly reflecting and reenforcing these trends, with challenging targets being set for the use of renewable fuels; for example, the European Union has targets for a 20 % reduction in CO2 emissions by 2020 and a 50 % reduction by 2050 (Dechamps, 2006; Farley, 2007, 2010). Co-firing biomass and waste derived fuels with coal offers one route to using such fuels in higher efficiency systems than can be used for the individual fuels. It is expected that there will continue to be significant
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challenges related to fuel sourcing, storage, handling preparation, feeding, even combustion, deposition, corrosion and environmental emissions (solid, liquid and gases). However, if power generating efficiencies are to be maintained, then these challenges are expected to increase as higher shares of the minor fuels are used, e.g. requiring different approaches to fuel feeding than is possible with 5–10 % additions. It may be possible to reduce the impact of some of these issues by careful matching of fuel compositions when creating the fuel mixes, or by using fuel additives. A possible future development is to re-start coal plants (decommissioned to reduce environmental emissions) with 100 % biomass fuel feeding. This would maximise the use of current power generating capacity, but would require modifications to enable biomass storage, handling, preparation, combustion, etc. Such systems could try to minimise the effects of the biomass compositions on deposition and corrosion by blending two (or more) biomass, or by using appropriate chemical additives. Finally, carbon capture and storage (CCS) is being investigated for use on large scale coal-fired power systems. There are several CO2 capture technologies that are at different stages in their development, including systems for either pre-combustion or post-combustion CO2 removal (Dechamps, 2006; Farley, 2010). Gasification of fuels (using either oxygen or steam as the oxidant) produces a gas that can be conditioned to enable pre-combustion CO2 removal. Post-combustion CO2 capture can be carried out using either solid or aqueous sorbent processes. Oxy-firing of fuels is a technology that would enable more efficient post-combustion CO2 capture. The various possible influences of all aspects of fuel compositions on these processes have yet to be determined, but research and development of specific processes with specific fuel compositions is currently active in many countries.
1.5
Sources of further information and advice
Fuel standards can be obtained from national and international standard institutions, for example: . . .
ASTM: www.astm.org/index.shtml BSI: www.bsigroup.com/en/ CEN: www.cen.eu/cen/Pages/default.aspx
For specific fuels, there are several books and websites that provide far more detailed information than was possible in this general introductory chapter. These books include: .
J G Speight, Handbook of Coal Analyses (John Wiley, 2005)
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S van Loo and J Koppeian, The Handbook of Biomass Combustion and Co-firing (Earthscan 2007) British Coal Utilisation Research Association (BCURA): http://www. bcura.org/ World Coal Institute: http://www.worldcoal.org/ International Energy Agency (IEA) Clean Coal Centre: http://www. ieacoal.org.uk/site/ieacoal/home
For fuel compositions, there are several databases that can be accessed via the internet, for example: . . . .
1.6
Energy Research Centre of the Netherlands, Phyllis, Database for Biomass and Waste: http://www.ecn.nl/phylli US Department of Energy, Energy Efficiency and Renewable Energy, Biomass Program, Biomass Feedstock Composition and Property Database: www1.eere.energy.gov/biomass/feedstock_databases.html BIOBIB – A Database for Biofuels, University of Technology Vienna: www.vt.tuwien.ac.at/biobib/biobib.html IEA Bioenergy Task 32, Biomass Combustion and Co-firing: www. ieabcc.nl/
References
ASTM International (2010): www.astm.org/index.shtml. ASTM-D3174-11 (2011), Standard Test Method for Ash in the Analysis Sample of Coal and Coke from Coal. British Standards Institute (BSI) (2010): www.bsigroup.com/en/. BS 1016 – Part 20 (1981), Hardgrove Grindability. BS 1016 (2010). Burnley S J (2007), ‘A review of municipal solid waste composition in the United Kingdom’, Waste Management, 27, 1274–1285. Carpenter A M and Skorupska N M (1993), Coal Combustion – Analysis and Testing, IEA Coal Research, IEACR/64. Clarke L B and Sloss L L (1992), Trace Elements – Emissions from Coal Combustion and Gasification, IEA Coal Research, IEACR/49. Colechin M (2005), Best Practice Brochure: Co-firing of Biomass (Main Report), DTI Report No. COAL R287, DTI Pub URN 05/1160 (2005). Commission of the European Communities (COM) (2005), 628 Final, Biomass Action Plan (2005). Couch G R (1995), Power from Coal – Where to Remove Impurties?, IEA Coal Research, IEACR/82. Council of the European Union (1999), Directive 1999/31/EC on the Landfill of Waste, Official Journal of the European Communities, vol. L 182, pp. 1–19. Davis C J and Pinder L W (2004), Fireside Corrosion of Boiler Materials – Effect of Co-Firing Biomass with Coal, UK Department of Trade and Industry, Report No. COAL R267 DTI/Pub URN 04/1795.
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Dechamps P (2006), ‘The EU research strategy towards zero emission fossil fuel power plants’, in Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Materials for Advanced Power Engineering 2006, Forschungszentrum Ju¨lich GmbH, pp. 25–40. Doran M (2009), ‘Feedstocks for thermal conversion’, in Bridgwater A V, Hofbauer H and van Loo S (eds), Thermal Biomass Conversion, CPL Press, pp. 129–156. EUBIONET2 (2007), Biomass Fuel Supply Chains for Solid Biofuels – From Small to Large Scale, VTT Jyva¨skyla¨, Finland. Euracoal (2010): http://www.euracoal.be/pages/layout1sp.php?idpage=2 [accessed 4 November 2010]. European Committee for Standardisation (CEN) (2010): www.cen.eu/cen/Pages/ default.aspx. Farley M (2007), ‘Clean coal technologies for power generation’, in Strang A, Banks W M, McColvin, G M, Oakey J E and Vanstone R W (eds), Parsons 2007: Power Generation in an Era of Climate Change, IoM Communications, pp. 335– 342. Farley M (2010), Overview of Capture Technologies for Pulverised Coal-Oxyfuel and Post Combustion Capture, Doosan Power Systems. Available from: http://www. specialmetalsforum.com/uploads/docs/12754049682.DoosanMFNamtec Harrogate2010.pdf [accessed 4 November 2010]. Foster D J, Livingston W R, Wells J, Williamson J, Gibb W H and Bailey D (2004), Particle Impact Erosion and Abrasion Wear – Predictive Methods and Remedial Measures, UK DTI, Report No. COAL R241 DTI/Pub URN 04/701. Francis W and Peters M C (1980), Fuel and Fuel Technology, Pergamon Press. Henderson P J, Karlsson A, Davis C, Rademakers P, Cizner J, Formanek B, Gorannsson K and Oakey J (2002), ‘In-situ corrosion testing of advanced boiler materials with diverse fuels’, in Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Materials for Advanced Power Engineering 2002, Forschungszentrum Ju¨lich GmbH, pp. 785–800. IEA Bioenergy (2003), Task 23. Co-firing of MSW and RDF, VTT Energy, Fuels and Combustion, New Energy Technologies http://www.ieabioenergytask36.org/ Publications/1998–2001%20Task%2023/Publications/ Cofiring_of_MSW_and_RDF.PDF. International Organisation for Standardisation (ISO) (2010): www.iso.org/iso/home. htm. Jones A (2005), 54th BCURA Robens Coal Science Lecture http://www.bcura.org/ csl05.pdf [accessed 4 November 2010]. Livingston W R (2009), ‘Fouling corrosion and erosion’, in Bridgwater A V, Hofbauer H and van Loo S (eds), Thermal Biomass Conversion, CPL Press, pp. 157–176. Livingston W R (2010), ‘Advanced biomass co-firing technologies for coal-fired boilers’ (publication details to be confirmed). Maciejewska A, Veringa H, Sanders J and Peteves S D (2006), Co-Firing of Biomass with Coal: Constraints and Role of Biomass Pre-Treatment, European Commission, EUR 22461 EN. Nalbandian H (2004), Air Pollution Control Technologies and Their Interactions, IEA Clean Coal Centre, CCC/92. Overgaard P, Sander B, Junker H, Friborg K and Larsen O H (2004), ‘Two years’
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operational experience and further development of full-scale co-firing of straw’, 2nd World Conference on Biomass for Energy, Industry and Climate Protection, Rome, May 2004. Phyllis (2010), Phyllis, Database for Biomass and Waste, Energy Research Centre of the Netherlands http://www.ecn.nl/phyllis. Pipatti R, Sharma C and Yamada M (2006), IPCC Guidelines for National Greenhouse Gas Inventories, Volume 5: Waste, Chapter 2: ‘Waste generation, composition and management data’, UN IPCC. Prewin (2011), ‘Performance, reliability and emissions reduction in waste incinerators’. Available from: http://www.prewin.eu/ [accessed 16 May 2011]. Raask E (1985), Mineral Impurities in Coal Combustion, Hemisphere Publishing Corporation. Simms N J, Kilgallon P J and Oakey J E (2007a), ‘Fireside issues in advanced power generation systems’, Energy Materials: Materials Science and Engineering for Energy Systems, 2, 154–160. Simms N J, Kilgallon P J and Oakey J E (2007b), ‘Degradation of heat exchanger materials under biomass co-firing conditions’, Materials at High Temperatures, 24, 333–342. Sloss L (1992), Halogen Emissions from Coal Combustion, IEA Coal Research, IEACR/45. Speight J G (1994), The Chemistry and Technology of Coal, M Dekker Inc. Speight J G (2005), Handbook of Coal Analyses, John Wiley. Stringer J (1995), ‘Practical experience with wastage at elevated temperatures in coal combustion systems’, Wear, 186–187, 11–27. UK Department of Trade and Industry (2007), Meeting the Energy Challenge, A White Paper on Energy. UK Department of Energy and Climate Change (DECC) (2009), UK Energy in Brief 2009, DECC/Pub 8876/3.5k/07/09/NP. URN 09D/220. van Loo S and Koppeian J (2007), The Handbook of Biomass Combustion and Cofiring, Earthscan. van Krevelen D W (1950), ‘Graphical – statistical method for the study of structure and reaction processes of coal’, Fuel, 29, 269–284. Vattenfall Research and Development (2005), ChlorOut, Available from: http:// www.vattenfall.com/en/file/ChlorOut_8459980.pdf [accessed 4 November 2010]. Waldron D (2010), ‘Options for biomass firing in utility boilers’, in Proceedings of Bioten Conference (to be published). White L P and Plaskett L G (1981), Biomass as Fuel, Academic Press. World Coal (2010): http://www.worldcoal.org/coal/what-is-coal/ [accessed 4 November 2010].
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2 Condition monitoring and assessment of power plant components C . D e M I C H E L I S , Independent Consultant, previously CESI, Italy, C . R I N A L D I , RSE, Italy, C . S A M P I E T R I , EXOVA, Italy and R . V A R I O , CESI, Italy
Abstract: This chapter describes condition monitoring and assessment methods used in thermal power plants and provides additional diagnostic information, complementary to that provided by plant control systems. The chapter sections deal with the main power plant items: the boiler (B) and heat recovery steam generator (HRSG), steam turbine (ST) and gas turbine (GT). For each item, the most relevant damage mechanisms are briefly mentioned and selected diagnostic methods, applicable to on-line monitoring or to maintenance inspections, are presented. Finally, an attempt is made to identify the most promising trends and future developments in these areas. Key words: power plant condition monitoring, on-line monitoring methods, non-destructive techniques, boiler, steam turbine, gas turbine, rotor, hot parts, material condition assessment, coating degradation.
2.1
Introduction
Optimized overall life-cycle costs, availability, energy efficiency and environmental performance are major objectives in a competitive global power generation market. Whereas major breakthroughs in the above areas can derive only from corresponding breakthroughs in process and material innovation, a substantial contribution to these objectives comes from an intelligent and cost-effective monitoring and assessment of the current health (integrity) and performance of the key components of fossil-fired power plants (PPs). Moreover, the adoption of reliable on-line condition monitoring techniques is a pre-condition for the implementation of
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condition-based maintenance practices, to optimize long-term maintenance planning and, more generally, asset life-cycle management. Plant condition monitoring (PCM) encompasses monitoring both process performance and structural integrity of main PP components. This task implies the combined use of appropriate – conventional and innovative – on-line condition monitoring sensors and instrumentation, in situ and offline non-destructive testing methods (NDTs), as well as laboratory testing and modelling tools (material damage, temperature/stress distribution, combustion process, heat exchange, etc.). With regards to structural integrity assessment, the usually higher cost and lower diagnostic reliability of on-line monitoring methods with respect to off-line NDT methods is more than compensated for by their provision of a time-continuous mitigation of structural risk. Only problem-specific operational experience can ensure that the adopted combination of on-line and off-line methods achieves the best cost–benefit compromise. Until about 1980 PCM systems supplied by the constructor (OEM) were basic and simplified, so that any improvement of PCM was often the concern of large PP operators, as a way to reduce costs and optimize profitability of their power generation assets. The revolution of the last decades in data acquisition, handling, processing, storage and long-distance communication has provided great improvements in the availability of performance and effective PCM systems. Today’s tendency to transfer the responsibility for long-term integrity and performance of PP components from the owner to the OEM by means of long-term global service contracts is shifting the effort for the development and implementation of costeffective PCM tools from the owner to the global service provider. Modern PPs are equipped with built-in, standard monitoring instrumentation, aimed at controlling the main process parameters (temperatures, pressures, flows, emissions, etc. see Fig. 2.1 and Table 2.1), in order to keep plant energy efficiency and emissions under control (as required by regulatory bodies). Little or no specific monitoring of actual structural degradation of key components is usually provided. Furthermore, some specific process monitoring aspects are not routinely monitored, although they can substantially impact the overall plant performance. This chapter covers some specific PCM techniques, complementary to those delivered by default by the OEMs to the utilities. Due to the extension of the subject, attention concentrates on monitoring and condition assessment of components of fossil-fuelled PPs (gas-, oil- and coal-fired, USC, combined cycle), while innovative PP designs are covered elsewhere in this book. A limited number of PP components of special relevance in terms of reliability, availability and cost are treated, namely:
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2.1 Levels of plant condition monitoring: (1) basic process control (sensors and actuators) through a distributed control system (DCS) + vibration monitoring system (VMS), built into the plant by the OEM; (2) condition monitoring system (CMS) supporting performance and maintenance optimization d = displacement; v.a. = vibration amplitude); (3) strategic plant control (SPC) completely automated system, similar to airplane systems, not yet implemented in power plants.
(a) (b) (c) (d)
boiler, superheater and reheater tubes; high-energy vessels and piping (headers, steam lines); steam turbines; gas turbines.
Special attention is given to the level of industrial acceptance of the methods discussed and to their acceptance in recommended practices or in national or international regulations and standards. The combined use of methods and tools for on-line condition monitoring and in situ, off-line condition assessment yields diagnostic insight that can be profitably used in the frame of fitness-for-service or life assessment/life extension procedures. In this chapter, however, neither fitness-for-service nor life assessment procedures nor models are treated; they are only briefly mentioned in order to place the assessment and monitoring methods in their proper frame.
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Monitoring boiler and heat recovery steam generator
Boiler supply normally includes a distributed control system (DCS) controlling the basic combustion and heat exchange processes by continuously monitoring the relevant process parameters: temperatures, pressures, mass flows and chemical species concentrations at crucial locations in both the water–steam cycle and the combustion cycle. Several model-based diagnostic systems exploit the above-mentioned set of data to provide insight on proper operation of the boiler or, where possible, guidance for corrective action and/or optimized performance. Due to the number and variety of components, to the diversity and complexity of the physical and chemical conditions alongside the boiler and to the corresponding variety of potential damage and failure mechanisms, widespread monitoring of individual components is not feasible and basic process monitoring is a necessary first approach, one that can timely detect potentially harmful deviations from recommended operating conditions and help to prevent structural failures. Nevertheless, failure of boilers still accounts for a substantial fraction – ~ 3 % in absolute value (Anon., 2010) – of total unplanned unavailability of thermal power units. Improvements can be sought in several ways. Previous operation experience or, conversely, process/design innovations, for which insufficient experience is available, can strongly motivate ‘a priori’ attention to assumedly critical items and justify an additional targeted monitoring effort using component-specific methods (sections 2.2.1 to 2.2.4). Inspections during maintenance shutdowns are also essential to assess the condition of actual equipment (section 2.2.5). Current critical problems/ items in individual power units are identified, thereby enabling the monitoring effort and available monitoring tools to be concentrated in areas where they are in fact most required. In-service inspections and on-line monitoring must be definitely regarded as coordinated actions within a unified strategy to ensure safe, efficient, reliable and cost-effective long-term operation and life-cycle management of the boiler. Besides the basic process control, advanced on-line structural integrity and performance monitoring of a boiler may include additional monitoring of: . .
integrity (leak, corrosion) of boiler tubes (waterwall (WW), superheater (SH) and reheater (RH)), steam headers and other high-pressure components; temperature and heat exchange in different portions of the boiler (WW panels, SH and RH coils);
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Table 2.1 Levels of plant condition monitoring with typical sensors and delivered information Monitoring level Sensors
Measured parameters
Type of information obtained
DCS = dynamic Thermocouples and control system pressure/position sensors placed at easily reachable positions
T, P measured out of critical regions (e.g. GT exhaust)
. Process control . T in critical regions is indirectly evaluated by process modelling . Inputs for performance monitoring
VPS = vibration monitoring system
Position/ vibration amplitude
. Dynamic control of rotating machinery
Proximity, acceleration or capacity sensors
Component Developed and specific installed on control purpose monitoring systems (CMS) Possibly nonintrusive, e.g. typically optical
Specific . Integrity of component components information . Ageing and damage (e.g. T in very level of component hot and harsh . Optimization of environment maintenance intervals or creep strain)
Strategic control — system
—
. .
Automatic feedback control
temperature distribution and chemical composition of flue gas in different portions of the combustion gas path; burners.
Further issues of specific interest for coal-fired units are monitoring of: . .
fineness and mass-flow balance of the pulverized coal–air mixture (PF) fed to the burners; content of unburnt carbon in fly ashes.
In the following a selection of available monitoring technologies is reviewed, addressing mainly structural integrity problems of crucial boiler components (see possible damage causes in Table 2.1), based on actual diagnostic experience on many thermal power units.
2.2.1 Leak detection Motivations for leak detection and monitoring in boilers Several damage processes can affect boiler tubes, like inside scaling, waterside corrosion and cracking, fireside corrosion and/or erosion, stress rupture due to overheat and creep, vibration-induced and thermal fatigue cracking, and defective welds (see Tables 2.2 and 2.3). If no timely counteraction is adopted these damages end up in a tube wall discontinuity and in a subsequent leak of
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pressurized process fluid. Such components are designed for safe and continued boiler operation in a leak-before-break condition. As a consequence, a monitoring approach aimed at leak detection is indicated. Leaks in boiler, SH and RH tubes can cause substantial efficiency losses, increasing with leak magnitude. Leaks in boiler tubes are somewhat better tolerated in the short term than leaks in SH and RH tubes. Moreover, and perhaps more critically, the high-pressure water or steam ejected from the leaking tube may impinge destructively on the nearby tubes and in a very short time cause unacceptably large damage. These motivations for leak monitoring also hold for other high-energy tubular components of the boiler, such as the high-pressure feedwater preheaters, where leaks have often occurred. Leak detection and monitoring in the boiler is therefore a self-justifying investment. Requirements on leak detection and monitoring methods To be effective, monitoring must first of all provide early detection of the leak: the sooner, the better. A second requirement is that the monitoring method enables – at least in a qualitative way – to control the time evolution of the leak, so as to prevent critical developments (e.g. destructive damage propagation). Thirdly, the method must be highly reliable, since it must help in the process of taking decisions with potentially relevant safety and cost fallouts: false alarms and erratic indications should be ruled out. An additional desirable feature is the capability to identify – even approximately – the leaking component or area. In practice a PP can often manage to tolerate a small or moderate leak in limited portions of the boiler volume for a few days, provided it is constantly and accurately monitored to prevent safety problems and damage extension. Typically, the leaking boiler is operated under strict surveillance till the next incoming period of low power demand (weekend, etc.). The resulting delay, ranging typically from several hours to a few days, enables the PP to properly plan the required repair work – and possibly other additional side tasks – and provide for power replacement in an orderly manner to attenuate the cost penalty of the unplanned outage. Mass flow balance methods Several sophisticated leak detection methodologies rely on models and detailed calculations of the mass flow balance of water and steam in the boiler (e.g. Pertew et al., 2008). They need process data of adequate accuracy to be provided by existing in-plant sensors and meters. Such methods, even if less affected by interfering factors and less prone to false alarms, are less
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sensitive and slower in their diagnostic response than the acoustic techniques described in the following. Methods based on leak-induced sound and vibration Among the methods that can detect leaks early, acoustic-based methods (acoustic leak monitoring, ALM) are the most cost-effective and most widely used in today’s power generation boilers (see Rovder et al., 1988; Kovacevich et al., 1995; also lists of installations of different manufacturers on their web sites). The users’ return of experience is reported to be positive (e.g. Studdard et al., 1992). An ALM system is today a frequent requirement in PP owner’s bids and a standard option in new PPs. The common physical basis of ALM systems is that turbulent outflow of highly pressurized fluid from a metal tube to an open volume (e.g. through a crack) generates a large amount of noise and vibration, as in the case of a leaking furnace, SH or RH tube. The overall released elastic energy increases with pressure drop and leak flow magnitude, the frequency content strongly depending on flaw shape and other parameters that are difficult to control. The vibrations transmitted by the outcoming fluid to the neighbouring atmosphere (e.g. in the combustion chamber) propagate in the form of sound waves in the boiler volume (‘air-borne noise’). Other wide-spectrum vibrations are generated and propagated through the metallic tube wall and its connections to the boiler structure (‘structureborne noise’). Both types of signals can be detected, processed and exploited for diagnosing the potential presence of leaks (see Figs 2.2 to 2.4). Acoustic leak detection and monitoring (‘air-borne noise’) Air-borne noise is detected by conventional microphones, engineered to withstand the harsh condition of the boiler environment. The acoustic signal reaching the sensor is a superposition of the high-level boiler background noise and of the acoustic signal specifically generated by the leak. Both signals are wide-band, with noise amplitude quickly decreasing with frequency, so that absolute sensitivity in not very relevant in this application. The key point for an early and reliable leak detection is the choice of proper filtering to select the most favourable frequency window, usually lying within the audible range (1–20 kHz). In any particular boiler a careful preliminary characterization of background noise is necessary, in order to find where the leak signal/background noise ratio is maximum and to define the proper ‘listening frequency window’. The whole volume of the boiler (combustion chamber and convection area, up to the economizer area, and the penthouse) can be covered by means of 10–20 sensors, according to the boiler size and to the specific
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Example of air-borne leak sensor (courtesy of CESI).
problem requirements. A typical example of sensor locations is shown in Fig. 2.3. Each sensor corresponds to an independent measuring channel, including signal conditioning, filtering, digitizing (these steps are carried out by local instrumentation modules) and transmission to a central display and diagnostic unit, which can be integrated at different levels into the plant supervision systems. Air-borne noise sensors are positioned outside the boiler, but must be physically coupled to the inside boiler volume by means of minimally intrusive acoustic waveguides, in the form of specifically adapted, aircleaned, short tube sections, to be welded to inspection windows or waterwall tube fins (Fig. 2.2). Leak detection and monitoring through microvibrations (‘structure-borne noise’) Structure-borne sensors (essentially high-frequency piezoelectric sensors) are in turn coupled to the metal structure (typically to the SH and RH headers in the penthouse) by means of waveguides, in the form of thin short metal
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2.3 Schematic of a boiler equipped with a leak detection system (black dots indicate sensor positions on the front side of the boiler, the superimposed squares indicate the corresponding positions of sensors on the back side).
rods externally welded to the component wall according to an appropriate welding procedure. Signal processing and diagnostic criteria for leak-induced sound and vibration Signal processing essentially consists of filtering and taking the signal root mean square (RMS). The basic diagnostic procedure is to compare the current RMS noise level to pre-set threshold levels, established during the system setup, when plant background noise is measured in different typical running conditions. A constant or very slowly increasing noise signal indicates a manageable leak, while a rapid increase with time suggests impending criticality (see the example in Fig. 2.4). Frequency analysis is also applied, but complex processing appears to add little to diagnostic performance and the simplest analysis level is in most cases to be preferred. The intelligent core of the system lies in the logical rules incorporated in
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2.4 Example of leak detection by means of acoustic monitoring. (a) The electrical load signal is plotted; the other signals are the output from the different acoustic sensors, in homogeneous arbitrary units. (b) The increase of signals and – to a lesser degree – (c) the increase of signals with time indicates a leak in the area monitored by the corresponding sensors (courtesy of CESI).
the software to recognize and rule out spurious noise sources (e.g. during sootblower operation intervals), changes in operation conditions and other potential sources of false alarms, and to assist plant personnel in the correct interpretation of system alarms. To confirm an active leak the operator’s attention is drawn to the signal time behaviour and to the possible convergent indications of neighbouring sensors.
2.2.2 Corrosion monitoring Boiler tubes undergo a large variety of corrosion phenomena, originating from either waterside or fireside corrosion. The former, common also to HRSG tubes, includes caustic corrosion, hydrogen damage, pitting and cleaning-induced damage; the latter includes coal and oil ash corrosion, lowtemperature flue gas corrosion as well as erosion (due to fly ash, coal particles, sootblower action). Combined effects such as corrosion fatigue failure must also be taken into account. Conventional and advanced corrosion-coupon probes The corrosion behaviour of materials of conventional boilers in standard operating conditions has been extensively characterized; it is well known and used normally in boiler design. Conventional corrosion coupons exposed in the boiler volume have been used for many years to characterize material behaviour in actual operating conditions. They provide largely qualitative time-integrated information on the comparative behaviour of
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different materials under the same operating conditions. The usefulness of cheap corrosion coupons greatly decreases when operating conditions vary substantially over relatively short times (e.g. variable fuel, variable combustion conditions). The safe implementation of process/design/material innovations in thermal boilers, besides requiring extensive preventive characterization of corrosion behaviour, may benefit from short-term feedback on actual corrosion behaviour of the installed materials, as operating conditions are not well known as often they are not sufficiently controlled. Semiquantitative information on the corrosion rate in actual operating conditions, although reliably collected over a relatively short operation period, can avoid potentially severe damage to the boiler, and thus large cost penalties. Different types of corrosion probes have been developed in past years, including advanced corrosion-coupon probes. These probes can be moved and controlled from outside the boiler, accurately positioned and temperature-controlled, to mimic at best the actual waterwall tube exposure and working conditions of the tube wall material in specific areas. Typically, the temperature of the coupon should be the same as the fireside surface temperature of the tube wall. The probe bears at its tip a small disc of the same material as the waterwall tubes, which can be easily extracted to measure time-integrated corrosion thinning or can be used for examination or replacement. These probes enable a reliable extrapolation of direct, relatively short-term (a few weeks) corrosion data to be obtained over a longer time range, and can also be used in co-combustion furnaces (Covino et al., 2005). Electrical resistance corrosion probes Other corrosion probes are based on electrical resistance measurements (DC or AC) that provide time-integrated estimates of corrosion thinning or, in more sophisticated electrochemical principles, that allow in principle a realtime estimate of the corrosion rate. Resistance probes exploit the increase in the electrical resistance of the conducting path due to corrosion-induced thinning. They can also be temperature controlled, with the same advantages described above, and can estimate the time-integrated corrosion thinning, though indirectly. A recent refinement of this concept is reported to enable wall integrity to be measured against corrosion over extended areas of boiler waterwalls, instead of a limited number of points. This system, already tested in the field, is based on a rectangular matrix of electrodes welded to the external, cold-side, tube surfaces, typically 1 m apart, and is minimally intrusive with respect to the boiler tube wall (Farrel and Robbins, 1998, 2002). Installations
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2.5 Example of boiler areas requiring corrosion mapping (left) and results (right) in terms of the corrosion rate (upper figure) and overall thinning (lower figure) (Farrel and Robbins, 2002; courtesy of Rowan Tech.).
of 100–200 sensors are reported, covering an area of about 100–200 m2 of waterwall. Pairs of electrodes in the matrix are scanned in a predetermined sequence to inject current and measure the equivalent circuit resistance, based on the same well-proven principle used in simple corrosion probes. The collected set of data is processed, taking into account the influence of current path geometry and temperature, to obtain maps of corrosion thinning over relatively large areas (see Fig. 2.5). The method is also reported to provide real-time estimates of fireside wall temperatures and temperature variations maps, as well as of heat flux, over the same area (Farrell et al., 2004; Robbins et al., 2004, 2009).
2.2.3 Overheating and heat exchange monitoring Motivations for monitoring temperature and heat exchange in boiler tubes A large portion of boiler-related power unit shutdowns are due to ruptures of high-pressure tubes in the steam generator, associated with some sort of temperature anomaly: stress rupture due to bulk overheating is ruled by the average temperature across the tube wall (determining its mechanical strength). Fireside corrosion is enhanced by local overheating and destructive damage takes place above a given fireside temperature threshold. Furthermore, thick oxide scales build up on the inner surface of waterwall
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Stress corrosion Corrosion due to coal ash Corrosion due to fuel oil ash Mismatching
Erosion due to ash Low-temperature corrosion
Erosion due to slag Erosion due to coal particles High cycle fatigue Thermal fatigue
SH and RH tubes SH and RH tubes SH and RH tubes Dissimilar welds in SH and RH tubes
RH tubes Economizer tubes
Hopper tubes Burners. Tubes near burners All tubes near anchors . Hopper tubes exposed to water jets . Tubes invaded by condensation due to blowing fluid . ECO and SH tubes subject to hightemperature variations . Zones with presence of corrosion phenomena and deformation due to cyclic constraints
Corrosion fatigue
Pitting corrosion Creep
Hydrogen embrittlement Flow accelerated corrosion
Waterwall LP tubes and headers in HRSG
All tubes SH and RH tubes
Caustic corrosion
Waterwall tubes SH tubes
Corrosion
Short-term overheating
Waterwall tubes SH and RH tubes
Tube sheets
Damage mechanism
Component
Fatigue in the presence of a corrosive environment
Metal temperature increasing due to: . Obstruction of the fluid flow due to thick internal oxide scale . Reduction of the fluid flow due to possible upstream leak . High heat fluxes due to irregular combustion Corrosion due to high concentration of NaOH in water in components characterized by high thermal flux and deposits (with the exception of stainless steel) Fractures due to service with low pH water due to the presence of salts Corrosion mechanism in which a normally protective oxide layer on a metal surface dissolves in a fast-flowing water or steam/water. The underlying metal corrodes to re-create the oxide, and thus the metal loss continues . High concentration of S and/or chlorides in the fuel . Misaligned burners . Formation of liquid slag on the external surface of tubes . Low oxygen combustion Oxygen and/or salts of high concentration in water Operation at temperatures close or higher to the maximum design temperature due to: . Excessive scale thickness . Insufficient circulation of the coolant . Excessive temperature of combustion gases Cracks generated by corrosion (high concentration of chlorides, sulphates or hydroxides) under high stresses Coal ash with corrosive slag and metal temperature between 600 8C and 700 8C Fuel oil ash with corrosive slag and metal temperature between 600 8C and 700 8C Cracks in the welds between heterogeneous metals due to temperatures and/or constraints of operation higher than those estimated, due to installation conditions Regions with higher fluid speed . Tube operating temperature lower than condensation temperature with formation of acid condensates on the tube external surface . Gas operating temperature lower than condensation temperature with formation of acid condensates on the ash particle surface Erosion due to slag fall on sloping surfaces Erosion of coal particles in the burners and on the tube surface due to the high velocity of the combustion air Tube vibrations induced by the energy of combustion gases or by their turbulence Fatigue caused by repeated changes of temperature
Cause of damage
Typical causes of damage to tubes in boilers and HRSGs
Table 2.2
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tubes and can produce both a decay in heat exchange efficiency and variations in temperature gradients across the tube wall (see Table 2.2). An accurate and reliable knowledge of the temperatures in different areas of the boiler helps to detect potentially dangerous local anomalies, preventing failures and ensuring the long-term integrity and functionality of the boiler. An approximate knowledge of the spatial distribution and efficiency of heat exchange between the combustion gases and working fluid is important during both transient and steady operation. This information is particularly relevant during test runs, first at startup and initial operation of boilers of non-conventional design (e.g. SH steam temperatures > 600 8C, new materials) and/or with major process changes (e.g. co-combustion). Design and/or process modifications have to be thoroughly tested in these boilers and running conditions consolidated. Operational experience can later point out other problems that may arise in individual power units during normal operation and that may require and justify the implementation of targeted temperature/heat flow monitoring within the boiler at any time. Infrared methods for temperature monitoring Thermal imaging cameras, working in the infrared (IR) range and specifically adapted to the harsh (because of temperature and environment) and complex (because of the need to exclude or separate — among others – the influence of the hot combustion gases on the thermal image) operating conditions of a running boiler, provide a quick, efficient and reliable means of detecting thermal anomalies in different areas of a boiler. Uneven thermal exchange due, for example to sootblower malfunction, hot spots, fluid flow blockages or overheated areas can be promptly identified and then monitored to check the effectiveness of corrective action (Lyon, 2000). Being a non-local (wide-area), remote monitoring method, thermal imaging is effectively used to compare temperature distributions in different areas or temperature differences within a given area. Use of permanently inserted thermocouples to monitor temperature and heat exchange in waterwall tubes If attention focuses on quantitative, long-term temperature and temperature profile measurement across the tube wall, several solutions based either on externally fixed thermocouples (TCs) or on TCs embedded in the tube wall are also commercially available. A proven embodiment of solutions based on embedded TCs is described in the following. Two grooves are spark-eroded inside a radial section of a waterwall tube portion; each groove hosts one TC (Fig. 2.6). The tip of each
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2.6 Monitoring temperature profile inside a boiler tube wall: (a) schematic diagram for the positioning of the thermocouples inside the tube wall; (b) section of a boiler tube showing the actual position of the thermocouples (courtesy of CESI).
TC lies at a known depth inside the tube wall. Both TCs lie in the same radial plane, their tips aligned along the same radius, normal to the fireside surface; embedded TCs are not exposed to the aggressive hot gas environment, thereby ensuring long-term survival and performance. The grooves are machined starting from the fire-shielded fin area, so that the TCs can be safely routed to data conditioning and acquisition modules outside the boiler and can be easily replaced, if required. A third TC in the unexposed portion of the tube wall measures the temperature of the process
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2.6 (continued)
fluid. In this way a fairly accurate temperature profile across the wall thickness, where the heat flow is maximum, is reconstructed and made available online. In particular the average wall temperature and the temperature at the outer fireside surface of the tube can be obtained, together with a qualitative appreciation of heat flow across the tube. In the actual implementation a tube section (~ 70 cm) was cut out of spare tubes, processed as above, equipped with TCs and extension cables and inserted by welding in a waterwall panel in place of an equivalent length of tube. Laboratory tests showed that the mechanical strength of the tube inserts was not impaired. Typically, from 10 to 20 instrumented tube inserts are installed in critical areas of the boiler to provide an overall picture of the ongoing heat transfer process: this is a necessary compromise between cost and coverage and usually proves to be adequate for diagnostic purposes, if the positioning of the tube inserts is judiciously selected according to engineering analysis and previous operating experience. Manufacturing and installation of these instrumented inserts is relatively expensive, if compared to externally applied thermocouples. On the other hand, this solution provides the temperature profile across the tube wall and has a useful operating life similar to that of the average tube life. Several boilers in different power units have been equipped with these instrumented inserts in the last three decades.
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Use of permanently inserted thermocouples to monitor temperature and heat exchange in SH and RH tubes Similar considerations hold true for SH and RH tubes. The average wall temperature of the hottest tube sections is in most cases the life-determining factor of SH and RH tubes: they mainly undergo high-temperature creep damage and overheating. Continuous monitoring of the temperature of SH and RH tube walls enables direct and accurate evaluation of creep life consumption and a true condition-based life management of these equipments. Unpredicted severe damage to SH and RH coils is more likely to require more lengthy shutdowns for replacement and considerably more cost penalties than incurred from damage to waterwall tubes. Due to the position and configuration of SH and RH coils, thermal imaging is not an easy exercise. Since surface-mounted TCs are generally short-lived and work as short-term solutions, embedded TCs can be used. In this case the average wall temperature at different radial sections along the tube length is of interest. As only one TC per radial section is required, two to four TCs are embedded in a longitudinal groove (typically 1 mm deep), machined on the outer surface of a straight portion of an SH or RH tube. The groove containing the TCs is sealed by welding a longitudinal metal strip to the underlying tube. TCs are thus completely shielded from contact with hot flue gases. Each TC measures the temperature at a different length along the tube, so that an accurate, reliable, long-term temperature profile of the tube is provided. This technology is applicable only to the terminal (inlet or outlet) portion of an SH or RH coil. At their exit from the upper end of the tube the TC cables are routed through the penthouse to the conditioning and data acquisition unit. One instrumented tube for each SH or RH coil is usually sufficient to obtain the required information, taking due consideration of the relative complexity of installation and relevant cost of each instrumented tube. Advanced fibre optic methods for temperature monitoring inside the boiler Thermocouples have well-known limitations in maximum measurable temperature, response, accuracy, stability and service life in harsh environments. If only surface temperature measurements are required, optical fibre thermometers (OFTs) (Yu and Chow, 2009) can be used instead, with the advantages of long-term stability, high sensitivity and quick response. Moreover, they are unaffected by electromagnetic interference, and can survive harsh environmental conditions. In the family of OFTs, blackbody and fluoroscopic sensors are widely used. Blackbody sensors consist of a high-temperature optical fibre with an opaque cavity
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attached to the sensing tip. The radiation flux spectrum detected at the end of the fibre is related to the temperature of the cavity via Planck’s law, which permits the temperature to be obtained from the measured spectral intensity or intensity distribution. This technique is particularly suitable at high temperatures (500–1900 8C), while at low temperatures fluorescent OFTs have better capabilities. Both types of OFT have high accuracy (<0.5 8C) in their effective measurement ranges and have the capability of remote measurement. The small size of the fibre and its electrical, chemical and thermal inertness make long-term installation of the sensor inside complex equipment possible; thereby access is provided to locations that are difficult to address. These techniques cover an unusually wide range of temperatures with adequate precision; they are also simple to calibrate or need no calibration of individual probes, as in the fluorescent lifetime-based OFT. In recent years intelligent monitoring systems utilizing distributed fibre optical sensors for real-time monitoring of high temperatures in a boiler furnace have been developed (Lee et al., 2004). Also special sensors based on the CO2 emission spectrum are nowadays on the market and are used to measure the gas temperature inside the boiler (e.g. see www.infra-view.com/ whitepaper.htm).
2.2.4 Structural integrity monitoring of headers and steam lines The main damage mechanisms acting on headers in boilers and HRSGs are reported in Table 2.2): creep in welds of high-temperature headers, thermomechanical fatigue cracks in header-to-stub or pipe welds and on the internal surface of header penetrations. The main damage mechanisms of steam lines are also reported in the same table: creep in welds of hightemperature steam lines, thermomechanical fatigue due to stress concentrations in welds or in connections/fittings. On-line monitoring of the integrity of high-temperature, high-pressure steam headers and steam lines is of primary interest mainly for the early detection of active cracks and leaks, which could finally lead to gross – and potentially dangerous – failure. The task is difficult due to the hostile environment and to the lack of access, even to an effective visual inspection, due to thermal insulation. Many attempts to develop effective monitoring tools (mechanical, optical, ultrasonic sensors to detect permanent deformation; high-temperature fibre optic sensors; see Carolan et al., 1997) have been pursued in past years for such high-temperature components, with variable claims of success. Although most methods do not reach a
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widespread acceptance, due to the lack of sensitivity and/or reliability, some examples are reported in the following. For integrity assessment and leak detection the ALD systems discussed in section 2.2.1 can be used, especially when limited volumes are concerned (e.g. steam headers in the penthouse). Monitoring the integrity of steam headers and steam lines by acoustic emission methods Another acoustic technique, acoustic emission (AE), is relevant to this purpose, though less widespread. It is based on the detection and processing of transient elastic waves generated by a sudden energy release from a localized source within a stressed material (e.g. caused by brittle crack propagation). Microvibration sensors are used, with a resonant frequency selected to high-pass the noise generated by the steam flow. Elastic waves travel a long distance along the structure and, by using an array of several sensors, this method can in principle keep the whole component under control in real time and localize as well the position of the elastic wave source (the area of potential damage). Typically, a set of strategically positioned AE sensors is coupled to the steam line, through thin (~4–6 mm) metal rods welded to the pipe wall (an appropriate welding procedure is necessary) (see Fig. 2.7). The stress-wave signals (high-frequency vibration signals typically in the 0.1–1.0 MHz range) are detected by the sensors, conditioned, processed and digitized by local instrumentation modules and then collected by a central processing and display unit, where further processing and display of diagnostically significant results takes place. To obtain information on source location, signals from different sensor channels are compared on a common time scale, so processing is time-correlated. Each signal is also processed independently; parameters like signal counting, signal amplitude distribution and frequency spectrum are used to discriminate between significant and non-significant, critical and non-critical sources. A sensitivity calibration must be carried out as a preliminary step, in order to get meaningful results. The source location algorithms, when used, must also be tested for proper operation, by using different types of artificial sources. Diagnostically significant and useful results were claimed by many authors, concerning for instance the identification and on-line assessment of damaged, potentially dangerous areas in SH and RH headers (Morgan and Foster, 1995; Cattaneo et al., 1998) and steam lines (Evans et al., 1988; Rodgers et al., 1996; Muravin et al., 2004; Rodgers, 2007). For example, the technique was used by EPRI to follow the growth of cracks in the thick wall of a header in the region of a tube weld (Morgan and Tilley, 1999); in this case a correlation between the AE rate and the stress peaks during thermal
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2.7 Structural monitoring of a steam header by means of acoustic emission: sensor coupling to the header wall.
transients was demonstrated. AE application guidelines have also been issued (Rodgers and Tilley, 2004). The method offers substantial perspective advantages in terms of an improvement in operation safety and optimized asset management at affordable costs (typically extending the safe operation life of an old steam line until the most convenient time for replacement). Nevertheless, many factors still appear to negatively affect its diagnostic reliability and accuracy. Hence the use of AE to monitor crack nucleation and propagation in steam lines appears to be still controversial and not yet widely accepted. Monitoring creep damage in steam headers and steam lines To make a reliable residual life evaluation of high-temperature components subjected to creep it would be necessary to measure creep strain and possibly also the creep strain rate. Monitoring creep strain in high-pressure steam pipes and other power plant components is a complex issue, due to the very demanding environmental and operational conditions that sensors must survive. Moreover, measurements must be able to single out creep deformation. Capacitive strain gauge sensors can be used (e.g. Baumann and Schulz, 1991) on pipes and dissimilar weldments for residual life
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evaluation. Some current state-of-the-art techniques in high-temperature strain measurement and extensive guidelines can be found in the ASTM E1319-98 (2009) standard, covering measurement of static strain in the temperature range from 425 to 650 8C (800 to 1200 8F). An optical technique, digital speckle photogrammetry (DSP), has also been proposed and has been under development since the 1980s (Kim, 1989). It is based on the detection and measurement of the interference fringes that appear when images – taken at different times – of an artificial grid permanently applied at a carefully selected location of the pipe surface are superimposed. Recently, sensors with a rugged design have been developed and installed across sections of welded steam pipe and other plant components in locations that provide the best monitoring points to detect early the onset of failure processes (Morris et al., 2006). Precision optics and a charge-coupled device (CCD) camera are used for uniaxial and biaxial strain measurement (Morris et al., 2007a, 2007b). Optical creep strain measurements are reported to have a 65 microstrain accuracy level with an error of less than 10 % (Morris et al., 2009). Data obtained during maintenance shut down periods can be used to assess the remaining life of steam pipes and other components operating in creep conditions. Nevertheless, the method is not widely used in practice because it requires that only creep strain contributes to the deformation of a component surface from one specklegram to another, which is actually fairly difficult to achieve in a real plant environment.
2.2.5 Assessment by in situ non-destructive techniques This section deals mostly with conventional condition assessment methods that can be applied only during boiler shutdowns and consequently are not amenable to continuous, on-line surveillance. They are certainly simpler and cheaper than all the on-line diagnostic methods discussed in previous sections and their results are more easily verified. They cannot answer, however, the ever-increasing need for higher reliability, availability and operational flexibility of thermal boilers unless they are combined with costeffective on-line condition monitoring methods. Table 2.3 summarizes some of the NDT methods (in addition to basic visual inspection) most commonly used to assess the condition of the different components of the boiler during maintenance shutdowns. Inspection of the boiler As accurate visual inspection of boiler furnace walls is a relatively lengthy and expensive task, attempts have been made to carry out visual inspection
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Table 2.3 Non-destructive testing (NDT) methods used on boiler and HRSG components Components Economizer headers Waterwalls Boilers drums Lower waterwalls and headers Junction headers Waterwall risers Waterwall headers Superheater headers (welds) Reheater headers (welds) Liners Nozzles HT superheater tubing Steam piping Feedwater piping
DC
TM
X X X
X
ME
HT
X X
X X X X X X X
PT/MT
X X X
X X X
X
X X X
X X
X X
ET
UT
RT
X X
X X X X X X X X X X
X X X X X X X X
X X
Acronyms used for non-destructive testing methods: DC = dimensional checks; TM = thickness measurement; ME = microscopic examination; HT = hardness testing; PT = dye penetrant testing; MT = magnetic particle testing; ET = eddy current testing; UT = ultrasonic testing; RT = radiographic testing.
of boiler walls by means of remotely operated robots carrying TV cameras, but the method is complex and not widely applied in practice. Mapping of residual tube thickness is mostly carried out by conventional ultrasonic methods that require scaffolding to permit direct access of personnel to the boiler walls. Accurate thickness mapping of boiler waterwall tubes enables significant comparisons to be made with maps taken in previous inspections and timely appraisals of the tube thinning rate and localization, as a function of operation history. This in turn makes it possible to take timely corrective or remedial action. Also, non-contact electromagnetic-ultrasonic transducers requiring minimal or no surface preparation have been demonstrated to be able to map residual tube thickness because of their use with remotely operated robots (Bozzetti et al., 2003). Methods for assessing inner oxide thickness in waterwall tubes Waterwall tube samples can be periodically extracted to assess the material condition across the wall and the condition of the inner surface: measurement of the oxide layer thickness at the tube inner surface has to be done very precisely on tube sections, after careful metallographic preparation, so as to avoid damaging the oxide layer. The oxide thickness determines the need and the time schedule of chemical cleaning.
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Several techniques for non-destructive evaluation of inner oxide thickness have been developed and tested, which are similar to standard ultrasonic wall thickness measurements. The reflections of an externally injected ultrasonic pulse from the metal–oxide and the oxide–air interfaces are detected and the time-of-flight of the various reflected pulses are measured and processed to yield wall and oxide thicknesses (e.g. see the GE web site). Inspection of SH and RH tubes The wall thicknesses of SH and RH tubes are mapped by means of ultrasonic thickness measurements on a discrete set of measuring points, by methods similar to those adopted for waterwall tubes. Dye penetrant and magnetic particle (for low-alloyed steels) methods are used to check the integrity of welds between tube ends and header inlet/outlet penetrations. A few tube samples can be extracted from the most critical locations to perform metallographic examination and mechanical tests in the laboratory in order to evaluate material creep life consumption. Inspection of steam headers and steam lines The inspection of steam headers and steam lines is performed by visual, borescope and ultrasonic methods to detect cracks, mostly due to thermal fatigue and originating from the inside surface of header nozzles and penetrations. Dye penetrant and magnetic particles (for low-alloy steels) are used to detect cracks, due to creep and/or thermal fatigue on the outer surface of header and steam line welds and penetrations. Ultrasonic methods are used to assess the through-thickness integrity of assembly welds in headers and steam lines and approximately to size the length, and above all the depth or height, of surface breaking or internal cracks, if any, in steam line welds and in header assembly welds and nozzles. Due to the large number of welds to be inspected on SH and RH steam lines and other high-energy lines, each inspection campaign is often limited to a selected number of representative welds, different for each campaign and chosen so that all welds are controlled after three or four shutdown maintenance inspections. Metallographic replicas are taken in functionally critical locations to assess material conditions in terms of creep damage in header and steam line welds. In situ metallography with portable microscopes is also used for the same purpose.
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Methods for assessing material condition in serviced steam headers and steam lines Material samples can also be eventually extracted from structurally noncritical portions of headers and steam lines to carry out mechanical tests in the laboratory, in order to assess current actual fracture toughness and creep life consumption of the material (Parker et al., 1995). A quasi non-invasive method, the so-called ‘small punch’ (SP) method, has been developed in past years to provide an insight into the mechanical properties of serviced materials in structural components, without implying damage or destruction. The method is more important for the assessment of fitness for further service and life consumption of turbine rotors than for headers, steam drums and steam lines, where it is relatively easy to access substantial amounts of material for conventional testing within the frame of repair and replacement actions. For these reasons the SP method will be discussed in more detail in section 2.3.5.
2.3
Steam turbines and generators
In thermal power plants the digital control system (DCS) has a vital and well-established role in ensuring long-term integrity, reliability and availability for the steam turbine, gas turbine, generator and major auxiliaries. Temperature and pressure time history is recorded systematically at crucial locations of steam turbines during startup, shutdown, trip and other operation transients, as well as during steady state operation. This information is extremely useful not only for day-by-day power unit management but also for performance monitoring and life assessment (Beebe, 2003).
2.3.1 On-line vibration and dynamic condition monitoring Vibration monitoring is the core of dynamic control for the main rotating machinery. It is the primary on-line, real-time source of diagnostic information on the actual dynamic condition of the machine (Doebling et al., 1996). This information is integrated by other parameters, ranging from current unit load to oil-film pressure, to help the diagnostic interpretation of vibration data. Nowadays vibration monitoring systems of high complexity and performance are able to deliver exhaustive, real-time diagnostic information to the power plant staff on the actual dynamic condition and performance of rotating machines. Standard high-performance vibration monitoring systems consist of pairs of accelerometers mounted on each bearing of the turbine shaft, at 908 from each other (see Fig. 2.8). An additional sensor (‘key phasor’) is also required
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to provide a common time reference for a correlated analysis of signals from the different sensors. The system necessarily includes on-site signal conditioning modules as well as a data acquisition and processing subsystem. Such a system can provide a real-time frequency analysis of turbine vibration and in principle enables a relatively sophisticated diagnostic appraisal of the machine dynamic condition. On-line dynamic monitoring is also being used in many instances to help power plant personnel during the crucial phase of startup of the large rotating machinery after planned or unplanned shutdown. The details of the dynamic behaviour of the machine when the critical rotor speeds are approached and overcome can be used by the experts to achieve a better understanding and control of machine behaviour and to go through the startup phase more safely. A very simple example is reported in Figs 2.9 and 2.10 (Lapini et al., 2005). It shows a cold run-up (with an important amplitude increase while going through critical speed). Additional reference data provide information needed to decide at any time whether the increase of amplitude is tolerable (a natural resonance frequency) or whether there may be some problem. The usefulness of an advanced monitoring system – not only for diagnostic purposes but also for the management of critical conditions or critical machines – has been demonstrated in many instances. DCS presentations are sufficient during normal operation, but a customized data display can be very helpful when anomalies or problems arise. An anomalous increase of vibration amplitude during transients can reveal the presence of growing defects in the rotor. An interesting example of the capability of vibration monitoring to detect the presence of a crack in a generator rotor can be found in Bicego et al., 1999. The vibration monitoring system revealed a gradual increase in vibration amplitude of a 370 MVA generator rotor after about 133 000 operating hours and about 400 startups; the plant was shut down and the rotor scrapped because further analyses demonstrated the presence of a deep fatigue crack (60 % of the rotor transverse section). Several modelling studies analyse the transient vibration response of a cracked rotor passing through its critical speed in an attempt to deliver results useful for crack detection and monitoring. A simple hinge model was proposed, for example, for small cracks under the breathing action induced by high-speed rotation (Sekhara and Prabhu, 1998); the effects of such factors as crack depth and unbalance eccentricity on vibration behaviour are also modelled in this work. The in-depth analysis of the huge amount of information contained in the time history of vibration signals from each sensor installed on a turbine, together with the time-correlated analysis of the signals from different sensors, can substantially improve the diagnostic performance of vibration monitoring. Both earlier detection of potential damage and better diagnostic
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2.8 Two couples of velocity and acceleration sensors radially mounted at 458 in correspondence to each bearing for vibration monitoring of large turbine shafts.
2.9 Rotor vibration amplitude versus rotation speed at a given bearing, in the presence and absence of a developing fault, during startup of a turbine-generator set (courtesy of CESI).
reliability and precision can be achieved, well beyond that obtained with systematic frequency analysis. Results are further improved if the analysis of collected data is supported by a good model of the dynamic behaviour of the investigated machine (Gregori et al., 2000; Bachschmid et al., 2002). This approach, beneficial in many instances, requires a level of specialized
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2.10 Vibration orbits along the shaft line at operation speed (3000 rpm), in the faulty conditions of Fig. 2.9 (courtesy of CESI).
manpower commitment that possibly conflicts with financial constraints. Clever and efficient automated strategies for data analysis and alarm generation can come into play, including: . . . . . .
algorithms for automated data processing and reduction; algorithms for the automated identification and periodic assessment of ‘normal’ baseline behaviour; algorithms for the identification of significant deviations from ‘normal’ behaviour; logical criteria for alarm identification and confirmation; where feasible, model- and/or knowledge-based support for malfunction or damage screening and identification; provisions for storing and making available all data associated with a given alarm, which enable the expert to carry out an independent analysis as completely as possible, since reliance on automated diagnosis only should be avoided.
These ‘smart alarms’ make it possible to carry out top-level on-line condition monitoring of critical items with a reduced expert manpower commitment (Lapini et al., 2005) and have been implemented – at different
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complexity levels – in machine monitoring packages available on the market (e.g. Killich, 2006, and www.ge-mcs.com). It must be stressed here that both standard and advanced on-line dynamic condition monitoring of main PP machinery have to rely on an efficient, flexible, well-organized and safe data exchange infrastructure between the machine site, the plant control room and remote expert locations, thus making it possible to exploit in real time – when necessary – high-level expertise and consultancy (see also section 2.6). In recent years vibration monitoring has contributed to the prevention of accidents by giving reliable warning of thermomechanical fatigue cracks in rotors of combined cycle (CC) power plants, initially designed for base load operation but actually severely cycled, as required by power market needs.
2.3.2 Damage mechanisms Steam turbine rotors are among the most critical and highly stressed components of steam power plants (see Table 2.4). In most cases, mainly for cyclic operation, thermomechanical fatigue damage can be more important than creep in determining the life consumption rate of most large steam turbines, so attention needs to be concentrated on transients. Temperature/ pressure data and time history are the basis for integrity assessment. If an FE thermomechanical (possibly simplified) model of the turbine is available, the recorded thermal transients can be used to derive the time-dependent temperature–stress distributions in the most critical positions along the rotor and the other turbine components. A wealth of algorithms, many of them implemented into proprietary or commercial SW, are available to convert the detailed features of the time-dependent temperature–stress distributions into a few simple parameters, which can be useful when evaluating cumulative life consumption due to thermal fatigue according to the acknowledged life-consumption models. Such information is also the basis for a safe life extension of steam turbines (where and if possible). For the sake of completeness, the prevailing damage mechanisms observed in other steam turbine components are summarized in Tables 2.4 to 2.7. Steam turbine casing damage generally consists of distortion and/or cracking. Cracking of casing can lead to steam leaks, whereas distortion can cause further damage due to contact between the stationary and rotating parts. Cracks in casing are typically located at inlets of the HP and IP sections, where the local thermal stresses are higher. Cracks may also be observed, although to a smaller degree, at the inlet section of LP casing. Generally the causes of cracking in the HP and IP casings are low-cycle/thermal fatigue occurring during startups and shutdowns (65 %), brittle fracture (30 %) and
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Table 2.4 Main damage locations and causes in steam turbine rotors Component Position
Cause of damage
HP-IP rotor
Fatigue, creep Skin peeling (if possible) Fatigue, creep Overboring crack propagation, brittle fracture Fatigue, creep —
LP rotor
Outer groove Centre bore
Blade-groove shoulder Centre bore
Remedial action
Fatigue and Overboring brittle fracture
creep (5 %). Brittle cracks occur during transients in casing with inherently low toughness or damaged material due to service-induced temper embrittlement and are the most dangerous, requiring immediate remedial action. The primary forces acting on blades are centrifugal tensile forces due to rotation and bending forces due to steam flow. In addition to steady forces, the blades are also submitted to non-uniform steam flow, causing fatigue phenomena. First-stage blades are also exposed to complex forces during partial steam admission and to occasional over-speed and fatigue stresses due to start–stop transients. In LP blades problems arise primarily from corrosion fatigue due to impurities condensing from the steam and water droplet erosion. Generally the main causes of blade damage are: stresscorrosion cracking (22 %), high-cycle fatigue (20 %), corrosion-fatigue cracking (7 %), creep rupture (6 %), low-cycle fatigue (5 %), corrosion (4 %), unknown (26 %), other causes (10 %). The main damage mechanisms of diaphragms and nozzle boxes are creep and thermal distortion of the nozzle chambers and cracking of the nozzle roots due to creep and thermal fatigue. Diaphragms and nozzle boxes are stationary components, and hence tolerant to cracks. Cracks can be weld repaired or the nozzle can be retired, when repair is no longer possible or convenient, or when distortions become too severe. Another service problem related to the first few stages of nozzles and blades in the HP and IP sections is that they can also undergo solid particle erosion, mainly by hard magnetite particles from boiler tubes. The main causes of damage in valves and steam chest are low-cycle fatigue to thermal gradients during transients, creep during steady state conditions at full load, erosion due to steam-borne particles and wear as they are key pressure vessel components operating at high temperature and pressure, their failure involves both safety and outage concerns.
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Condition monitoring and assessment of power plant components Table 2.5 Component
Damage locations and causes in steam turbine blades Position
Cause of damage
HP-IP blades Tenon and root Creep (few first of blades stages)
Tenon and root High cycle of blades fatigue
Airfoil of blades Solid particle (first few erosion stages) Shroud bands and tenons Root of blades
LP blades
67
Remedial action Tenon: grinding of tenon above the blade tip and weld repair Root of blades: replacement* Tenon: grinding of tenon above the blade tip and weld repair Root of blades: replacement* Airfoil of blades: blending out and, if necessary, weld repair of affected zones of the blade Shroud bands and tenons: weld repair Root of blades: replacement*
High stress concentration Airfoil of blades Corrosion Airfoil of blades: grinding and (few last fatigue weld repair** stages) and Root of blades: replacement* root of blades Blade edges Erosion/ Weld repair of affected zones corrosion of the blade (water Replacement of damaged droplet) blades by new ones and of new design Application of protective coatings to guard against corrosion and erosion damage Tie wires Stress corrosion Tie wires: full penetration weld Tie wire holes cracking repair, split sleeve repair, Blade covers plug insertion Tenon holes Tie wire holes: grinding*** Blade edges Blade covers: weld repair Root of blades Tenon holes: grinding and weld repair*** Blade edges: weld repair** Root of blades: replacement* Root of blades High stress Root of blades: replacement* concentration
* At this time, repair of the roots of rotating blades is not recommended because of the high stresses inherent in this attachment area. ** At this time, the damaged blades are usually replaced, as repairing is difficult *** If the damage zone is small
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Table 2.6 Damage locations and causes in steam turbine diaphragms and nozzle boxes Component
Position
Diaphragms and nozzles Chambers and roots Chambers and vanes Airfoil of blades (first few stages)
Cause of damage
Remedial action
Fatigue, Weld repairing * creep Thermal Replacement (in some distortion cases re-machining **) Solid particle Replacement erosion
* The operating life can be extended by weld repairing; the total life is spent when weld repairing is no longer applicable. ** The operating life can be extended by realignment and re-machining; the total life is spent when such actions are considered to be no longer realistic. Table 2.7 Locations and causes of damage in steam turbine valves and steam chests Component
Position
Valves and steam chest Casing Valve seat Stems, bushings and seal rings
Cause of damage Fatigue, creep Erosion Wear
Remedial action Weld repairing * Replacement Replacement
* The operating life can be extended by weld repairing; the total life is spent when weld repairing is no longer applicable.
2.3.3 Non-destructive examination As many of the degradation modes previously mentioned can only be detected by visual inspection or non-destructive testing, casings do require to be opened for inspection. This is particularly true for the many large power generation machines continuing in service beyond their intended design life. The variety of defects to be potentially found in serviced steam turbines consists mainly in surface breaking defects on relatively easily accessible surfaces, to be detected by means of visual, dye penetrant and magnetic particle methods, or by replicas or in situ metallography, where material degradation is implied. Bolt inspection against fatigue cracks may require ultrasonic examination. However, the main technical problem is to assess the structural integrity and reliability of turbine rotors. This holds true for both hollow rotors of
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older design and new welded rotors. In both cases accurate inspection during planned maintenance shutdowns is essential, since on-line methods are not very effective in early detection of such service-induced defects as fatigue or thermal fatigue cracks in the rotor shaft, possibly originating from pre-existing manufacturing defects. Ultrasonic inspection is in fact the main tool, even if not the only one, to address rotor integrity problems. The integrity of older hollow turbine rotors is currently being investigated by means of advanced boresonic examination techniques and equipments (Bontempi et al., 2005). In fact, a higher sensitivity to the most frequent defects, those originating from near the inside surface of the rotor, is achieved by inspecting the rotor from inside the bore. For this purpose the inside surface must be first cleaned and conditioned by a honing treatment; then a motorized ultrasonic probe scans helicoidally the inside surface of the rotor, while reliable acoustic coupling is ensured between the probe and surface. Automated probe scanning, data acquisition and processing, in order to provide some form of ultrasonic imaging of the scanned volume, is usually required, since the time window available for inspection and diagnostics during shutdown is extremely limited, and probe scanning can be carried out only once. The ultrasonic inspection from inside the rotor is complemented by eddycurrent scanning of the same inside surface, to detect fine and small surfacebreaking cracks. Hardness measurements and replicas in selected rotor areas exposed during operation to very high temperatures are carried out in order to check potential material damage. Specific manual or mechanically assisted ultrasonic procedures are also used to inspect rotor discs and blade attachments for corrosion-fatigue cracks. In this context, where probe manipulation can be difficult, ultrasonic beam scanning phased-array probes can be used with good results. In the last decade phased-array ultrasonic (PAU) techniques have been increasingly used and have gradually gained industrial acceptance in several applications. A linear array of piezoelectric elements excited in sequence with fixed time delays generates an ultrasonic beam, whose orientation depends on element spacing and time delay. By varying in steps the time delay between the excitation of the individual elements of the array, according to the Huyghens principle (in similiar to the mode of operation of the radar), an ultrasonic beam of variable inclination is obtained, which dynamically scans a portion of the underlying area in a fan-like mode. This principle has been known and tested for a long time, but has only recently become a widely applied condition assessment tool, by exploiting simultaneous progress in sensor technology, electronics, data and imaging processing, together with modelling. The basic advantage is that the electronically steered ultrasonic beam requires no probe manipulation to
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span a section of the underlying volume, which can therefore be investigated even from a hardly accessible area, where the probe can be positioned but not manipulated. Furthermore, electronic scanning is more controllable and precise than manual scanning and is easily amenable to presentation of results in the form of an ultrasonic image, which is easier to interpret. This technique has interesting advantages, such as the reduction in inspection time, resulting in cost savings as well as the reduction in the logistics with respect to conventional methods; it is applied in the examination of thick components and welds. The PAU technique was also applied to the inspection of turbine discs at blade attachments and of bores of turbine discs (Abbasi and Fair, 2006). Developments in PAU techniques were also devoted to achieve detection of defects in the most critical region (the two serrations) of large last-stage blades with curved axial entry fir tree roots while in situ mounted on the low-pressure steam turbine rotor (Charlesworth, 2010).
2.3.4 Assessment of the structural integrity of main steam turbine components A reliable condition assessment of a steam turbine component (mainly rotor, but also casing, nozzles, etc.) requires simultaneous knowledge of: . .
.
the position, orientation and size of defects lying in the component; the mechanical properties of the material in the most critical areas (in terms of stress and/or flaws), i.e. the local load bearing capability (yield strength σYS, ultimate tensile strength σUTS) and fracture resistance (fracture appearance transition temperature FATT50, toughness KIC); the effective stresses acting on the existing defects during both steady state and transient operation.
The above three factors – defects, mechanical properties and stress distribution and history – should be accurately known for each individual component. Given the material, the mechanical properties of individual rotors may show a substantial spread with respect to standard reference values. The different weights of material ageing processes due to the specific operation history can further affect these properties. Moreover, the operation history, in both steady state and – more critically – transient conditions, determines the actual space and time distribution of stress. Transient conditions, in particular, are characterized by a large variability and high thermal and mechanical gradients. The more accurately these factors are known, the more reliable is the assessment of fitness for further service and residual life of the component. In practice, however, assumptions and simplifications are introduced, both
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to make up for the lack of knowledge of some parameters and to reduce manpower commitment and cost of the assessment. In conventional approaches these assumptions are very conservative: the most severe operating conditions are considered; lower bound mechanical properties are used; defect orientation is assumed to be orthogonal to stress. This may lead to a substantial underestimation of the residual life and to the rejection of components that could still operate under appropriate conditions. Developments in FFS and life assessment methodologies for ageing rotors (as well as for other major power plant components) have therefore been mainly aimed at: . .
.
determining with higher accuracy the specific input data (defect position and size, material properties, relevant operation history) of the individual rotor to be examined; improving the methods to calculate temperature/stress distributions and to evaluate fitness for service, so as to properly take into account all relevant parameters and to enhance the reliability of results, while keeping processing complexity at a manageable level; pursuing the above improvements while simultaneously reducing assessment duration and cost as compared to more conventional practices.
2.3.5 Miniaturized tests for rotor material assessment Determining the actual material properties, above all mechanical strength (yield and tensile strength) and toughness (FATT and KIC) in serviced critical components is essential to ensure their fitness for safe further service. This is particularly true for steam turbine rotors and attachments. The fracture toughness of serviced material needs in particular to be correctly estimated, in order to determine the maximum acceptable defect size for safe operation according to fracture mechanics criteria, to be compared with the results of the NDE of the rotor and discs. It can be questionable, however, if not impossible, to extract from the rotor the significant amount of material necessary to carry out conventional mechanical tests. In the most favourable case, 10–15 mm thick rings of material are directly extracted from the rotor in two different positions (‘hot’ and ‘cold’), from which standard test specimens are worked out. Heavy equipment and several days are required to extract such large samples from the rotor. Moreover, there are heavy constraints on the positions in which samples can be extracted and very often the allowed positions are not the most significant ones. An attractive alternative of growing importance to the power industry is offered by the use of miniaturized specimens. Among methods based on miniature specimens, the ‘small punch’ (SP) method has been proposed
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(Mao and Takahaski, 1987; Takahaski, et al., 1988) and used to obtain accurate estimates of the tensile and fracture resistance parameters of serviced rotor materials. The SP method relies on the extraction, in a minimally invasive way, of very tiny portions of material from the component; neither post-sampling repair nor derated operation of the component are required. The non-invasivity of the method makes it possible to select the most significant positions for material sampling. Sample extraction is carried out by means of specially designed miniaturized electromechanical devices, based on cutting or spark erosion. It is essential that the extraction process does not alter locally the mechanical properties of the extracted sample. The cutting equipment is portable and easy to position and use in field conditions. A typical in situ sample extraction takes 2–3 days and consists of three samples 1208 apart on the ‘cold’ and ‘hot’ rotor sections. A small disc-shaped portion (typically 20 mm in diameter, 3– 4 mm thick) of material is extracted by soft methods (e.g. spark erosion) in the relevant positions of the component (Fig. 2.11). About six to eight SP cylindrical test specimens 8 mm in diameter and 0.5 mm thick are machined out of each spherical segment of sampled material and submitted to a sort of punch test, with a semi-cylindrical head (1 mm in diameter). Deformation under increasing load is recorded and can be correlated with the main mechanical properties of the bulk material in the given location, namely to fracture toughness. The area below the load–deformation diagram, up to maximum load, is the fracture energy ESP. SP tests carried out at different temperatures yield a sigmoid ESP versus temperature curve, which is used to determine a corresponding transition temperature Tsp based on statistical correlations previously established by comparison with the results of extensive standard tests This value can be correlated with the FATT50 (Foulds et al., 1992), the transition temperature measured by conventional impact tests and defined in the ASTM standard A 370-03a. Normally, Tsp is well below room temperature and 15–20 SP tests have to be carried out at several temperatures down to liquid nitrogen temperature, for an exhaustive characterization of the material of a given rotor section. Once Tsp has been estimated by SP tests, the fracture toughness, KIC can also be determined indirectly, via established Tsp versus KIC correlations (Bicego et al., 1995). The key problem remains the reliability and accuracy of such correlations, still a controversial subject on which much work is still to be done. Some of the correlations are freely available from the literature, but are generally affected by poor accuracy. Much work carried out in this field, however, is unpublished and sometimes exploited in service activities as an element of proprietary competitive advantage. For the low-alloy steels used in steam turbine rotors and other power unit components, more refined multiparameter correlations were obtained through extensive testing (SP and standard specimens) on scrapped rotors. These improved correlations
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2.11 (a) Extraction of a sample of material from a serviced steam turbine rotor for ‘small punch’ testing; (b) aspect of the rotor after cutting away a small material sample for SP testing; (c) actual size of the extracted material sample (courtesy of CESI).
enable estimations be made of σYS, σUTS, FATT50 and KIC as functions of Tsp (obtained from the SP tests) and of the additional parameters HV (Vickers hardness) and d (ASTM grain diameter), according to the logical framework of Fig. 2.12. HV and d can be measured in a completely nondestructive way. The uncertainty of the estimated FATT50 value is ±20 8C (Bontempi et al., 2005). Interest in the SP method for power generation concentrates mainly on the assessment of fitness for further service and of the residual life of power plant components, in particular HP–IP–LP steam turbine rotors, but also steam drums and headers. The industrial use of the SP method is limited, mainly because of a lack of recognized standard codes and procedures for deriving the relevant mechanical properties. Some international, interlaboratory cooperation tried to overcome this problem. In the frame of CEN activities, Workshop 21 on ‘Small Punch Test Method for Metallic
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2.12 Flow chart of a procedure adopted to estimate FATT50 and KIC of a serviced rotor material from ‘small punch’ test result and other non– destructive measurements (Hv = hardness; d = grain size determined with metallographic replicas) (courtesy of CESI).
Materials’ resulted in the publication of the European standard practice CWA 15627 (2006). CWA 15627 Part A describes the use of SP tests to determine creep properties of aged rotor materials; it is based on the results of a European interlaboratory exercise performed on a CrMoV steel (Ule et al., 1999). Japan is also going to publish a standard for this technique, as presented at the 1st International Conference on ‘Determination of Mechanical Properties by Small Punch and Other Miniature Testing Techniques’, held at Ostrava in 2010.
2.4
Condition monitoring of gas turbines
In the last two decades gas-fired power units have conquered the front stage in power generation; with their sophisticated design and manufacturing technology they are now a special challenge for on-line condition monitoring, due to the extremely high temperatures and harsh operating conditions of the gas path (Fig. 2.13). Historically, the first monitoring systems were mainly concerned with vibrations and parameters related to plant performance. Vibration monitoring in gas turbines is as well established and important as in steam turbines. Most of what has been said in section 2.3.1 for dynamic monitoring of steam turbines holds also for
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2.13 Schematic of a gas turbine: hot gas path is evidenced in grey; numbers indicate critical components: combustion chambers and liners (1), rotating blades (2) and stationary vanes (3).
gas turbines; for specific aspects see the overview by Maalouf, 2005, and books suggested at the end of this chapter. Nowadays standard monitoring systems can deliver more information, such as thermodynamic performance, flame detection, overspeed detection, combustor instability and (in some cases, as in the oil and gas industry) hazardous gas detection. Different monitoring modules can be purchased by the utility with the plant (details are given in the OEM’s web sites). Gas turbines used across many different industries are often supplied with transducers and monitoring systems in accordance with API 670 and API 616. Recently the ISO TC 192 GT committee has published the ISO 19860:2005 standard about trend monitoring systems (TMSs). This standard identifies specific advantages for the OEM and for the plant user, indicates a common terminology, lists parameters to be measured (by direct or indirect methods) and provides guidelines on how to compare plant performances. The typical TMS enables the user to analyse the current plant condition, to identify and reduce possible failures, and to detect short/longterm drifts; the short-term trend indicates the need for compressor washing while the long-term trends reflect the ageing of the unit. Appendix B of this standard underlines the need to develop physically based and/or analytical and empirical models, in order to increase the reliability of data interpretation. The standard does not mention at all the degradation of materials and coatings. Within the frame of current process and material technology, the early detection of operation-induced damage to critical components is crucial, not
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only for reliable and safe operation but also for cost-effective planning of maintenance interventions. Component-specific, performant innovative diagnostic techniques are required, as described in the following sections 2.4.1 to 2.4.3.
2.4.1 Combustion instabilities Some gas turbines for power generation were affected in the 1990s by a specific dynamic problem – strong pressure pulsations in the hot gases due to dynamic flame instability – that caused substantial damage to hot parts. Gas turbines with dry low emission (DLE; with lean pre-mixed particles) combustion systems are required by present regulations to meet extremely low emission limits; as a consequence they must operate on the verge of flame-out and this lean-burn operation can cause combustion instabilities, resulting in pressure pulsations in the combustor. Table 2.8 summarizes the type of oscillations that may occur, their possible causes, the components incurring potential damage and suggested mitigation strategies. Metal fretting of casing and support structures, as well as premature wear of sheet metal structures, such as combustor liners, are among the possible consequences (see also Brun and Kurz, 2009). Great effort was made in developing ad hoc combustion stability monitoring techniques to be implemented in an active control loop (Dowling and Morgans, 2005). Today’s large land-based gas turbines on the market are equipped with dynamic pressure sensors that can immediately detect and analyse the onset of pressure pulsations, thus preventing potential component damage. Modern dynamic pressure sensors enable the pulsation frequency range (0–5 kHz; see Table 2.8) of practical diagnostic interest to be kept under control; they are cheap and robust, designed and manufactured to withstand the harsh operating conditions inside the gas turbine. A combustion monitoring system (CMS) is often combined with the emission monitoring system (EMS), according to the above-mentioned ISO 19860 standard. Innovative techniques are being developed to study the fluid dynamics of new combustion chambers, which enable the investigation of process fluid density with high time and space resolution, even in harsh combustion chamber environments. Optical techniques are in principle most suitable, being non-invasive. One of the most promising seems to be a non-invasive optical technique based on Rayleigh scattering (Raush et al., 2010). Laser vibrometry is also interesting, since it is able to detect the main modes of resonant flame dynamics, to determine the density fluctuations and to obtain time and space maps of the instability dynamics (Giuliani et al., 2006, 2009).
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Table 2.8 Matrix assisting in identifying causes of combustion instabilities Description
Freq. range Component at Potential causes (Hz) risk of damage
Low0 to 25 frequency 25 to 100 dynamics
Swirler basket Flashback Nozzle indications Lean blowout Damaged swirler(s) Air flow restrictions High injection flow rates Pilot-nozzle distress
Intermediate- 100 to 500 frequency dynamics
Transition panel and/or seals (fretting/ wear) Intermediate- 500 to 1500 Downstream frequency components dynamics (fretting/ wear) 1500 to 5000 Basket HighCross-flame frequency tubes dynamics Flashback TCs
Fuel composition Fuel splits Bypass-valve distress
Mitigation strategies Increase pilot-stage fuel fraction Increase C-stage fuel fraction repair / replace the basket Remove air-side obstructions Reduce the injection flow rate Combustion tuning Active tuning
Equipment distress Inspect and repair/ replace combustor components Install Helmholtz Overfiring resonators IGV position error Adjust IGV position Fuel composition Increase steam System dampling injection Basket distress Pre-heat the fuel
Source: adapted from Combined Cycle, 3Q/2006, ‘Monitoring and mitigating combustion dynamics’, page OH-62
2.4.2 Electrically charged particulate at turbine exhaust for early detection of hot part damage Early warning of incipient damage of the high-temperature parts in modern gas turbines can help operators in taking the right decisions to avoid accidents and unplanned shutdowns. A method has been developed for this purpose, based on the detection and measurement of the electrostatic charge carried by the particulate contained in the exhaust gas flow at the turbine outlet. The amount of electrostatic charge carried by the particulate varies with operating conditions and is essentially a consequence of the highly turbulent flow. This ‘normal background’ is used as a reference ‘baseline’, to be established by a preliminary period of monitoring during normal operation. Significant deviations from this ‘baseline’ can be promptly detected and evaluated for diagnostic purposes. Abnormal loss of coating or blade material, due to on-going damage, will result for instance in an abnormal increase of the particulate, causing in turn an increase of the
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electrostatic charge signal and an unusual time-dependence of the signal with respect to the baseline behaviour (e.g. large bursts). Several processes can cause an increase of particulate and can therefore be detected by electrostatic charge monitoring: damage of combustion chamber and hot section materials, wear of seals and coatings, blade tip rubbing, as well as lubricating oil losses from bearing seals and foreign material ingestion. Basically, a small number (3–4) of charge sensors (a simple disc- or rodshaped conductive electrode mounted on the exhaust gas duct, exposed to the gas stream and connected to a charge amplifier) each provide a lowimpedance electric signal proportional to the currently detected local electrostatic charge. The working principle of the sensor is electrostatic induction. In the simplest configuration all signals are independently processed (HP filtering, A/D conversion). The parameters usually considered for diagnostic evaluation are: . . .
RMS of the charge signal over a time integration interval (typically 1 s). It correlates to the current rate of fine particulate in the exhaust gas. Number of signal samples above an assigned or dynamically updated threshold in the same time interval. This parameter marks the occurrence of relatively large particulate (e.g. > 50 μm). Frequency analysis of raw or envelope charge signal. It can provide hints on the possible origin of the particulate (e.g. moving or static machine components).
The system can be made rugged enough to operate indefinitely in a heavyduty gas turbine environment. It is important to note that not all processes causing a particulate increase are necessarily correlated to ongoing damage (a typical example is periodic compressor washing). Data processing and evaluation software must be tailored automatically to recognize and filter situations belonging to normal operation (e.g. compressor washing), which give rise to abnormal charge signal behaviour without diagnostic significance, so as to avoid unacceptable false alarms. A preliminary automatic screening and reduction of the large flux of incoming data is also highly desirable. The method is attractive because it is simple, robust and cheap. It was first considered in the 1970s for application to aeronautical engines and turbines, mainly out of safety concerns. Only later, in the 1990s, was the effort also directed to the development of the potential of using the method as a ‘condition monitoring’ method for power generation gas turbines (Lapini et al., 2001; Lapini and Zippo, 2003), specifically for giving early warning of those damages that cannot be detected with a satisfactory lead time through conventional methods of machine performance monitoring. Recently this method was also presented at the ASME Turbo EXPO Conference by the SWRI (Wilcox et al., 2010).
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2.14 EDMS sensor and cable (courtesy of CESI).
An example of results obtained on a 40 MW aero-derivative gas turbine during a monitoring campaign that lasted over 2 years (Lapini and Zippo, 2003) is shown in Figs 2.14 to 2.16. A single charge sensor (Fig. 2.14) was mounted on the exhaust diffuser (Fig. 2.15), in the upper part of the duct, free of obstacles inside. The electrostatic charge, as well as rotation speed and load level, were monitored for over 5 months, to identify an average baseline activity level (see Fig. 2.16). At a certain time (November 12th in the diagram) a sharp increase is observed and continues for a while, with observed peak amplitudes 3 times the baseline. An endoscopic inspection carried out as a follow-up of the charge monitoring data showed a loss of coating on about 20 % of the first-stage turbine blades. The turbine was judged to be fit for further service and operated for an additional 4 months, until it was shut down for refurbishment. This was done while the charge monitoring system provided further damage indications (right-hand end of the diagram). No other monitoring system installed on the machine was able to detect any structural or functional anomaly during such an operating period. The counterpart for the simplicity, ruggedness and low cost of the method is a lack of selectivity and precision in the diagnostic indication and, consequently, the risk of false alarms. Many causes can contribute to an increase of the electrostatic charge level, not all of them diagnostically
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2.15 EDMS sensor installed in the exhaust duct of a gas turbine (courtesy of CESI).
2.16 Example of EDMS results over a monitoring period of six months (courtesy of CESI) .
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significant; moreover, it is not easy to discriminate between the different potential causes of deviation from normal charge signal behaviour. This method should be used as a first alarm of potential onset of structural damage and a good reason to start, as soon as possible, more precise and targeted investigations. Preliminary collection and analysis of data on a significant number and variety of gas turbines for an equally significant time duration is the best way to consolidate its diagnostic potential and reliability.
2.4.3 Gas temperature and integrity of hot parts The temperature of hot gas at the turbine inlet (TIT) is a crucial parameter for the energy efficiency and the operation life of gas turbines. At present temperature and pressure of the gases in the hot section of the turbine are estimated by calculations and indirect measurements (typically exhaust gas temperature, measured by groups of two thermocouples for each burner). Such temperature values are introduced into analytical models of the turbine CMS, which indirectly estimates the TIT and the temperature of hot parts. Such an indirect approach has an intrinsic uncertainty of about 10– 15 8C. As a consequence the turbine has to be operated below its optimum efficiency, by about 1 %. Both in design validation and during normal operation a direct measurement of the temperature of gases entering the turbine would be very desirable, for the reasons stated above. Moreover, a direct, accurate measurement of the temperature of hot parts would be very useful in optimizing maintenance intervals, as a variation of ±10 8C in component temperature can induce a +30 %/–20 % variation in the lifetime due to oxidation and a +100 %/–50 % variation in creep life prediction. Currently available instrumentation cannot survive for very long in the turbine hot gas path due to the hostile environment – very high temperatures and pressures, high dynamical stresses, corrosive atmosphere – presenting unique challenges. An interesting summary of the technology gap in this area was worked out at the European level by evi-gti.com in 2006 (Lab Gap Matrix is available from www.evi-gti.com). A similar analysis was performed in the United States by the Propulsion Instrumentation Working Group (PIWG, www.piwg.org), which issued specifications for the missing instrumentation capabilities required ‘to fill the gaps’. Recent summaries of the state-of-the art in sensor needs and measurement requirements in the hot gas path of gas turbines can be found in Roberts (2004) and Niska (2009). The recently completed European project HEATTOP should also be mentioned among the efforts in this field. The project was finalized to reduce hot gas path measurement uncertainties, by developing new sensors and technologies for GT monitoring and by validating their use in test facilities
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reproducing the conditions of actual turbines (Mersinligil et al., 2010; see the website: http://ec.europa.eu/research/transport/news/article_8780_en. html). Optical techniques are the most promising to measure very high temperatures, being non-intrusive (Yu and Chow, 2009). On-line infrared (IR) pyrometry, aimed at a direct on-line measurement of the temperature of hot blades, is the most mature optical technique available nowadays. IR pyrometry is a line-of-sight non-contact optical technique, which requires a suitable fixture to enter the GT at a position that enables a direct view on the rotating blades. The mounting system (Fig. 2.17) provides the right view through a pressure-proof sight glass assembly; the viewing window can remain clean due to a purging air flow derived from the air compressor discharge. The optical signal collected by the pyrometer is transmitted through a rugged, flexible, fibre optic light guide to a remote electronic signal conditioner. If temperature is continuously recorded, a sequence of peaks is observed; each peak corresponds to an individual rotating blade passing in front of the pyrometer. This method can detect significant bladeto-blade temperature differences, possibly due to manufacturing variations or to obstructions of cooling passages. An overheated blade would show an associated peak amplitude higher than the measured mean amplitude (Fig. 2.18). A defective cooling would expose, for example, the first-stage blades to temperatures beyond 1500 8C, above their material limit, so that in many instances on-line pyrometry can be useful to check cooling efficiency and, consequently, to assure coating and blade integrity. The first trials with on-line pyrometry started in the 1970s (Mossey patent 1972) supported by the manufacturers’ interest. Several plant applications have been implemented since the 1980s (e.g. Beynon, 1981). The efforts for
2.17 Schematic of pyrometer installation into the GT (adapted from Amory and Hovan, 2000, LAND web site).
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2.18 Example of pyrometers installed in a GT (right) and results indicating an overheated blade (left).
the application of pyrometry on operating gas turbines were continued by R&D centres in the 1990s (e.g. EPRI in USA, Kema and CESI in Europe), in support of utilities confronted by new generations of GTs coming to the market. Examples of applications of pyrometry at the development level can be found in a recent GT handbook (Soares, 2007). When blades coated with zirconia–ittria ceramics were introduced, pyrometers using IR wavelengths of around 10 μm were also tested (Markham et al., 2002), due to the optical properties of this ceramic material. However, the near-IR wavelengths (around 1 μm) have noticeable advantages from a technological point of view for the following reasons: . . .
Glass is not transparent at 10 μm, so more expensive and complex optical windows and lenses would be required for this wavelength. At the shorter wavelength the combustion gases are transparent; they do not disturb the measurement and a higher sensitivity can be reached. A thin oxide layer grows on the surface of the hot parts during operation, which results in an emissivity very similar to that of the hot parts with metallic coatings used before the application of TBCs.
Today IR pyrometry is rather well known and applicable to the most important and common gas turbines (see the web sites of LAND and ROTAMAP). For many gas turbines the mounting system is designed and can be installed by the manufacturer. IR pyrometry can contribute to significant cost savings, by helping to prevent equipment damage and unscheduled shutdowns due to failure of overheated blades and by optimizing operating conditions. Furthermore, it could usefully contribute to the optimization of maintenance schedules and to the development of new GTs.
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2.5
In situ assessment of gas turbine hot parts by nondestructive techniques
Gas turbine users usually follow the maintenance schedule suggested by the OEM, which is based on very conservative criteria to increase plant reliability. During periodic inspections a visual check of the surface of hot gas path components can be done by microendoscopic examination, without opening the turbine case, to verify the integrity of the components and ceramic coatings. A small video camera connected to optical fibres allows a visual check to be made of critical positions on hot parts (mainly rotating blades) and to find possible TBC adhesion failures or local damage due to foreign objects (FOD). Very small cameras are nowadays on the market, with a very high number of pixels (30 000) and very high resolution. If the TBC is absent in very large areas a severe overheating of the substrate during further operation could occur and cause component failure. This in situ damage assessment enables sensible decisions to be taken about the possibility to refurbish damaged components and to avoid their premature replacement. Even if visual inspection gives a positive response, coating degradation can nevertheless proceed slowly according to several degradation mechanisms: deterioration of the protective capability of the bond coat (BC) due to interdiffusion with the base alloy; incipient defects and detachments; microcracking of the thermally grown oxide (TGO) and TBC, sintering of the TBC layer with consequent thermal conductivity and stiffness increase; growth of TGO at the metal/ceramic interface, with the development of high compressive stresses at the bottom of the TBC (Fig. 2.19). In practice, the coatings of hot parts of individual gas turbines that have undergone the same operation period could nevertheless have different life fractions, depending on the particular operation history experienced. Life prediction modelling of metallic coatings is nowadays well developed, but the actual reduction of coating durability due to fuel changes and/or higher degradation due to cycling cannot yet be reliably predicted. Moreover, coating assessment (both in the as-delivered and serviced condition) is even more critical for ceramic thermal barriers (TBC), as their premature failure can dramatically reduce component life. For these reasons a noticeable effort has been made to develop nondestructive techniques that make it possible to evaluate both the initial coating quality (coating thickness) and (possibly in situ) the residual protective capability of the coating after a known operation period (Ellingson et al., 2005; Osgerby et al., 2006). Some innovative inspection and NDT methods to assess hot part coating damage are presented in the following. These methods can be used to support cost-effective decisions about refurbishment needs, avoiding, on the
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2.19 Example of TBC coating as-delivered (left) and aged in cyclic oxidation at 970 8C (right).
one hand, the heavy cost penalties of irreversible damage of base material and, on the other hand, an equally penalizing premature scrapping of very expensive components.
2.5.1 Thermography to check blade cooling and coating adhesion As already mentioned for boiler waterwalls, thermographic techniques can be used to map temperature distributions in complex components; a classic example is the check of cooling channels in a gas turbine blade (Fig. 2.20). Thermography can also be used to detect coating delaminations. The area to be inspected is heated through either conduction or radiation; the surface temperature is monitored using a suitable imaging system, such as a scanning infrared camera with mK resolution. Delaminated areas appear as cold or hot spots, respectively. The technique is unable to detect cracks that are normal to the surface (Grice et al., 1983; Bento and Almond, 1995), but can used to determine some characteristics of the coating, like thermal properties. Pulsed thermography (Fig. 2.20), in particular, is effective – and is currently used – in detecting delaminated areas of metallic and ceramic coatings deposited using the most common deposition techniques. One or two flash lamps heat the component surface and an infrared camera records a time sequence of colour digital images during the cooling phase; at each
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2.20 Examples of a pulsed thermography configuration and (b) a thermal image of a blade with evidence of cooling channels (adapted from Ellingson et al., 2005).
time a cooling curve can be obtained from the image processing system. The cooling process is slower where a coating delamination is present, because the air layer between the coating and substrate interrupts the heat flow into the base material. The resolution of pulsed thermography (PT) depends mainly on the characteristics of the IR camera, on the thermal properties of materials, on coating thickness and on defect geometry (e.g. size and depth). Specimens with artificial defects of different sizes (several millimetres) are used for calibration purposes. Typical physical limitations of thermography, such as the influence of the emissivity of the inspected coating surface, can be overcome. If the emissivity of the coating surface is too low, as in the case of ceramic thermal barriers, a thin reversible graphite layer is first sprayed on to the surface. PT can be used as a quality control technique on new components, to detect coating delaminations and to find out local differences in coating thickness. Moreover, the reliability of PT was increased by the introduction of an algorithm based on the apparent effusivity (Marinetti et al., 2007), which permits discrimination between delaminations and variations in coating thickness. Being a fast, non-contact, single-side technique, PT allows the inspection of wide coated surfaces at a time. The PT equipment is transportable and applicable in situ (Rinaldi et al., 2008) if the components can be reached by the flash lamp radiation. Where a heat pulse cannot be artificially induced, natural thermal transients can be exploited by recording a sequence of images by an IR camera; differences in cooling times can indicate local anomalies. An on-line application of IR transient thermography was developed to monitor local overheating of rotating blades due to TBC delamination or blockage of cooling holes (Zombo, 1997; Zombo and Shannon, 2006).
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Transient infrared methods are now so mature that an OEM has developed a quality control system to evaluate turbine airfoil cooling holes during blade manufacturing (Allen et al., 2010). The blades are mounted on a fixture where fast internal heating with hot air is obtained in a few seconds; during the following cooling phase an infrared camera guided by a robotized system takes IR pictures at fixed positions. The images recorded on each blade during cooling are compared with reference images of a sound blade, taken at the same positions. From this comparison hole manufacturing defects can be quickly and automatically detected in a much more reliable way than with the previously used manual methods.
2.5.2 Advanced eddy current method to evaluate coating thickness and depletion In many cases the actual thickness of the metallic coatings applied to hot parts for protection against oxidation, corrosion and high temperatures cannot be reliably controlled after deposition with traditional NDTs, for one or more of the following reasons: very low overall thickness (some hundreds of micrometres), multilayer structure, small difference in conductivity between some of the layers. Indirect methods are normally used during thermal spraying of big parts, but a direct method is desirable. Moreover, the protective power of MCrAlY bond coats is due to the presence of the aluminum-rich β phase; its dissolution is a thermally activated, diffusion controlled phenomenon, which can today be reliably modelled. During operation two β-phase depleted regions form on the two sides of the coating, near the base material and at the upper surface. Knowing the thickness of such depleted layers, the residual life of the bond coat can be estimated using diffusion and oxidation laws together with inverse problem solution techniques (Krubovsky et al., 2001; Rinaldi and Mandelli, 2009). Consequently, a non-destructive method that enables accurate measurements of both the thickness of the as-delivered parts and the residual effective coating thickness after an operation period is highly desirable, in order to check the integrity and to evaluate the residual operation life of the coatings. The frequency scanning eddy current method (F-SECT) (Antonelli et al., 1997, 1998, 2001; Rinaldi and Antonelli, 2005) is able to detect not only the thickness of the BC and TBC but also the ß-phase depletion of the BC. The method is based on the same physical principle as the well-known eddy current NDT method, but the eddy current probe is excited with a variable frequency (typically from 50–100 kHz to 1–10 MHz) to gradually scan higher to lower depths and to detect the layers with different electrical properties. The spatial resolution is determined substantially by the probe
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diameter (about 8–10 mm). The material response (electrical impedance versus frequency) is recorded; the qualitative trend of the curve is already an indication of coating degradation with respect to a new coating (Fig. 2.21). If the results are analysed with suitable multilayer models, quantitative data can be obtained, such as the thickness of a MCrAlY coating on an Ni-based alloy of a hot part. The measurement of the metallic coating thickness is possible notwithstanding the small conductivity difference between the base metal and the MCrAlY coating. The F-SECT method is also applicable in the presence of yttria-stabilized zirconia (YSZ) TBC coatings, because the external ceramic layer is transparent to the probing electromagnetic field; if the probe is in good contact with the TBC surface, the TBC thickness can be estimated by exploiting the lift-off signal. Taking account of some error sources (small variations in the base metal conductivity, surface roughness, geometry effects due to the curvature of the blade surface), the technique can be used during the manufacturing of gas turbine blades and vanes for coating thickness control. An accuracy of 15 % or 30 % was demonstrated in the measurement of the thickness of as-sprayed and smoothened
2.21 Electrical impedance vs frequency plots obtained on a new blade (left) and a blade after service exposure in the plant (right). Anomalous curves on the right correspond to coating areas with service-induced damage (ß-phase depleted areas). The micrograph is taken on the section of the worst situation detected (opposite curvature): A = TBC; B = effective bond coat (BC); C = depleted BC; D = anomalous precipitation at base material interface.
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MCrAlYs, directly or through a thick ceramic thermal barrier coating, respectively (Jungbluth et al., 2008). The qualitative shape of the impedance versus frequency (from 50 kHz to 1 MHz) plots taken at different positions after service exposure in the plant already gives evidence of the degradation level of the coated material Fig. 2.21). The MCrAlY coatings change their electromagnetic properties according to the temperature history and operation time; F-SECT is able to detect the level of metallic coating degradation, in particular the β-phase depletion of the bond coat for γ/β coatings. A quantitative parameter of the degradation level is determined on basis of a multiple (n)–layer model with a task-dependent allocation of unknowns (e.g. the thickness of different layers). A good correspondence is generally found between the degradation levels measured by this NDT and by metallography after sectioning, as shown in the example of a failed component analysed in Rinaldi et al. (2006). The dependency of the impedance (Z) on temperature (T) and the effective coating thickness varies with the chemical composition both in γ/βand in γ/γ´-type coatings (in the latter case differences in volume fraction of the precipitated Cr-rich phases are important). If the Z–T relationships have been previously determined experimentally for each class of metallic coating, e.g. on coated specimens aged in a furnace at known temperatures for increasing times, the F-SECT technique can be used to determine the temperature of the measured area. As a consequence, using this technique and an automated robotic system it is possible to establish accurate and high-resolution maps of the surface temperature distribution over the entire airfoil and platform areas (Schnell et al., 2007). The exact location of hot spots on the airfoil can easily be identified. The F-SECT technique can also be successfully applied in situ after service exposure in the plant of the blades still mounted during maintenance stops in industrial GTs (Rinaldi and Antonelli, 2005, or Rinaldi et al., 2006). The sensor developed for on-site use and an example of results obtained in field applications are shown respectively in Figs 2.22 and 2.23. Finally, in bare regions of ex-service components (coating eroded or delaminated) the technique can determine the depth of a base material layer affected by nitridation, the precipitation of acicular nitrides severely embrittling the alloy (Jungbluth et al., 2008).
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2.22 Example of in situ frequency-scanned eddy current (F-SECT) measurement (adapted from Rinaldi and Antonelli, 2005).
2.23 Results of frequency-scanned eddy current measurements (FSECT) on a blade section after service exposure in the plant.
2.5.3 Photoluminescence piezospectroscopy to detect electron beam physical vapour deposition TBC degradation This technique was invented by Clarke (Tolpigo and Clarke, 2000), and is based on: .
the transparency to the visible light of the zirconia–yttria with columnar
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morphology made by electron beam physical vapour deposition (EBPVD) and the fluorescence in the red of the Cr3+ ions present inside the TGO alumina layer. If excited by a green laser light, the outer electrons of the Cr3+ ions can rise to a higher energy level. Falling again to a lower energy level they produce a red fluorescence radiation (photoluminescence). Initially the technique was studied through a typical Raman spectroscopy instrument combined with an optical microscope, as shown in Fig. 2.24.
It is well known that the position and the shape of fluorescence peaks depend on the crystalline structure of the material and on its elastic strain level (and corresponding stress). A spectroscopic analysis of the fluorescence spectrum of TGO highlights the presence of different alumina phases (α, θ, γ), as shown in the example of Fig. 2.25, which concerns a vertically cracked air plasma sprayed YSZ. Here the doublet relative to the α phase (R1 and R2) and the doublet relative to the θ phase (T1 and T2) are clearly visible, together with the broad peak of γ alumina. As a consequence, this technique can be helpful in quality control of columnar TBC coatings, since the presence of the undesired θ phase gives evidence of improper heat treatments during manufacturing. Since the frequency shift Δν of the main peaks (R1 and R2), compared to the corresponding peaks of a stress-free alumina (see Fig. 2.26 top left), is proportional to the residual stress level of the alumina layer, the technique is called photoluminescence piezospectroscopy (PLPS). The following formula permits quantitative TGO stress evaluations to be obtained: σ
TGO
= – Πij Δν
where the tensor Πij was determined once by laboratory measurements on an alumina sheet at increasing levels of controlled bending strain (and corresponding stress). Since the TGO peaks shifted towards lower energy levels, the corresponding residual stress σTGO is compressive (note the minus sign in the formula). At the beginning of the TBC life, the TGO is growing while remaining flat and in-plane, and the order of magnitude of σTGO is usually about 2–4 GPa. High temperatures and thermal cycling during operation cause local irregularities to appear at the TGO/BC or TGO/TBC interfaces. As a consequence, the TGO residual stress is relieved and the PLPS doublet shifts toward a stress-free condition, as shown in the sequence of Fig. 2.26. Several parameters have been proposed to evaluate the evolution of the PLPS spectra during TGO degradation (Selc¸uk and Atkinson, 2002), but the R2 peak position seems to be the most stable (minimum dependence on temperature) and significant parameter. In order to determine it reliably, a
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2.24 (a) Schematic of the PLPS technique: excitation by means of an incident green laser (e.g. 514 nm) and detection of TGO (thermally grown oxide) photoluminescence (693 nm) through the EB-PVD thermal barrier coating (TBC). (b) Columnar structure of a typical EB-PVD TBC.
deconvolution process of the doublet R1 and R2 has to be performed by using the sum of a Gaussian and a Lorentzian function. If the bond coat is sufficiently rigid, no significant rumpling occurs and the compressive stress is progressively released, so that the spectra do not contain several different doublets to be deconvoluted with complex routines. This can happen with
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2.25 PLPS spectra of the TGO with evidence of different alumina phases (α, θ, γ). The spectrum with the γ-phase peak was detected through the segmented APS TBC (see the bottom micrograph) after reducing the TBC thickness to 120 μm (adapted from Antonelli et al., 2006).
the much softer bond coats used in the aeronautic field and may give problems. For calibration purposes, a measurement on an alumina stressfree sample has to be done just before and after a set of measurements is taken. The capability of this technique to follow the degradation of the EB-PVD TBC system has been demonstrated by several authors (Gell et al., 2004, on aeronautic bond coats and Rinaldi et al., 2008). The required condition is that a monotonic behaviour of σTGO is measured on specimens aged in laboratory tests. In several laboratories aged specimens have been measured and a monotonic trend of σTGO with exposure time and number of thermal cycles has been found. A portable system to perform measurements in situ was developed and applied during a maintenance stop (De Maria et al., 2006; see Fig. 2.27). The
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2.26 Left: comparision of PLPS spectra (black) of specimens aged in cyclic oxidation tests at 1100 8C with reference spectra of stress-free alumina (grey); Δλ of resonance peaks decreases with increasing damage due to cycling. Centre: corresponding cumulative frequency distributions of σTGO values mapped over the whole surface of the exposed disc coated by EB-PVD TBC (300 μm thick). Right: morphology of TGO at the TBC/BC interface.
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different values of the TGO residual stress highlight local damage differences between well film-cooled areas and hot areas of the rotating blades. The parameter σTGO has been shown to be an indication of TBC life expenditure on blades with NiCoCrAlY–Re bond coats of land-based GTs (Rinaldi et al., 2010). Moreover a fast PLPS mapping of σTGO on hot parts coated with EBPVD TBCs points out the most critical areas of the component; statistical analyses of hundreds of measured σTGO values give reliable indications of σTGO reduction and corresponding interface degradation, enabling an early detection of damage before dangerous delaminations occur (Del Corno et al., 2010). During laboratory ageing of an EB-PVD TBC (Fig. 2.26), mapping of the specimens was performed at increasing numbers of thermal cycles; the cumulative probability distributions of the hundreds of measured σTGO values change significantly (Fig. 2.26; mean value decreases to < 1 GPa and standard deviation increases as the interface damage increases). This shows that this non-contact optical technique is very robust if used with an automatic system. Finally a case can be made for the advantages of an integrated approach, based on the cooperation of the three different NDT techniques (the capabilities are summarized in Table 2.9). These techniques provide complementary diagnostic information and performances: while the capability of IR imaging as a wide-area non-contact technique can be exploited for quick and effective in-shop screening of coating adhesion, FSECT can provide more detailed information on coating layer structure and on coating functional integrity and fitness for service, both in-shop and infield, even during short-term shutdowns, thus becoming a valuable condition assessment tool. PLPS, in turn, can be useful in the shop both during the process of developing and qualifying new coatings and in giving very early warning of the onset of EB-PVD TBC coating structural and functional degradation after an operation campaign. Such an integrated approach is widely explained in Antonelli et al. (2006), where different coating damage aspects are detected both on thermal-cycled samples and on service-exposed components coated with TBCs: BC depletion is detected by means of an F-SECT eddy current system, TGO/ BC interface degradation by the portable PLPS system and delaminations by pulsed thermography. As an example, compare the results shown in Figs 2.23 and 2.27 on the same component section.
2.6
Remote monitoring solutions
Advanced power plant diagnostics requires the availability of experts, an ever more valued resource in times of hard competition in the power sector. One effective solution to this problem is to centralize the diagnostic service,
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2.27 Example of PLPS results on a blade section after service exposure in the plant (compare with Fig. 2.23). PLPS can detect the regions of an ex-service blade exposed to the highest temperatures (lowest TGO stress values).
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Table 2.9 Comparison of methods available for in situ non-destructive evaluation of hot part coatings Technique
Detection capabilities Limitations
Pulse . Detachment at thermography TBC/BC or at BC/ substrate Interfaces
F-SECT
PLPS
Applicability in field
Applicable to There must be a components if they discontinuity at an can be reached by interface, typically flash lamps wider than a few mm parallel to the interface Requires preliminary Applicable also on mounted blades, calibration on each with open turbine type of metallic case coating
. MCrAlY or bond coat degradation (quantitative parameters) . TBC thickness . Very early stages Only for EB-PVD of damage at TBC/ TBCs BC interface can be Application to detected on EBvertically cracked PVD TBC (up to APS up to 100 μm 350 μm thick) thick is under . TGO residual study stress level (quantitative)
Portable instrument with optical fibre; applicable on dismounted blades
sharing all the relevant data coming from several plants – or from specific subsystems of these plants – within a common expert centre able to provide the assistance of highly qualified experts to solve the problems of individual plants. The unceasing development of data processing, storage and transmission capabilities makes it relatively easy to share data in real time with a remote centre. Moreover, the relevant parameters describing the current status of a main component can be stored to make possible the reconstruction of its previous operation history. The problem is then to build efficient data processing and screening procedures to extract diagnostically significant information preventing local and remote diagnostic facilities from being submerged by huge amounts of non-significant information. The associated aspect of data security can also be solved efficiently and satisfactorily in response to any well-defined requirement. Different access levels can be provided to data of increasing sensitivity. This approach offers a number of advantages. Personnel of the power production plant can get real-time assistance from top-level experts when required, at drastically reduced costs. The diagnostic centre experts can use their time efficiently, since they mostly avoid time-consuming travel and use their potential at its best, since they have their analysis tools and collected case histories readily to hand. Such centralized acquisition and analysis of diagnostic data is done according to unified criteria. This greatly accelerates
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the growth of diagnostic capabilities within the company, enhancing early malfunction or damage prevention and detection, as well as the possibility of corrective actions. This aspect is of particular relevance for manufacturers (see, for example, Killich, 2006, or www.ge-mcs.com), since they can collect in a relatively short time a large amount of operational experience on the components or machines they supply, and can optimize their service activity to their own benefit as well as to their clients. A variety of such centres has been developed in past years: . .
.
2.7
OEM centres very efficiently assure on-line assistance to their customers, as required in a long-term global assistance contract. Power generation company centres contribute to managing the generation fleet in the most efficient way. They are not limited to diagnostic monitoring, but rather support overall optimization of the power production strategy. Third party service company centres, including specialized university centres.
Future trends
2.7.1 Trends in condition assessment and monitoring policies Strong market competition in a frame of safety and environmental constraints is the main driver for developments in condition assessment and monitoring policies in thermal power plants. Besides assuring the required level of safety inside and outside the plant and compliance with environmental regulations, condition assessment and monitoring activities must be functional to the economic optimization of the entire power unit life cycle and are an integral part thereof. The individual plant itself may undergo different condition assessment and monitoring policies, depending on its specific operation mode within the power generation fleet (base load, intermittent, peak only, backup, etc.). For example, an older combined cycle of comparatively low efficiency, which is deemed to be replaced in a couple of years by a new-generation, high-efficiency counterpart, requires focus on short-term availability and reliability issues. The focus will be quite different for the new plant, although the available condition monitoring tools are essentially the same. Another important point is that utilities increasingly tend to contract the technical task of condition assessment and monitoring, as far as main plant items are concerned (boiler, steam and gas turbine, generator) to the manufacturers in the form of long-term global service contracts. The potentially unpredictable maintenance costs, as well as the related
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unavailability costs, are correctly treated as a risk and transformed into a fixed cost for the utility by means of long-term service contracts, essentially insurance contracts. The risk is thus transferred to the manufacturer, who is in turn stimulated to improve condition assessment and monitoring capabilities, as one way to control risk and to make the service business profitable. As a consequence, the most recent innovations in the field of condition assessment and monitoring for power plants have been introduced or are under development by manufacturers. As far as technical aspects are concerned, besides the above-mentioned thrust towards remote monitoring, the following general trends are observed: . . .
.
.
increasing automation, both in manipulation (ultrasonic probe scanners, etc.) and in data processing, screening and presentation; increasing tendency to present diagnostic results in the form of immediately usable images (ultrasonic, tomographic, etc.); continued interest for non-contact, non-local, wide-area diagnostic tools (a good recent example is the examination of long portions of piping by guided ultrasonic waves to detect corrosion and other types of damage in non-accessible areas); continued interest for improved models of the power unit or main subsystems thereof, to be used in connection with (reliable) monitoring data to enable relatively cheap, quick and precise understanding of the current condition; development of sensors resistant to harsh environments.
Efforts are also ongoing towards the use of neural networks and artificial intelligence (AI) techniques to support decisions about the machine condition (e.g. Adgar et al., 2009).
2.7.2 Some examples of technology innovations Substantial technical innovations in the past decade have widened the range of problems that can be addressed and the range of solutions available for condition monitoring and assessment. The dramatic progress in telecommunications has also had relevant fallouts in the field of industrial diagnostics. A significant example is provided by wireless sensors and local wireless data collection and exchange nets, which have achieved noticeable practical success and industrial penetration in the last decade. Wireless solutions have become available for a variety of monitoring tasks, from vibration monitoring of rotating machinery to environmental sensing. Wireless sensor subsystems can be an integral part of large local and remote monitoring systems for power generation units. The international standard for wireless communication in process automation (IEC 62591Ed. 1.0) was
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approved on 26 March 2010. The availability of wireless sensors and wireless data collection systems allows a substantially lower time, complexity and cost for the installation of monitoring systems and is of special relevance where a condition assessment and monitoring intervention is required at short notice and for a definite time duration (e.g. on a relatively isolated rotating machine). Some examples of efforts to develop new sensors are shortly described in the following. The aim is to obtain sensors and monitoring solutions that withstand harsh environments and are applicable in closed-loop control or can be implemented into an open-loop advisory/diagnostic system. Optical technique to measure gas turbine inlet temperature An innovative optical non-contact system to measure on line the temperature of gases at the gas turbine inlet (turbine inlet temperature, TIT) has been developed and is actually in the testing phase. The technique (Gianinoni et al., 2007) is based on the measurement of the IR emission of hot gases. The IR emission is detected by an optical probe mounted in a through-thickness penetration of the turbine case and positioned in such a way that the line of sight of the probe is perpendicular to the gas flow direction. Within the angle of view of the probe, the system carries out photometric measurements of the IR radiation emitted by the CO2 molecules present in the combustion gases in an appropriate wavelength band. The IR wavelength band is selected to meet the following requirements: it should undergo strong absorption along the test optical path (to avoid contributions from the opposite hot wall) and, at the same time, be transparent enough to allow the emitted radiation to cross a significant portion of the test region. The prototype sensor is mechanically robust and has been tested on a full-scale combustion rig (ENEL – Sesta Testing Facility, Italy). Results have shown that the recorded signal trend is in good agreement with the adiabatic temperature derived from the combustion process parameters. Recently, in the frame of the EU project HEATTOP, a solution based on multiwavelength detection, aimed at obtaining an absolute measurement of the gas temperature, has been investigated Silicon carbide (SiC) high-temperature sensor A novel silicon carbide (SiC) optical ultraviolet (UV) dual-diode flame temperature sensor (FTS) has been developed for an active combustion pattern factor controller (APFC) for gas turbines by General Electric’s
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Global Research Center and GE Energy and applied in harsh environments simulating a real combustion situation (Palmer et al., 2009). Smart components Continuous on-line condition monitoring of life-limiting components in the hot gas path sections of turbines would be highly desirable to optimize maintenance. For some years efforts have been ongoing (Mitchell et al., 2009) to build smart, self-aware engine components that incorporate embedded, harsh-environment-capable sensors and high-temperature-capable wireless telemetry systems. The innovative technology approach involves two aspects. First, sensors are embedded on complex shapes, such as turbine blades, by the MesoScribe’s MesoPlasma™ robotic thermal spray deposition system (Longtin et al., 2004), which can deposit finefeature patterns on flat and conformal surfaces (both thermocouples and strain gauges are under development). Second, wireless telemetry systems are embedded in less hot and harsh regions of the component (e.g. the root of the rotating blade), where temperatures prevent the use of conventional silicon-based devices but high-temperature electronics can be used.
2.8
Sources of further information and advice
A general and systematic description of the basic concepts and techniques of condition monitoring for power plants can be found in Handbook of Condition Monitoring (Rao, 1996, pages 285–324). For more details about boilers, see Boiler Operator’s Guide (Kohan, 1997) and Power Boiler Design, Inspection, and Repair (Malek, 2005) can also be used. An example of a leak monitoring system can be found at www. acousticmonitoring.com. A wide literature exists about vibration monitoring (see, for example, Kelly, 2006, for Advanced Vibration Analysis). For a systematic description of conventional non-destructive testing methods, vol. 17 of ASTM Metals Handbook: NDE and Quality Control is recommended; for a recent update on thermographic techniques, see, for example, chapter 14 of Chen (2007). Regarding the use of Small Inspection Vehicles for Non-destructive Testing Applications, see, for example, Friedrich et al., 2006. A systematic description of the most common techniques used in gas turbine condition monitoring can be found in the recent Gas Turbines: A Handbook of Air, Land, and Sea Applications (Soares, 2007). The applications of pyrometry and thermography are particularly significant. A comprehensive review of the non-destructive techniques developed for hot part coatings and applicable to thermal barrier coatings can be found in
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the dedicated chapter 8 of the book on High Temperature Coatings (Bose, 2007).
Web sites Corrosion monitoring http://www.kema.com/Images/KEMCOP.pdf http://www.lesman.com/unleashd/catalog/analytical/Honeywell-CET5000/ 34-SC-03-02.pdf http://www.lesman.com/unleashd/catalog/analytical/Honeywell-CET5000/ 34-SC-03-01.pdf http://www.osti.gov/bridge/purl.cover.jsp;jsessionid=AC1FD1732C579 B9B2D6326D60C080F49?purl=/895400-wsRNR4/ Boiler temperatures http://www.infra-view.com/whitepaper.htm Gap matrix EVI-GTI, the lab gap matrix, on the WWW, at http://www.evi-gti.com. Excel file PIWG, sensor specifications, on the WWW, at http://www.piwg.org http://gltrs.grc.nasa.gov/reports/2004/TM-2004-213202.pdf Further information on the HEATTOP project http://ec.europa.eu/research/transport/projects/article_6499_en.html http://ec.europa.eu/research/transport/news/article_8780_en.html http://www.evi-gti.com/default.asp?contentid=1227 http://www.msm.cam.ac.uk/UTC/thermocouple/pages/HEATTOP.html Pyrometers http://www.lirkorea.com/Landinstruments.net%20Website/combustion/ downloads/index.htm http://www.rotadata.com/pages/products/optical_pyrometry1.php Ceramic thermocouples http://www.rowantechnologies.co.uk/research_and_development.htm Thermography Grote K H, Antonsson E K (2009) Handbook of Mechanical Engineering, Springer edition, vol. 10, pp. 135–136.
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References
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testing methods,’ in Proceedings of ASME Turbo Expo 2008, 9–13 June 2008, Berlin, Germany, paper GT2008-51535. Kelly S G (2006) Advanced Vibration Analysis, CRC Press, USA, ISBN 9780849334191. Killich M (2006) ‘Operational flexibility for steam turbines based on service contracts with diagnosic tools’, in PowerGen 2006, Koeln Germany. Kim J S (1989) ‘Range and accuracy of speckle displacement measurement in double-exposure speckle photography’, Journal Optical Society of America, 6, 675–681. Kohan A L (1997) Boiler Operator’s Guide, McGraw-Hill. Kovacevich J J, Sanders D P, Robertson M O, Nuspl S P (1995) ‘Recent advances in the application of acoustic leak detection to process recovery boilers’, in TAPPI Engineering Conference, 11–14 September, 1995 Dallas, Texas, paper BR-1594. Krubovsky P, Kolarik V, et al. (2001) ‘Theoretical and experimental approach for long term modelling of oxidation and diffusion processes in MCrAIY coatings’, in Lifetime Modelling of High Temperature Corrosion Processes, Eds M Schultze, W J Quaddakers and J R Nicholls, EFC Publication no. 34, pp. 233–245, Maney Pub, ISSN 1354-5116. Lapini G L, Zippo M, Tirone G (2001) ‘The use of electrostatic charge measurements as an early warning of distress in heavy-duty gas turbines’, in Proceedings of ASME Turbo Expo 2001, June 2001, New Orleans , Louisiana. Lapini G, Zippo M (2003) ‘Some experiences about the use of electrostatic charge measurements in the exhaust gases as an early warning of distresses in industrial gas turbines’, in Proceedings of CAME-GT Conference, 10–11 July 2003, Brussels. Lapini G, Vario R, Zanetta G A (2005) ‘Experience in advanced remote vibration monitoring and diagnostics of largine turbine–generator sets’, in Power Gen, 2005, Milan, Italy. Lee K Y, Kim B -H, Velas J P (2004) ‘Development of an intelligent monitoring system with high temperature distributed fiber-optic sensor for fossil-fuel power plants’, in Proceedings of Power Engineering Society General Meeting, 2004, IEEE, vol. 2, pp. 1989–1994, ISBN: 0-7803-8465-2, DOI: 10.1109/ PES.2004.1372729. Longtin J, Sampath S, Tankiewicz S, Gambino R J, Greenlaw R J (2004) ‘Sensors in harsh environments by direct-write thermal spray’, IEEE Sensors Journal, 4 (1), February. Lyon R (2000) ‘The development of condition monitoring tools for the power generation industry’; available from http://www.ndt.net/article/wcndt2004/pdf/ power_generation/798_lyon.pdf. Maalouf (2005) ‘Gas turbine vibration monitoring – an overview’, Orbit, 25(1), 48– 62. Malek M A (2005) Power Boiler Design Inspection, and Repair: ASME Code Simplified, McGraw-Hill. Mao X, Takahashi H (1987) ‘Development of a further-miniaturized specimen of 3 mm diameter for TEM disk (f3 mm) small punch tests’, Journal of Nuclear Materials, 150, 42–52. Markham, Latvakoski, Frank, Lu¨dtke (2002) ‘Pyrometry of combustion turbine blades with thermal barrier coating’, EPRI Technical Report 1004336.
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Marinetti S, Robba D, Cernuschi F, Bison P G, Grinzato E (2007) ‘Thermographic inspection of TBC coated gas turbine blades: discrimination between coating over-thicknesses and adhesion defects’, Infrared Physics and Technology, 49, 281–285. Mersinligil M, Brouckaert J F, Desset J (2010) ‘First unsteady pressure measurements with a fast response cooled total pressure probe in high temperature gas turbine environments’ in Proceedings of ASME Turbo Expo 2010, paper GT2010-23630. Mitchell D, Kulkarni A, Lostetter A, Schupbach M, Fraley J (2009) ‘Development and testing of harsh environment, wireless sensor systems for industrial gas turbines’, in Proceedings of ASME Turbo Expo 2009: Power for Land, Sea and Air, paper GT2009-60316. Morgan B C, Foster C L (1995) ‘Acoustic emission monitoring of high-energy steam piping, Volume 1: acoustic emission monitoring guidelines for hot reheat piping’, EPRI, TR-105265-V1, Project 1893-20. Morgan B, Tilley R (1999) ‘Inspection of power plant headers utilizing acoustic emission monitoring’, NDT& E International Journal, 32(3), 167–175; available from http://pndmx.comze.com/articulos/Inspection_of_Power_Plant_Headers_ Utilizing_Acoustic_Emission_Monitor.pdf. Morris A, Dear J, Kourmpetis M (2006) ‘Monitoring high temperature steam pipes using optical strain measurement techniques’, Applied Mechanics and Materials, 5–6, 145–152. Morris A, Dear J, Kourmpetis M, Maharajai C, Puri A, Fergusson A (2007a) ‘Monitoring creep strain in power station engineering plant’, Applied Mechanics and Materials, 7–8, 31–36. Morris A, Kourmpetis M, Dear I D, Sjo¨dahl M, Dear J P (2007b) ‘Optical strain monitoring techniques for life assessment of components in power generation plants’ Proc. IMechE, Part A: Journal of Power and Energy, 221, 1141–1152. Morris A, Kourmpetis M, Dear I D, Sjo¨dahl M, Dear J P (2009) ‘Optical strain measurement techniques to assist in life monitoring of power plant components’, Journal of Pressure Vessel Technology, 131(2), 024502 (8 pp.) DOI: 10.1115/1.3062935. Muravin B, Muravin G, Lezvinsky, L, (2004) ‘Revealing creep associated flaws and industrial flaws in operating high energy piping by quantitative acoustic emission method’, in EPRI 4th International Conference on Advances in Material Technology for Fossil Power Plants, 25–28 October 2004, Hilton Head Island, South Carolina, pp. 646–658. Niska H (2009) ‘Requirements for advanced high temperature instrumentation and measurements in gas turbine engines’, in Advanced High Temperature Instrumentation for Gas Turbine Applications, VKI LS 2009-06. Osgerby S, Rinaldi C, De Maria L, (2006), ‘Recent advances in NDE of gas turbine coatings’, in Proceedings of Liege Conference on Materials for Advanced Power Engineering, Eds. J Lecompte et al., vol. 1, pp 217–227. Palmer C A, Abel R L, Sandvik P (2009) ‘Application of silicon carbide photodiode flame temperature sensors in an active combustion pattern factor control system’, in Proceedings of ASME Turbo Expo 2009, paper GT2009-59023. Parker J D, Stratford G C, Shaw N, Spink G, Tate E (1995) ‘Deformation and fracture processes in miniature disc tests of CrMoV rotor steel’, in Proceedings
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of Third International Charles Parsons Turbine Conference, Institute of Materials, vol. 2, pp. 418–428. Pertew M, Sun, X, Kent Golden R, Marquez H (2008) ‘A new blowdown compensation scheme for boiler leak detection’, Proceedings of American Control Conference, 11–13 June 2008, Seattle, Washington. Rao B K N (1996) Handbook of Condition Monitoring, Elsevier Science Ltd, pp. 285– 324, ISBN: 1 85671 234 1. Rausch A, Fischer A, Konle H, Gaertlein, A, Nitsch S, Knobloch K, Bake F, Ro¨hle I (2010) ‘Measurements of density pulsations in the outlet nozzle of a combustion chamber by Rayleigh-scattering searching entropy waves’, in Proceedings of ASME Turbo Expo 2010, paper GT2010-22492. Rinaldi C, Antonelli G (2005) ‘Epitaxial repair and in situ damage assessment for turbine blades’, Proc. IMechE, Part A: Journal of Power and Energy, 219, 93– 99. Rinaldi C, Mandelli M (2009) ‘The role of oxidation and diffusion kinetics in a coating life prediction code – application to components’, in ASME Turbo Expo 2009, 6–12 June 2009, Orland, Florida, paper GT 2009–60063. Rinaldi C, Bicego V, Colombo P P (2006) ‘Validation of life management system by case histories, on line measurements and in situ NDT’, Journal of Engineering Gas Turbine and Power, 128, 73–80. Rinaldi C, De Maria L, Del Corno A, Cernuschi F, Gianinoni I (2008) ‘Development of innovative non destructive methods for TBCs and monitoring techniques in the frame of an Italian R&D programme’, in Proceedings of the 4th International Conference: The Future of Gas Turbine Technology, 15–16 October 2008, Brussels, Belgium, paper IGTC08_P11. Rinaldi C, De Maria L, Mandelli M (2010) ‘Assessment of the spent life fraction of gas turbine blades by coating life modelling and photoluminescence piezospectroscopy’, Journal of Engineering Gas Turbine and Power, 132(11), 114501 doi: 10.1115/1.4000804. Robbins B J, Farrell D M, Sikka P, Seaman M (2004) ‘Non-intrusive sensor systems for monitoring the thermal conditions of furnace walls’, in EPRI 5th Intelligent Sootblowing Workshop, Nashville, Tennessee. Robbins B J, Farrell D M, Wilkins C J (2009) ‘Non-intrusive scanners for monitoring the thermal conditions of furnace walls – recent applications’, Rowan Technologies Ltd, Manchester, UK; available from www. rowantechnologies.co.uk. Roberts J (2004) ‘Requirements for instrumentation technology for gas turbine propulsion systems’, in Advanced Measurement Techniques for Aero Engines and Stationary Gas Turbines, paper VKI LS 2004-04. Rodgers J M (2007) ‘Acoustic emission of seam-welded high energy piping systems in fossil power plants’, Journal of Acoustic Emission, 25, 286–293. Rodgers J, Tilley R (2004) ‘Standardization of acoustic emission testing of fossil power plant seam-welded high energy piping’, in ASME Pressure Vessels and Piping Conference, 25–29 July 2004, San Diego, California, vol. 471, pp 113– 131. Rodgers J M, Morgan B C, Tilley R M (1996) ‘Acoustic emission monitoring for inspection of seam-welded hot reheat piping in fossil power plants’, in Proceedings SPIE, 2947, 126.
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Rovder J S, Yavorsky R W, Carlyle J M (1988) ‘Acoustic detection of leaks in steam piping and headers’, in Proceedings of 3rd EPRI Incipient Failure Detection Conference, EPRI CS-5395, 10–31 through 10–42, Electric Power Research Institute, Palo Alto, California, August 1988. Schnell A, Germendonk K, et al. (2007) ‘An innovative non-destructive method for the in-service metal surface temperature estimation of coated GT parts’, in Proceedings ASME Turbo Expo 2007, 14–17 May 2007, Montreal, Canada, ASME paper GT2007-27392. Sekhara A S, Prabhu B S (1998) ‘Condition monitoring of cracked rotors through transient response’, Mechanism and Machine Theory, 33, 1167–1175. Selc¸uk A, Atkinson A (2002) ‘Analysis of the Cr3+ luminescence spectra from thermally grown oxide in thermal barrier coatings’, Materials Science Engineering, A335, 147–156. Soares C (2007) Gas Turbines: A Handbook of Air, Land, and Sea Applications, Butterworth-Heinemann. Studdard B, Arrington P, Rechner M (1992) ‘Operating experience using acoustic leak detection at Gaston Station’, in Power-Gen ’92, 17–19 November 1992, Orlando, Florida. Takahashi H, Shoji T, Mao X, Hamaguchi Y, Misawa T, Saito M, Oku T, Kodaira T, Fukaya K, Nishi H, Suzuki M (1988) ‘Recommended practice for small punch (SP) testing of metallic materials’, paper JAERI-M 88–172. Tolpigo V K, Clarke D R (2000) Materials at High Temperature 17(1) 59–70. Ule B, Sustar T, Dobes F, Milika K, Bicego V, Tettamanti S, Maile K, Schwarzkopf C, Whelan M P, Kozlowski R K, Klaput J (1999) ‘Small punch test method assessment for the determination of the residual creep life of service exposed components: outcomes from an interlaboratory exercise’, Nuclear Engineering and Design, 192(1, 2), 1–11. Wilcox M, Ransom D, Platt J, Henry M (2010) ‘Engine distress detection in gas turbines with electrostatic sensors’, in Proceedings of ASME Turbo Expo 2010: Power for Land, Sea and Air, 14–18 June 2010, Glasgow, UK, paper ASME GT2010-22349. Yu Y B, Chow W K (2009) ‘Review on an advanced high-temperature measurement technology: the optical fiber thermometry’, Journal of Thermodynamics, Article ID 823482, 11 pages, DOI:10.1155/2009/823482. Zombo P J (1997) ‘Developing NDE methods for coated combustion turbine components’, Proceedings of TBC Workshop, NASA Lewis Research Center, pp. 127–137. Zombo P J and Shannon R E (2006) ‘Advanced NDE systems for flexible operation and maintenance of gas turbine components’, in Power Gen International Conference, 2006, Orlando, Florida; available from http://www.energy.siemens. com/hq/pool/hq/energy-topics/pdfs/en/service /3_Advanced _NDE_ Systems. pdf.
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3 Availability analysis of integrated gasification combined cycle (IGCC) power plants A . L A U G W I T Z , M . G R A¨ B N E R a n d B . M E Y E R , TU Bergakademie Freiberg, Germany
Abstract: In this chapter each of the four main units (air separation, gasification, gas treatment, power block) from 13 integrated gasification combined cycle (IGCC) plants are critically analyzed in terms of planned and unplanned downtimes. It can be seen that these plants cannot compete with conventional coal based power plants in terms of availability, especially due to unplanned production losses. Reasons for unplanned downtimes are many and vary from plant to plant, so a generalization concerning lessons learned is difficult to make. However, one element is that coal based integrated gasification plants need four to five years to achieve their average availability of 50–70%. On the other hand, refinery residue based integrated gasification plants need only two to three years to achieve higher average availabilities of 80–90 %. Key words: IGCC, gasification, combined cycle, power plant, ASU, availability.
3.1
Introduction
From 1994 to 1997, different coal based integrated gasification combined cycle (IGCC) plants of a size between 100 and 400 MWel were realized. The goal was to demonstrate that IGCC technologies were ready for commercial implementation. Moreover, several IGCCs, some of even bigger size, have been erected in the last 13 years, gasifying refinery residues for hydrogen and electricity production. The large-scale operation of this technology has led to the accumulation of a large store of lessons learned in the industry. Verification of low SOx, NOx or dust emissions as well as an acceptable quality of waste water and slag was achieved. Moreover, expectations concerning high thermal efficiency were confirmed. Nonetheless, ongoing improvements in conventional combustion technologies inhibited transla110 © Woodhead Publishing Limited, 2011
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tion of specific IGCC characteristics into a striking competitive advantage over these conventional power generation systems. One reason is that availability of these IGCC demonstration plants could not satisfy the demands of the power industry. Reaching commercial breakthrough will therefore also depend on investigating the reasons for low availability and conducting appropriate remedies. In the following, the main availability issues are pointed out. The objective is to minimize planned and unplanned downtimes by deriving an advantage from the lessons learned, which have been published in open source literature.
3.2
Basic structure of integrated gasification
Each IGCC comprises the same fundamental units shown in Fig. 3.1. Their particular technical realization and configuration, however, depends on specific boundary conditions. The air separation unit (ASU) is responsible for delivery of oxygen (gasification agent) and/or nitrogen (transport agent, purging gas, diluent). The gasification unit comprises fuel preparation (grinding, milling, drying, slurrying), the reactor itself and the raw gas cooling unit. Subsequently, raw gas has to be treated prior to its combustion in the combined cycle unit (CCU).
3.2.1 Air separation unit The overwhelming majority of gasifiers in IGCC applications employ oxygen of at least 85 % purity (usually 95 %) as gasification agent. The Pin˜on Pine IGCC was operated using air, but at the moment there is only the MHI gasifier at Nakoso IGCC employing air (slightly enriched) as gasification agent. The targeted oxygen purity of the other investigated IGCCs lies between 85 % and 99.5 %. Another very important and often discussed variable to be considered is the degree of air-side integration between the ASU and gas turbine, defined as the proportion of the air requirement of the ASU that is extracted from the gas turbine. Integration can reach from 0 % (Wabash) to 100 % (Buggenum, Puertollano). Air-side integration of 100 % leads to long start-up periods due to a sequential start sequence (Hannemann, 2002). In these cases the ASU can deliver required oxygen during the start-up process with a time lag only. This could be counteracted by oxygen and nitrogen storage, which would lead to increased investment costs. Additionally, the longer start-up procedure requires excessive amounts of expensive back-up fuel. So far this option seems economically unfavorable, especially in times of higher natural gas prices. On the other hand, there may be a need to extract a certain amount of air from the gas turbine air compressor in order to facilitate optimal performance over a wide range of ambient conditions and in case of
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Basic structure IGCC.
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load changes. Moreover, the auxiliary load of the ASU is markedly reduced by decreasing necessary compression effort, giving rise to higher IGCC efficiency. The ideal degree of integration depends on many conditions (e.g. chosen gas turbine, scope of environmental conditions, load changes, CO2 sequestration). However, Gra¨bner et al. (2010) recommend 70 % integration for coal fed IGCCs featuring two gas and one steam turbine.
3.2.2 Gasification unit Different options and entailed issues of the sub-units of the gasification unit strongly depend on the chosen reactor technology. Entrained flow gasifiers (GE, ConocoPhillips, Shell, Siemens, MHI and others) require coal particle sizes of less than 100 μm, whereas fluidized bed gasifiers are fed with particles of 2–5 mm in size and moving bed systems can process particles in an order of magnitude of centimeters. Of course, these differences have an influence on chosen milling technology. Moreover, it should be considered whether a dry feed (Siemens, Prenflo, KBR, MHI) or a slurry feed (GE, ConocoPhillips) is used. Usually dry feed leads to application of roller mills and the necessity of a dryer. Slurry preparation is usually done in rod mills. Besides fuel preparation, the pressurization of fuel can be conducted in different ways. Slurry fed systems employ pumps for transportation and pressurization purposes. A dry feed requires lock hoppers for pressurization and usually pneumatic dense phase conveying (using N2 or CO2 as the transport agent) for transportation purposes. The gasifier itself is typically, though not always, an entrained flow reactor. Moving bed gasifiers (like those in Vresova) are employed as well in some industrial applications, but the targeted product in these cases is chemicals or fuels rather than power. Flow direction of entrained reactors can be downdraft (GE, Siemens) or updraft (ConocoPhillips, Shell, Prenflo). The gasifier can be a refractory lined vessel (GE, ConocoPhillips) or a membrane wall construction (Shell, Siemens, MHI). Most of the reactors are single staged (GE, Shell, Siemens, Prenflo), which means that coal is fed at one level only. In order to establish a chemical quench, some gasifiers feature a second feed level. These are called two staged reactors (ConocoPhillips, MHI gasifier). Generated gases have a temperature between approximately 1000 8C (two staged gasifiers) and 1600 8C (single staged gasifiers). Prior to subsequent gas treatment, a temperature decrease is mandatory. Raw gas cooling can be realized in different ways. Usually, steam is raised while the raw gas is cooled in order to use it for power production via the steam turbine. Employed heat exchangers are radiant heat exchangers (employed only by GE, but Siemens is likely to use it (Hannemann, 2010)) and convective heat exchangers (water tube or fire tube). Another way to reduce raw gas temperature is to apply a recycle gas
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quench as performed by Shell and Prenflo. An additional option is a full water quench (offered by GE and Siemens), which is not used in coal fired IGCC plants at the moment due to a related decrease in thermal efficiency since no additional high pressure steam can be fed to the steam turbine. A full water quench is a reasonable cooling option if a downstream CO shift conversion is desired, which would be the case in chemical applications or if carbon capture and storage (CCS) is planned. Of course different combinations of mentioned cooling options are conceivable. Pre-treatment of raw gas includes separation of solid matter as well as ammonia and chlorides. Employed devices are candle filters (ceramic or sintered metal), water wash or scrubbers. Finally, slag removal devices are part of the gasification unit. Depressurization is, in most cases, done via lock hoppers. Only the ConocoPhillips gasification unit is equipped with a continuous slag depressurization system without lock hoppers (Holt, 2006). Additional information on gasifiers and related sub-units can be found in Schmalfeld (2008) and Higman (2008).
3.2.3 Acid gas removal and sulfur recovery Sulfur compounds like H2S, COS and mercury are the main impurities that have to be removed. Acid gas removal (AGR) technologies can be categorized into chemical (usually amine gas treatment, e.g. MDEA, ADIP) or physical (e.g. Selexol, Rectisol, Purisol, Genosorb) absorption processes. Additionally, there are mixed approaches like Sulfinol or Amisol combining features of both categories. Each option can be characterized as a mature process since they have been practised for many years in different industrial applications featuring very high availabilities. Rectisol is capable of guaranteeing demanding purity requirements of a possible subsequent synthesis (< 0.1 ppmv total sulfur). Moreover, it can handle a raw gas containing various trace elements. On the other hand, it is an expensive process generating sulfur levels one order of magnitude lower than required for a combined cycle. This is why IGCC plants usually employ alternative approaches like MDEA, Selexol or others, which are much cheaper. Opting for one or the other will be based on a trade-off between very high purity (Rectisol) and low investment costs (MDEA) with Selexol and Sulfinol lying in between. Moreover, COS handling needs to be considered as well. Conversion of COS into H2S via hydrolysis has to be conducted prior to cold gas cleanup, due to the fact that chemical scrubbers are not capable of absorbing COS and Selexol would be uneconomic if designed to do so. Only Rectisol does not need an upstream hydrolysis at all. Once the H2S was separated it is usually treated in a sulfur recovery unit. Typically, a Claus plant is employed for that purpose. The only exceptions in IGCC applications are Tampa and Vresova, where sulfuric acid is produced
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instead of elementary sulfur, and Nakoso IGCC, where a limestone–gypsum unit is employed. Each IGCC Claus plant employs oxygen to operate the Claus furnace in order to reduce capital costs. Oxygen demand is low compared to the oxygen consumption of the gasifier. Differences occur in the handling of tail gas treatment. In a typical refinery Claus plant the tail gas is hydrolyzed and cleaned from sulfur using an amine solution prior to being post-combusted and released into the atmosphere (traditional tail gas treating–TGT). An alternative concept is to recycle hydrogenated tail gas. This results in both elimination of a permanent sulfur emission source and increased mass flow to gas turbine. The point of recycle stream integration differs from plant to plant. Besides H2S, CO2 can be separated from the raw gas as well for capture and sequestration purposes (pre-combustion capture of CO2). The basic principle is to convert the CO of the generated gas in a water gas shift reaction into CO2 and H2 with subsequent separation of the CO2 in the gas cleaning units. Hence, a conditioned fuel gas containing mainly hydrogen can be fed to the combined cycle unit. Respecting economic considerations, suitable CO2 separation processes are selected depending on the required CO2 purity of the separated stream (selectivity), the partial pressure of CO2 in the raw gas stream and the required separation efficiency. Whereas the carbon recovery rate determines the amount of CO2 captured, the properties of the separated CO2 stream are defined by the CO2 transport and storage technology or other utilization approaches. Today, separated CO2 is used for enhanced oil recovery (EOR) in Canada. Potential technologies for CO2 separation are membranes, pressure swing adsorption (PSA) or physical and chemical absorption. Chemical scrubbers usually show a low co-absorption of non-acidic components. Thus these scrubbers show a high selectivity for CO2 if the raw gas was desulfurized in a prior step. Physical scrubbers are especially suitable for IGCC processes because CO2 separation is favorable at higher partial pressure compared to chemical scrubbers. Moreover, these systems are characterized by a higher solvent loading, which is especially advantageous if higher mass flows of CO2 have to be separated, as the amount of scrubbing agent can be reduced, resulting in a reduced auxiliary power demand for pumping. A comprehensive comparison of IGCC applications with and without CO2 separation is given by von Morstein et al. (2009) and Gra¨bner et al. (2010). However, because CO2 capture is not applied in large scale IGCC applications, its influence on the availability of a gas treatment system can be analyzed only from a theoretical point of view. It should be mentioned that the Puertollano IGCC accomplishes tests to separate CO2. Since late 2010, a new 14 MW pilot plant has demonstrated H2 production and CO2 capture (ELCOGAS, 2010).
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3.2.4 Combined cycle The combined cycle unit includes a gas turbine, a heat recovery steam cycle and associated auxiliary devices and is comparable to a natural gas driven combined cycle. One main difference is the application of NOx reduction techniques. Operation with significant quantities of hydrogen is not possible with the pre-mix ‘dry low NOx’ burners used in modern natural gas machines and so diffusion flame burners are used. NOx control is therefore performed by dilution either with N2 or steam or a combination of both. There is little experience with selective catalytic reduction (SCR) in the flue gas stream. Waste water treatment and the flare do not significantly contribute to overall downtime of an IGCC plant and will not be regarded in the following. An overview and detailed information on state-of-the art IGCC plants can be found in Holt (2006), Anon. (2010) and Nykomb-Synergetics (2007). Tables 3.1 to 3.3 summarize the most important characteristics of analyzed IGCC plants.
3.3
Availability issues of the ASU
Industrial gasification plant operators have two different options for oxygen supply: onsite or offsite production. Typically, coal based IGCC plants chose the onsite option. On the other hand, most refinery IGCCs opted for the offsite oxygen generation. In this way a gas company can operate the ASU by employing specially trained staff, which usually leads to better reliability of oxygen supply in these cases. Generally, air separation units in industrial applications have a typical availability of 98.5 % with downtimes of approximately 1 % due to planned outages and 0.5 % due to unplanned outages respectively (EPRI, 2007a). Even though planned downtimes of an IGCC-ASU are scheduled to coincide with other planned maintenance work, the performance of existing units has not matched that of other industrial ASU performances, as shown below.
3.3.1 Reported failures Even though there are no published numbers, it seems quite evident, though surprising, that the electric motor of the air compressor was a main reason for poor availability (< 70 % within the first 2 years) of the Falconara IGCC plant. Subsequently, the unreliable electric motor was replaced. It is believed that lack of reliability was caused by the employed cooling concept. Fresh sea water was used for cooling purposes instead of a closed circuit. This design caused considerable corrosion problems and entailed leakages
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Nippon Oil Corp.
Clean Coal Power R&D Co. LtD
Falconara, ITA
Negishi, JAP
Nakoso, JAP
SGC – synthesis gas cooler.
api
Sarlux, ITA
*
ERG Petroli; LUKOIL Saras S.p.A.
ISAB, ITA
2007
2003
2001
2001
2000
1997
Shell
Pernis, NL
1997
1997
Sustec
SVZ, GER
1996
250gross
350
241
550
512
130gross
300
80
26200
100 253
Puertollano, ESP ELCOGAS
Sokolovska Uhelna
Vresova, CZ
250
1998 1994
250
1996
SG Solutions/ 1995 ConocoPhillips
Wabash River, USA
Polk Power Tampa Station, USA Electric Pin˜on Pine, USA SPPC Buggenum, NL Nuon (part of Vattenfall)
Operator
Plant
28
22 26
30
29
71 N.A.
MHI; air blown
66
39
GE (oil)
GE (oil)
GE (oil)
68
61
Shell (oil) GE (oil)
26
Prenflo
Coal
Power
Power
Chemical quench; convective SGC
McDaniel, 2002
Dowd, 2000
References (for Tables 3.1, 3.2 and 3.3)
Breakler and Kamka, 2004 Holt, 2006 Siemens V94.3 ELCOGAS, 2001 Pen˜a, 2009 Pen˜a, 2010 GE MS 6541 B de Graaf et al., 1999 EPRI, 2007a 26Siemens Anon. 2009 V94.2K EPRI, 2007a 36GE 9E Sharp et al., 2002 Collodi and Brkic, 2003 Alstom (ABB) Sharp et al., 2002 13E2 Arienti et al., 2005 MHI 701F Ono, 2003 Yamaguchi, 2004 Hatayama, 2006 MHI 701 DA Ishibashi and Shinada, 2008 Ishibashi, 2009 Watanabe, 2010
GE 6 B
GE 6 FA NETL, 2002 Siemens V94.2 Hannemann, 2002 Kanaar, 2002 Holt, 2006 26GE 9E Buryan et al., 2008 Holt, 2006
GE 7 FA
GE 7 FA
Cooling concept Gas turbine
Chemical quench; fire tube SGC* Coal; Power Radiant SGC; petcoke fire tube SGC Coal Power Convective SGC Coal; Power Recycle gas biomass quench; water tube SGC Lignite Power Full water quench for Siemens– gasifier–raw gases Coal; Methanol; Diverse biomass; power waste; plastics Coal Power Recycle gas quench; water tube SGC Refinery H2; power Fire tube SGC residues Refinery H2; power Full water quench residues Refinery H2; power Full water quench residues Refinery H2; power Full water quench residues Refinery Power Full water residues quench
Coal; petcoke
Pressure (bar) Feed-stock Product
76Lurgi fixed bed 25 16BGL 16Siemens
266Lurgi fixed bed 16Siemens
KRW; air blown Shell (coal)
GE (coal)
ConocoPhillips
Size Start-up (MWel,net) Gasifier
Table 3.1 Overview of analyzed IGCC plants
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Table 3.2
Overview gas treatment and sulfur recovery in gasification plants
Plant
Gas treatment
Sulfur recovery
Wabash River Polk Power Station Pin˜on Pine Buggenum
Claus unit Sulfuric acid CaO + SO2 + 12O2 → CaSO4 Claus unit
Vresova
COS hydrolysis + MDEA COS hydrolysis + MDEA Transport desulfurizer HCN/COS-hydrolysis + Sulfinol M Rectisol
SVZ Puertollano Pernis ISAB Sarlux Falconara Negishi Nakoso
CO shift + Rectisol COS hydrolysis + MDEA CO shift + Rectisol COS hydrolysis + MDEA COS hydrolysis + Selexol COS hydrolysis + Selexol COS/HCN hydrolysis + ADIP COS hydrolysis + MDEA
Table 3.3
Summary IGCC ASU
Plant
LOX Capacity % O2 % Integration O2 production storage ASU–GT (t/d)
Wabash River 2060 Polk Power 1840 Station Pin˜on Pine Buggenum Vresova SVZ Puertollano Pernis
1780 N.A. N.A. 2400 3175
95 95
0 Initially 0; 15 since 2005
95 100 96 0 94 50 85 100 99.5 0
ISAB 261850 95 Sarlux 262300 95 api Falconara 95 Negishi 2400 95 Nakoso
Haldor Topsøe (wet sulfuric acid plant) — Claus unit Claus unit Claus unit Claus unit Claus unit Claus unit Limestone–gypsum unit
0 0 0 0
Onsite Onsite
Air blown Onsite Onsite Offsite Onsite Offsite
Offsite Offsite Onsite Onsite
Supplier
No No
Air Liquide Air Products
Yes N.A. N.A. Yes Yes, plus 7 km pipeline Yes Yes No No
Air Products Linde Linde Air Liquide Air Products
Air Liquide Air Liquide Praxair Air Liquide
Enriched air blown gasifier
(Arienti et al., 2005). Likewise problems occurred at Wabash. Windings of the electric motor were exposed to moisture, which caused serious problems but no outages. The root of related troubles is that the motor operated outdoors even though it was not designed for outdoor operation. Tampa and Negishi suffered from production loss of 3 weeks (McDaniel, 2001; Yamaguchi, 2004), which was caused by a damaged air compressor rotor and required repair at the manufacturers’ factories. Neither Tampa nor
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Negishi had a backup rotor, which could have reduced downtime. According to the Final Technical Report from the Wabash River Project (Dowd, 2000), the most important issue concerning reliability of its air compressors is the inlet guide vane. Furthermore, Tampa operators experienced some trouble with that particular device and there were several other gasification plants, which had additional problems with instrumentation and control devices associated with the air compressor. In some cases, molecular sieve heat exchangers suffered from leakage of steam to the gas side, which can lead to intrusion of moisture into the molecular sieve during regeneration. As a consequence, the cold box might be exposed to moisture and subsequently freeze up. One possible reason for leakage was reported to be vibration of the heat exchanger. Furthermore, ductwork leakage within the cold box was experienced at Tampa and Wabash. The 401 hours of unplanned downtime in 2000 were caused by welding failures on an instrument tube at Tampa (McDaniel, 2000). Repair took quite a long time because an inaccessible area was concerned. The Wabash River Final Technical Report (Dowd, 2000) reports 299 hours of unplanned downtime due to inaccurate welding on a de-icing collector as well. During a durability test of the recently erected Nakoso plant (air blown MHI gasifier) leakage of extraction air cooler tubes was observed. This was caused by corrosion due to inadequate and irregular tube materials (Watanabe, 2010). None of these failures can be attributed to design features specific to an IGCC application. However, when designing an ASU for an IGCC, it is clearly necessary that similar criteria are applied to issues of reliabilityrelated design and specification as elsewhere in the industrial gas industry in order to ensure similar reliabilities of 99–99.5 %.
3.3.2 Liquid oxygen storage Short-term outages of ASU can be bridged by buffered oxygen, though this is limited for economic reasons to about half a day. Regarding the above explained failures and root causes, it is evident that LOX storage would not have been capable to avoid all production losses. These considerations caused some owners not to consider oxygen storage in future gasification plants (ELCOGAS, 2001). On the other hand, the presence of a storage system seems reasonable for certain plants, as operators make some effort to improve their systems (Buggenum). Several considerations should be taken into account when opting for or against liquid oxygen storage. First, due to the fact that motors from air compressors are vulnerable to voltage dips, stability of the external grid must be ensured. Second, economical penalties for a gasifier outage (mainly backup fuel for the gas turbine and restart process) need to be considered.
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3.4
Availability issues of the gasification unit
Due to the fact that the gasification unit is the most distinctive part of an IGCC plant, its sub-units are itemized in more detail compared to the other main units.
3.4.1 Feed system Issues for slurry feed and dry feed systems tend to be different. One main issue related to the slurry feed system is settlement of particulate matter in both storage tanks and suction pipes (especially upstream of the slurry pump during downtimes). This can be avoided by ensuring permanent motion of the slurry. Most slurry fed IGCCs employ slurry storage tanks, which can bridge unplanned downtime of the rod mills. Moreover, most plants (except Tampa) have two slurry pumps of 50–100 % capacity each. A similar decision is necessary for dry fed systems. The number and capacity of mills have a decisive influence on the availability of the fuel preparation unit. The 2660 % roller mills from Puertollano have been identified as being insufficiently robust (Pen˜a, 2005) and to be important contributors to outages. The 3650 % mills used in Buggenum ensure increased availability of the fuel preparation unit. Dry fed IGCCs are of course concerned with common issues related to transport and storage of ground coal, e.g. bridging in sluicing devices or clogging of conveyors as experienced at Puertollano. Moreover, it is not trivial to maintain a stable fluidization and adequate pressure control in dense phase transport systems and to establish a vital and reasonable coal dust explosion prevention system. An issue common to both feed systems is the proper blending of raw materials in order to guarantee adequate and predictable characteristics of the gasifier feedstock. This is of decisive importance, since adjustment of gasification conditions, slag removal and raw gas cooling parameters is based on predicted feedstock properties.
3.4.2 Gasifier Feedstock It is absolutely necessary to avoid slag solidification within the reaction chamber of slagging gasifiers. Solid slag lumps would plug the bottom discharge system. Thus it has to be ensured that operation conditions are constantly above the ash melting point. In the case of altered coal feed it is essential to adapt the operation temperature to new ash properties. As an example, Buggenum suffered from formation of slag lumps in the
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lower part of the reaction chamber in the early days. Until 2002 the IGCC operation management was not responsible for delivered coal quality. Coal was supplied by nearby conventional plants. In this way, uncontrolled changes in coal quality caused several outages (Kanaar, 2002). Reactor wall Decisions made on refractory lined vessel or cooling jacket concepts or membrane wall construction have a decisive impact on unplanned and especially planned downtime. Aggressive and corrosive slag stresses different layers of refractory lining, which leads to increased abrasion and wear. Consequently, the lifetime of such vessel linings is usually less than 2 years, even though Wabash achieved 3 years by 2006 (Holt, 2006). In case of replacement of refractory it takes up to 35 days to cool down the gasifier, change the refractory and reheat the system. On the other hand, Payonk (2008) reports that refractory repairs at the Wabash River IGCC can be done in 17 days. If possible, it is reasonable to make these changes during planned maintenance of other devices (e.g. the gas turbine). Moreover, employment of a refractory lining requires slow start-up and shut-down to avoid high temperature gradients in the bricks. Comparatively short lifetimes of these systems make employment of a spare reactor especially desirable even though Payonk (2008) claims that the spare reactor at the Wabash River IGCC is no longer needed. An alternative design is provided by membrane wall construction. The metallic vessel is protected by a water wall with a studded ramming mass. Pressurized water flows through the tubes of the wall, is heated and is used for steam generation. The inner side of the ramming mass is exposed to the hot synthesis gas. Liquid slag reaching the inner wall of the reactor transfers its heat into the membrane wall and thus solidifies. A protective layer of solid slag grows. If this insulating coat is thick enough, the outer slag will no longer solidify and will flow downwards into a slag chamber. The insulating layer regenerates immediately whenever a part of the wall is unprotected. This construction makes it absolutely necessary to operate at temperatures above the ash melting point. When gasifying low ash feedstock (pet coke, low ash coals) about 5 % of the ash or flux must be added to the feed, so as to maintain the solid slag layer. Advantages compared to refractory lined vessels are reduced weight of the gasifier, faster start-up and faster shutdown because temperature gradients are unproblematic. Lifetime of the Buggenum membrane wall is expected to be 25 years with more than 16 years of successful operation so far (Chhoa, 2005). Prenflo and Siemens gasifiers expect comparable lifetimes (Radtke, 2007; Schingnitz, 2008). Thus, the lifetime of membrane walls does not influence the overall availability. However, it should be mentioned that the bottom part of the Buggenum
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reactor was refractory lined and required periodic maintenance. It is expected that Shell will equip future gasifiers with a water cooled bottom region (Chhoa, 2005). On the other hand, there are some issues related to membrane walls at Puertollano, where the wall suffered from local leakage caused by flow blockages and erosion on the water side (Pena˜, 2005). Burner Slurry feed systems are stressed by frequent burner changes because stress corrosion cracking leads to short lifetimes of injector tips (Holt, 2006). Associated downtime is approximately half a day. In the case of single train concepts, downtime may be even longer because downstream gas cleaning units have to be restarted after the maintenance period, which is timeconsuming. Lifetime of burners from dry fed gasifiers (up to more than 16 000 hours) does not affect overall availability at all (Chhoa, 2005). However, for the dry feed gasifier at Nakoso, Watanabe reported an erosion issue at the burner cooling tube caused by inadequate positioning of the burner (Watanabe, 2010). Corrosion Pen˜a (2005) reports of ‘cold ends’ corrosion and downtime corrosion at the Puertollano IGCC. Characteristic locations are fittings for the measurement equipment and nozzles, where wet syngas is cooled below its dewpoint. There are two ways to eliminate such destructive corrosion. Either utilization of sour gas corrosion resistant materials or deployment of trace heating should be considered.
3.4.3 Slag removal Issues connected with the slag removal section have centered around plugging and erosion/corrosion aspects. Usually, plugging emerged if slag properties changed due to an unexpected change in feedstock properties as described above. Polk Power Station experienced how solidified slag lumps plugged the slag discharge line as well. It takes up to 10 days to cool down the system and to mine out the clump (Holt, 2006). An important issue in the slag removal unit is erosion and corrosion in the piping of circulating slag water from the slag bath. This circulating water contains fine, sharp solids, which cause erosion (Holt, 2006). Additionally, the pH value needs to be monitored in order to avoid corrosion. As reported by EPRI (2007a), the Buggenum slag disposal tubes were stressed by erosion and corrosion, leading to approximately 120 hours of downtime in 2002. Meanwhile, these tubes have been replaced by more resistant duplex steel. It
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has also been reported that particle separation from slag bath water caused problems. Thus it was necessary to retrofit a lamella separator. It is well known that Buggenum decided to replace the ceramic heat skirt, which was installed to protect the pressure vessel wall of the slag bath, and retrofitted a membrane wall construction. Puertollano operators observed that erosion will occur if velocities in solids separation units (e.g. slag removal, entrained fly ash) are high. As in Buggenum, stressed devices were replaced by wearresistant materials and design. Also, operation control has been improved (Pen˜a, 2005). A unique pressure let-down system is used in ConocoPhillips systems (Wabash) where lock hoppers are redundant. So far, no problems have been reported, apart from slag plugging, which emerged once.
3.4.4 Raw gas cooling Fouling, plugging and leakage are the possible reasons for a significant number of failures. Unplanned maintenance is reported as a regular issue by each IGCC operator, though reasonss vary due to the different cooling approaches taken by the respective technology suppliers. Alteration of the feedstock quality can cause entrained fly ash to plug the entry zone of cooling units in a comparable manner to the plugging of the slag discharge line as described earlier. Adapted operation control is required in these cases, as done by Buggenum operators, which consequently could reduce such fouling issues (Eurlings and Ploeg, 1999). On the other hand, leakage has remained an issue in the last few years. Waste heat boilers in Puertollano were stressed by both sticky and fluffy fly ash (Pena˜, 2005). As a consequence, the recycle gas stream (for gas quench purposes) was increased in order to decrease the gas inlet temperature at the entry zone of the heat exchanger. In this way entrained ash was much less sticky and both plugging and fouling were reduced. Issues contributing to fluffy fly ash were solved by increasing gas velocity. The gasification system at Tampa employs both a radiant syngas cooler (RSG) and convective syngas coolers (CSCs). The RSG is much less stressed by fouling than expected; thus it cools the raw gas even more than designed. Only a few outages of 288 hours in 1999 and 96 hours in 2001 were reported by EPRI (2007a). The inlet region of the CSCs is one main problem and significantly reduces overall availability. Ten outages between 1999 and 2001 caused almost 1500 hours of unavailability. As a consequence, accessibility of the heat exchangers was improved and it was decided to eliminate convective coolers for the future reference plant (Rigdon and Schmoe, 2005). Even after more experience was gained this issue could not be overcome. Up to 5 % lost availability was caused by plugged coolers in 2006
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and 2007 (McDaniel, 2007). Plugging forces a shut-down in order to clean the tubesheets every 30 to 45 days (McDaniel, 2008). Because the ConocoPhillips gasifier features a second stage (chemical quench leads to comparatively low outlet temperatures of approximately 1040 8C), there is only one vertically arranged convective syngas cooler employed at Wabash River. Fouling was also an early problem of this syngas cooler. It was caused by entrained tar matter that survived in the secondary gasification zone. Adapted operation control and feeding of petcoke solved that problem (Dowd, 2000). An additional correction was to improve accessibility of the syngas cooler, as was done at Tampa as well. Another, very specific incident resulted in 607 hours of lost production. Sodium carbonate condensed at cooling pipes, which led to an increased pressure drop and impeded heat exchange. The source of sodium might have been waste water recycled to the rod mills that was neutralized using caustic soda (EPRI, 2007a).
3.4.5 Raw gas pre-cleaning Dry filtration Scheibner and Wolters (2002) report of a successful cooperation between Nuon and Pall Schumacher to improve performance of ceramic candle filters at the Buggenum IGCC. Lifetime could be doubled to two years because plugging and breakage of candles were overcome. One remedy was to fix candles in a grid in order to avoid horizontal movements of the candles. Unfortunately, Puertollano could not extend the lifetime of employed ceramic candle filters above 4000 operating hours (8000 hours were initially expected), because plugging had not yet been overcome. Another issue related to ceramic filters is breakage. This is why Wabash operators switched to sintered metal filters in 1996. These show a reasonable performance and lifetimes of 10 000 hours (Payonk, 2008). Watanabe (2010) also reports inadequate tightening of a package. This resulted in gas leakage from the base of a rotary valve below the filter. The failed IGCC project Pin˜on Pine had its main problems with a hot gas filtration system. As reported in a DOE Assessment Paper (NETL, 2002), operators could not overcome problems with the filter-fines removal system (accumulation of fines in the bottom region of the hot gas filter). During a last start-up attempt of the Pin˜on Pine IGCC, some material from the desulfurizer (installed just prior to the hot gas filter) was entrained, reached the filter and ignited. The consequential fire and ensuring damage caused the termination of the project in 2001.
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Wet scrubbing The whole scrubbing unit and recycle system is acidic. Therefore, monitoring and adjustment of the pH value is necessary. If entrained particles are not separated in a dry filter (Tampa), they will be separated in the scrubber and the correct design is required in order to avoid erosion and corrosion issues in the black water piping. Typically fivedimensional radius bends and ensuring a low chloride content in the circulating water are usually sufficient to avoid any concern. Attention should be paid to another potential problem as well. Raw gases could entrain droplets from the scrubbing unit, especially during the startup process. Droplets contain halides and metals, which act as a catalyst poison in the subsequent COS hydrolysis. In order to avoid damaged catalysts, as experienced at Tampa (McDaniel, 2002), the COS hydrolysis should be started only if the prior scrubber has reached stable operation. Although the gasification area has not been the most frequent contributor to unplanned outages, some issues are less amenable to improvement than those in other parts of the plant. Some problems like issues with the convective syngas coolers in Polk can only be resolved in a new plant and the demonstration plant will have to continue with the design it has. Work has been done on extending the life of refractory linings, although it will be a while before the improvements can be quantified. There are inherent issues with any solids handling systems and the gasification unit is no exception. For dry-feed units some solutions are available from conventional power plant technology. Clearly the lessons learned from the demonstration units will benefit new plants, but one must still be cautious about taking too much credit for this at this time.
3.5
Availability issues of acid gas removal (AGR) and sulfur recovery
Even though AGR is a mature technology and only a small amount of unplanned downtime can be attributed to the related units, there are some lessons to be learned.
3.5.1 Acid gas removal COS hydrolysis Several operators reported the formation of formic acid on hydrolysis catalysts. This acid can give rise to corrosion issues (directly or indirectly) for the downstream units. Formation may be due to the following reaction (EPRI, 2007a): HCN þ 2H2 O $ NH3 þ CHOOH;
DH ¼ 76:14 kJ=mol
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The only plants that seem to be untroubled by that issue are Negishi (they have an ammonia wash prior to their ADIP desulfurization unit) and Buggenum (they explicitly call their hydrolysis a HCN/COS hydrolysis). Moreover, catalyst poisoning and damage caused by condensation must be avoided. For this reason it needs to be ensured that alkali, heavy metals and halides do not reach the COS catalysts and that operation is not close to dewpoint conditions. Tampa operators had to change their start-up procedure to avoid entrainment of catalyst poisons from the scrubber into the COS hydrolysis unit (McDaniel, 2002). There was no water wash prior to the COS hydrolysis unit in the original design for Wabash, so chlorides and metals damaged the COS catalyst. After the first damaged catalyst a water wash was retrofitted (Dowd, 2000). The first catalysts at Puertollano were alumina based and featured a very poor durability leading to 2–3 changes p.a. (Pen˜a, 2005). Later, a new and 8 times more expensive catalyst, based on titanium, was employed. Additionally, inlet temperature was increased and subsequently the point of operation was shifted further away from the dewpoint. Both remedies increased the lifetime of the catalyst, which operated successfully for 4 years. Selexol Operators of Selexol utilizing plants reported the entrainment of formic acid from the COS hydrolysis unit during the start-up process. They discovered that formic acid is absorbed by Selexol solvent and is vaporized in the regenerator. Subsequently, it condenses overhead, is refluxed back into the column and accumulates in the system in this way. Consequently, the pH value is reduced, which leads to considerable corrosion (Sharp et al., 2002). Rectisol There is one issue related to Rectisol that causes significant downtimes: plugging of the methanol–water column, as reported, for example, by Shell for its Pernis IGCC (EPRI, 2007a). MDEA Entrainment of formic acid from COS hydrolysis caused significant degradation of MDEA solution. This caused the formation of corrosive heat stable salts. As a consequence, additional maintenance was necessary to fix corroded and fouled devices. Installation of an ion exchanger solved that issue at Tampa and Wabash by removing heat stable salts from the solution cycle (Dowd, 2000; McDaniel, 2002). Since the new titanium catalyst was
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employed, Puertollano operators had no further trouble with fouling of the AGR solvent.
3.5.2 Sulfur recovery Even though Claus plants in refinery and natural gas environments are extremely reliable, they represent a critical bottleneck as these units are in charge of assuring permitted annual sulfur emission of the whole plant. In the case of a Claus plant outage, these emission values will be reached quickly, forcing a shut-down of the IGCC. As this circumstance is generally recognized, most commercially operated IGCCs employ 2666 % Claus units. Since in general the ICGG will not be operating with the highest sulfur design coal, this still leaves scope for an operation reasonably close to full load in the event that one unit is down. However, there are some reported hours of lost production connected to the Claus unit. Wolters (2003) reports of almost 600 hours of downtime in 2003 due to corroded welds and entailed leakage of the Buggenum Claus unit. The remedy was found by redesigning the tube sheets. Acid gas removal and sulfur recovery units have not been a major contributor to unplanned outages in IGCCs. Nonetheless, their performance in the demonstration units has not been as good as similar equipment has shown in the oil and gas or chemical industries. Many of the problems can be attributed to issues of plant integration, e.g. trace formic acid formation in the COS hydrolysis causing corrosion in the downstream acid gas removal unit. These issues have now been largely identified and a closer approach to the reliability levels of other industries can be expected.
3.6
Availability issues of the combined cycle
During examination of data provided by public reports and databases (ORAP®), in order to judge the availability of syngas fired combined cycles it turns out that these units cannot match the reliability performance of natural gas driven combined cycles. Natural gas driven gas turbines achieve availabilities of 90–95 % whereas CC units in IGCC applications achieve in most cases less than 90 % (Higman, 2005). Analyzing these data, one should be aware of the fact that Tampa, Wabash and Puertollano employed new F-class gas turbines (see Table 3.1). Each represents an early model of the respective series. Even natural gas driven turbines of the same class exhibited availability-reducing problems in the early years and were subject to a major product recall action. The assumption that lower availability is not caused by the syngas application itself is supported by the fact that Italian gasification plants and Buggenum have been equipped with E-class gas turbines, showing a much higher
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availability. However, some published reasons for unplanned downtimes are listed below. Tampa and Wabash operators experienced quite similar problems with their CCUs. An outage lasting 100 days in 1999 due to a damaged blade in the compressor section of the gas turbine is mentioned in the Final Technical Report from the Wabash IGCC (Dowd, 2000). In 2005, a damaged compressor also caused an unplanned outage of 100 days in Tampa. The damage was caused by a creeping compressor casing and accompanied breakage of the compressor’s stationary elements. Moreover, Wabash suffered from an outage of 19 days due to pipe leakages in HRSG (Payonk, 2000). Tampa experienced problems with the backup fuel system, leading to 416 hours of lost production in 2000 (McDaniel, 2000). Additional 24 days of unplanned outage was reported for 2003. The 7FA rotor was replaced and both generators were rewound. Horizontal silo burners caused several problems at the Puertollano gas turbine. Before a burner modification was made in 2003, preventive inspection of the hot gas pass was required every 500–1000 operation hours (a 3 year interval was desired). After improvement the inspection interval could be extended to 4000 hours. Me´ndez-Vigo (2002) reports of 617 hours of downtime in 2001 and 105 hours in 2002. These were necessary, among others, to revise and change ceramic tiles in gas turbine combustion chambers. In 2003 a production loss lasting three months was caused by deformation of the generator casing flange. In 2004 and 2005, issues with the gas turbine transformer, DCS and compressor guide vane occurred. Several operators reported difficulties caused by humming and oscillating gas turbine burners. One of the latest incidents attributed to CCU occurred at the ISAB IGCC. It was reported that an explosion damaged one train of the combined cycle, in October 2008. Since December 2008, the plant has run on the left train. The accident happened during maintenance activities of one train. For some reason, syngas reached the non-operating train past the gas turbine and ignited as it contacted with air in the subsequent HRSG. The explosion pushed back into the gas turbine and destroyed the train. It seems that inertization (using nitrogen) of the maintained train was insufficient (Anon., 2009). Another explosion is reported for the SVZ IGCC by Braekler and Kamka (2004). At SVZ, light oil is fed to the gas turbine for start-up purposes. The explosion occurred during test runs while switching from light oil to syngas. As a consequence, the gas skid was retrofitted by a nitrogen buffer. In the demonstration plants the combined cycle unit has statistically been the largest source of unplanned outage. In evaluating this finding, it is necessary to separate issues associated with syngas firing and those due to
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other reasons, which would also impact availability of a natural gas-fired turbine. With the exception of the Puertollano machine (which is no longer offered by the vendor), the syngas firing issues have been resolved and provide a sound basis for implementing the conversion of newer natural gas models to syngas firing. Nonetheless, it may be wise to assume that a syngas fired turbine will not have an above-average availability compared to a natural gas turbine, or at least not initially.
3.7
Summary of existing plants
Four of the plants (Buggenum, Puertollano, Tampa and Wabash) reviewed in the previous sections were built in the mid-1990s as demonstration plants. For the gasification section three of the four plants represented a major upscaling on previous references. For Puertollano the factor was 40, for Buggenum 9 and for Tampa 2. Additionally, the gas turbines from Tampa, Wabash and Puertollano were early examples of their respective technologies (7FA and V94.3), both of which required several years before fleet teething problems in natural gas service had been resolved. Nevertheless, all the plants achieved their targets for high efficiency and low emissions. While they have also made the transition from demonstration to commercial service, the availability has been less than satisfactory. The following figures summarize the contributions of different units to overall downtime (each reflecting the latest data that are available from public domain sources) and development of IGCC availability of these IGCCs over operation years. It is important to note that published data concerning availability are edited in different ways. Some authors, for example, report on reliability instead of availability or it is not mentioned if stated figures are based on syngas operation only or include backup fuel operation as well. Additionally, one has to be aware of the fact that the contribution of specific main units to overall downtime differs from year to year, as shown in Fig. 3.2 for Puertollano. This is caused by several planned maintenance activities, which consume much time but are necessary in certain years only. It would therefore be best practice to take the mean value of all operational years to calculate representative numbers of their contribution to overall downtime. Unfortunately, there is a lack of published data for most IGCC plants and therefore Fig. 3.3 refers to a specific year only, rather than presenting an average of several years. Nevertheless, general trends can be derived. Additionally, Fig. 3.3 includes coal based IGCCs only, due to the fact that this information is not reported for refinery based IGCCs in open source literature. Note that the Wabash River IGCC was not in operation in 2004 and most of 2003 because of contract restructuring and ownership change (Payonk, 2008). Therefore availability from these years is missing in Fig. 3.4.
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3.2 Main units downtime at Puertollano (2006–2009) (Pen˜a, 2007, 2008, 2009, 2010).
3.3 Contribution of different main units to total downtimes (numbers in brackets refer to the year of operation when numbers were recorded) (Wolters, 2003; McDaniel, 2008; Payonk, 2008; Pen˜a, 2010).
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3.4 Development of IGCC availability (data adapted from Zuideveld, 2001, 2002; Kanaar and von Dongen, 2006; EPRI, 2007a; Hornick, 2007; Ishibashi and Shinada, 2008; Payonk, 2008; Ciccotosto et al., 2009; Domenichini, 2009; Ishibashi, 2009; Pen˜a, 2009, 2010).
It is generally found that refinery based IGCCs show higher availabilities compared to coal based plants. One remarkable fact about availability of coal based IGCCs is the low reliability of ASUs and especially of the combined cycle. The gasification unit itself contributes less to overall production loss than the combined cycle. A fair amount of ASU downtimes could have been avoided and are avoided, as a matter of fact, in the industrial gas industry. Nevertheless, the explained reasons for unplanned downtimes show that poor availability of ASU is not a consequence of its application in an IGCC. Analyzing outages related to the gasifier itself, as done in section 3.4, shows that another conclusion has to be drawn. Even though some lessons learned could have been anticipated, like settlement of solid matter in slurry tanks or downtime corrosion, there are some difficulties left. Extending lifetimes of the refractory lining and injector tips (for slurry feed gasifiers) continue to be the subject of research and results will also benefit existing plants. Other issues, especially fouling in the entry zone of the convective syngas cooler in Polk, remain of concern in these plants, but have been addressed in the design of newer plants. Future research activities will continue to address these problems in order to achieve availabilities comparable to those of Italian refinery residue gasification units. In the case
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of refractory lined reactors, planned maintenance (lining replacement) and its timing with inspections of gas turbines is essential to maximize availability as these two tasks are on the critical path. It is useful to address raw gas treatment units as well, even though reported downtimes hardly contribute to overall production loss. One lesson learned is that entrainment of formic acid from COS hydrolysis to Selexol or MDEA units causes degradation. A second valuable consideration is employment of 2660 % Claus plants due to the fact that availability of the Claus unit is critical to overall plant availability. Low availability of the combined cycle is more a question of the gas turbine applied and the reliability of its auxiliary units rather than application of syngas instead of natural gas. Increasing availability of that unit seems probable for future IGCC plants. However, appropriate scheduling and timing of planned maintenance and inspection activities are of great importance concerning availability, as shown below. Finally, activities to establish a maintenance program for IGCC plants are comparable to those of conventional power plants. A detailed development of a maintenance program for an IGCC plant including, besides others, reliability, availability, maintainability (RAM), hazard and operability (HAZOP), safety integrity level (SIL), failure mode effect and critical analysis (FMECA) and critical spare parts analysis can be found in Verwilligen (2010).
3.8
Forecast based on RAM modeling
A reliability, availability and maintainability analysis (RAM) can be carried out if a database for the various components of an IGCC plant is available comprising mean values and statistical deviations for times between failures and time to repair. Since some international studies incorporate this kind of data, which is typically not gained at IGCC plants, a plausibility check with operating plants is required. Simulations were accomplished by different authors (EPRI, 2007b; Higman, 2008; Sutor et al., 2009) yielding theoretical availabilities and reliability factors for different plant configurations on a higher level of detail. Table 3.4 shows five cases comprising two common IGCC plants with Shell (case A) and GE gasification technology (case B) (EPRI, 2007b), with the latter including a spare gasifier. Case C represents a standard IGCC for hard coal with Shell gasification and CO2 capture derived from von Morstein et al. (2009). The last two cases, D and E, were recently published by Sutor et al. (2009), exhibiting a polygeneration IGCC plant that can produce methane as well as electric power, hence including a methanation unit.
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98.84%
n Cap. RF.
Case B – GE IGCC with spare gasifier (EPRI, 2007b)
97.85%
Fly ash removal
89.36% 5.11% 84.80%
91.02% 5.38% 86.12%
1 100% 97.28% Methanation
1 100% 99.51% Methanation
2 60% 97.81%
2 100% 99.06%
2 60% 99.52%
2 60% 99.23%
95.53%
91.45% 5.11% 86.77%
1 100% 99.50%
85.49%
1 100% 98.76%
87.89%
2 60% 99.00%
1 100% 99.80% CO2 compression 1 100% 99.80% CO2 compression 2 60% 99.84% and dehumidification and dehumidification 1 100% 97.87% 1 100% 96.66% 1 100% 97.02%
1 100% 98.83% Claus process
1 100% 99.40% Acid gas removal
1 100% 99.04% CO shift conversion
94.46%
1 100% 99.55% Claus process 1 100% 99.28%
1 100% 99.66%
1 100% 99.74% Acid gas removal
1 100% 99.77% 1 100% 99.74% CO shift conversion
97.08%
1 100% 99.28%
n: number of units, Cap.: capacity, RF: reliability factor; AF: availability factor.
Total Reliability Factor Scheduled Outage Factor Equivalent Availability
2 50% 99.16%
2 50% 99.95% Water wash 2 50% 99.90% Gas cooling, 2 50% 99.88% CO shift boiler feed water, conversion and knockout 2 50% 99.95% Mercury removal 2 50% 99.95% Rectisol absorption 2 50% 99.55% COS hydrolysis 2 50% 99.55% Rectisol regeneration 1 100% 99.87% Selexol - acid 2 50% 99.71% gas removal 2 66% 99.36% Claus process 2 66% 99.36% Claus process 1 100% 99.50% Hydrogenation 1 100% 99.50% Hydrogenation reactor and reactor and gas cooler gas cooler 2 50% 99.87% Syngas reheat 2 50% 99.87% Syngas humidification and reheat CO2 compression and dehumidification 1 100% 94.91% 1 100% 94.91%
97.98%
1 100% 99.74%
4 25% 96.16%
3 60% 99.61%
95.78%
n Cap. AF.
Case E – Polygeneration IGCC – Full Redundant (Sutor et al., 2009)
2 50% 98.59% Coal milling and drying 4 25% 96.16% Siemens gasifier full water quench
94.80%
n Cap. AF.
Case D – Polygeneration IGCC – Non redundant (Sutor et al., 2009)
2 100% 99.78% Coal milling and drying 1 100% 97.90% Siemens gasifier full water quench
96.73%
n Cap. RF.
Case C – Shell CCS IGCC (Higman, 2008)
2 100% 100.00% Coal handling, 2 100% 99.99% conveyor 3 50% 99.75% Slurry preparation 2 70% 99.36% Coal milling and drying 2 60% 99.54% Slurry feeding 3 50% 99.91% Shell gasifier, syngas cooling, cyclone, filter 2 50% 98.25% GE gasifier, 3 50% 99.58% Syngas cooling cooler, scrubber
97.55%
n Cap. RF.
Air Separation Unit 1 100% 98.52%
Combined Cycle
Sulfinol – acid gas removal Claus process Hydrogenation reactor and gas cooler Syngas humidification and reheat
COS hydrolysis
Water wash Gas cooling, boiler feed water and knockout Mercury removal
Gas Treatment
Shell gasifier, syngas cooling, cyclone, filter
Coal handling, conveyor Coal milling and drying Coal feeding
Gas Production
Case A – Shell IGCC (EPRI, 2007b)
Table 3.4 Results of RAM studies
Availability analysis of IGCC power plants 133
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All the derived figures from Table 3.4 are the results of software aided modeling of detailed analyses of each block of the system. The blocks consist typically not only of the main devices but also controls, instrumentation, valves and piping respectively. Hence, reliability factors (RF), scheduled outage factors and equivalent availabilities are computed for each block as defined by the equations below: forced outage time Reliability factorðRFÞ ¼ 1 6100 % ½3:1 unit period time Scheduled outage factor ¼
scheduled outage time 6100 % unit period time
scheduled outage factor Equivalent availability ¼ RF6 1 100 % Availability factor ¼
½3:2
½3:3
forced outage time þ scheduled outage time 1 unit period time 6100 %
½3:4
Detailed maintenance plans allow shortening the total planned downtime to a minimum, which was assessed for cases A to C (EPRI, 2007b; Higman, 2008). As indicated above, coal gasification holds a potential for unplanned outages. For cases A and B, the averaged values for the reliability factor show that a double-train Shell gasification unit achieves 98.25 %. The same configuration for the GE gasifier would result in a lower reliability of 97.70 %, resulting from burner failures due to erosion and problems with the slag let-down system. Moreover, a design is chosen in which heavily fouling convective syngas coolers are replaced by a water quench. In order to enhance reliability a final configuration of three trains with 50 % capacity was selected for case B, exhibiting reliabilities up to 99.58 %. The gas treatment in cases A and B differs in configuration but shows high reliability values of 97.98 % in case A and 97.85 % in case B dominated mainly by the Claus process. The combined cycle is assumed to have an identical reliability of 94.91 % in both cases. Due to the higher oxygen demand of the slurry fed GE gasifier a setup with two air separation units is reasonable, improving reliability from 98.52 to 99.16 %. The total reliability factor accounts for 89.36 % for case A and 91.02 % for case B, which is equivalent to an annual downtime of 39 and 32 days respectively, if there were no scheduled outages. Defining the scheduled downtimes, the studies used optimized maintenance plans, yielding a minimum annual scheduled
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downtime of 19 days in case A and up to 20 days in case B due to burner replacements because of erosion problems. Hence, an equivalent availability can be calculated, being 84.80 % in case A and 86.12 % in case B. Regarding case C, the whole unit has a stringent single train configuration except for the coal mills. Consequently, the gas production unit has a low reliability of 96.73 %, although employing Shell gasification. Due to CO shift conversion and additional CO2 separation and CO2 compression, the gas treatment also shows a slightly decreased reliability of 97.08 %. The air separation unit and the combined cycle deploying an F-class turbine show the best achievable reported values of 99.50 and 97.87 % respectively. Since the scheduled outage factor is the same as in case A, the equivalent availability accounts for 86.77%, which is fairly high in comparison to the experience-based assessment, published in von Morstein et al. (2009) (see Table 3.4). The study carried out by Sutor et al. (2009) investigates the influence of a non-redundant and all-redundant polygeneration IGCC. It is shown that a fully redundant configuration allows an increase of availability from 85.49 % to 87.89 % by only 2.4 % pts. The simulation reveals that an additional 210 hours of operation per year lead to a steep increase of capital costs and are not favorable. This clearly shows the necessity of understanding the performance of the individual components of the system and using such analysis to ensure that redundancy is only applied where it really brings ‘value for money’.
3.9
Future trends
The 1990s generation of IGCCs in the 250–300 MWel range have paved the way for improved designs, which are inherently more reliable than these demonstration units. The lessons learned number in the thousands, covering design as well as operation and maintenance practices; some of these are sufficiently minor as to appear trivial, while others are a matter of applying best practices from other known technologies and some are at a more fundamental level. Examples of the latter include the move away from 100 % air integration by all systems builders, GE’s decision to replace the convection syngas cooler that has given so much trouble at Polk by a proven water quench design as in the Edwardsport 650 MWel unit (Rigdon and Schmoe, 2005) and Shell’s decision to replace the refractory in the slag quench area by a membrane wall in all its post-Buggenum gasifiers (Chhoa, 2005). Complex and highly integrated systems like IGCC power plants need a comparatively long time to be adjusted and to reach high availability. After start-up, refinery residues fed IGCCs have shown that they can achieve maximum availability up to 90 % in the second or third year of operation.
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In contrast, the demonstration coal fed systems did not reach stable availability until the fourth or fifth year of operation (compare with Fig. 3.4), achieving up to 30 % in the first, up to 50 % in the second and up to 60 % in the third year depending on the duration of the compulsory main revision of the gas turbine. Given the store of lessons learned – in particular the use of best practices in areas such as the ASU, acid gas removal and combined cycle – one can expect a faster ramp-up of availability in the next generation of plants. However, one should not forget that the third year is likely to need a hot gas path inspection of the gas turbine, so that even with much improved performance in years two and three, peak availability is unlikely until the fourth year of operation. Today’s coal based IGCC plants show maximum availabilities of 82.4 % (Hornick, 2007). In contrast, some chemical plants, such as those in Kingsport and Coffeyville, achieve availabilities in the 95–98 % range on a regular basis (EPRI, 2007a). While one cannot expect such figures from an IGCC simply on the basis of necessary gas turbine inspection procedures, it does show that the ASU and acid gas removal (which in both cases include CO shift and CO2 capture) are able to support such availability. In both cases the gasifiers used are refractory lined, so that a spare gasifier is provided to allow on-stream refractory replacement. This in turn does provide a window for some preventative maintenance for a limited portion of the gasification area. Different studies (EPRI, 2007a; von Morstein et al., 2009) have been published indicating the long-term best achievable values for each unit based on manufacturer experiences (see Table 3.5). In order to describe trends towards higher stable availability of a mature IGCC plant, it is instructive to identify the longest foreseeable planned downtime, which is typically the inspection of the hot gas duct of the gas turbine. All other maintenance activities take less than these tasks and should ideally be accomplished in parallel. According to the gas turbine downtime, the advent of the F-class model in IGCC technology results not only in increased efficiency but also lower availability. (This lower availability is due at least in part to longer planned maintenance periods, which are required because of the larger physical size of the machine, number of burners, etc.) Reasonable values can be derived from natural gas fired F-class turbines exhibiting 20 to 25 days planned maintenance (5.5–6.8 %) (von Morstein et al., 2009; ORAP) and 8–11 days unplanned production loss (2.2–2.9 %) (Higman, 2005; von Morstein et al., 2009). Since there is only limited experience in adapting advanced natural gas turbines to syngas, more unplanned downtimes can be expected in the first few years. In the case of an E-class turbine, the planned downtime could be shortened to 15 days (4.3 %), being just sufficient to replace the gasifier’s refractory wall. Since all other units have shorter maintenance periods than
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1.2 E-class 7.2
2.2 F-class 15.9
7.9 76.2
Planned maintenance (% p.a.) Availability
6.8 85.0
2.2 F-class 8.2
0.5 n/a Claus No
0.5 5.0 Shell/GE
EPRI prognosis
Sources: McDaniel (2008); EPRI (2007a); von Morstein et al. (2009).
4.3 88.5
0.5 n/a Claus No
1.2 MDEA Sulfuric acid No
0.5 5.0 Shell/GE
EPRI prognosis
Gas treatment AGR Sulfur recovery CO2 separation CO shift Combined cycle GT type Total unplanned downtime
3.4 9.1 GE
Polk IGCC average 2006–2008
5.5 83.7
2.9 F-class 10.8
1.7 MDEA OxyClaus No
1.3 4.9 Shell
von Morstein et al. case 1 prognosis
5.5 84.1
2.9 F-class 10.4
1.2 Flexsorb Clinsulf No
1.3 5.0 HTW
von Morstein et al. case 2 prognosis
Comparison of availability experiences of Polk Power Station to several studies
Unplanned downtime (% p.a.) ASU Gas production Gasifier
Table 3.5
5.5 83.5
1.3 4.9 Dry feed entrained flow + water quench 1.9 Rectisol OxyClaus Yes Sour 2.9 F-class 11.0
von Morstein et al. case 3 prognosis
5.5 83.4
1.9 Rectisol OxyClaus Yes Sour 2.9 F-class 11.0
1.3 5.0 HTW
von Morstein et al. case 4 prognosis
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the gas turbine, ideally only unplanned outages account for additional losses in availability. The unplanned outage of typical ASU in the industrial gas industry is about 2 days (0.5 %) total in a year, which can be mitigated in part by liquid oxygen storage. In a worst case an unplanned de-riming cycle of about 5 days (1.3 %) may be needed. The unplanned outage to be expected from the gasification unit is more difficult to quantify. Despite applying many of the lessons learned, including many of those mentioned earlier, some malfunctions and problems must still be expected, particularly in connection with equipment involved in solids handling. Undetected fluctuations in feedstock quality are another potential source of downtime. A conservative assessment would foresee a potential for about 18 days (5.0 %) unplanned outage. This value would take credit for some small but not dramatic improvement over the performance of the demonstration units. In principle, gas treatment systems are mature technology and have high availabilities. Hence, they do not influence significantly the global availability of an IGCC plant. For the sulfur recovery a double-train strategy (2660 %) is favorable, allowing short-term partial load operation while avoiding cooling down and restarting the gasifier. Despite deployment of completely different washing processes in different studies, Table 3.5 shows expected outages of 2–6 days (0.5–1.7 %), which extends to 7 days (1.9 %) in the case of CO2 separation. Here, a comparison with the performance of the chemical plants mentioned above, which also employ different solvents (Rectisol in Kingsport and Selexol in Coffeyville) for desulfurization and include CO shift and CO2 capture, demonstrates that such values are on the conservative side. Using the above values, an overall availability of 83.7–85.0 % can be expected for a generic IGCC plant deploying an F-class turbine system. Integration of CO2 capture may reduce availabilities to about 83.5 % due to the additional CO shift reactor and CO2 compressor. The most recent demo plant at Nakoso, Japan, already demonstrates the improvements and lessons learned since the mid-1990s. Within the first year a 2000 hour continuous reliability run was achieved. During the following year a 5000 hour durability test has been performed with five minor incidents. The conclusion by the operator, Clean Coal Power R&D Co. Ltd., is that ‘Every incident could be solved, and know-how has been accumulated so that commercial plants will be realized with high reliability.’ (Watanabe, 2010).
3.10
References
Anon. (2009), Fluor to rebuild fire-damaged ISAB Energy IGCC, Online news in Modern Power Systems, 9 October 2009, Global Trade Media.
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Anon. (2010), IGCC state-of-the-art report, EU-FP7, Sub Project 4: H2-IGCC. WP 1: System Analysis, Department of Mechanical and Structural Engineering and Material Science, University of Stavanger. Arienti, S., Boeri, C., Sinisi, M. and Ciccarelli, P. (2005), Availability improvement of an integrated gasification combined cycle plant: a successful example at api Energia SPA – Falconara facility, PowerGen Europe, Milan. Braekler, K. U. and Kamka, R. (2004), IGCC operation experience of 50.000 hours, PowerGen Europe, Barcelona. Buryan, P., Mika, P., Bucko, Z. and Higman, C. (2008), Reporting from Vresova: 12 years of operating experience with the worlds largest coal-fuelled IGCC. Modern Power Systems, October 2008. Chhoa, T. (2005) Shell gasification business in action, in Gasification Technologies Conference, San Francisco, California. Ciccotosto, V. F., Mezzanotte, P., Pitari, G., Galletta, F., Rottino, V. and Brkic, D. (2009), 17 years of experience gained from three gasification plants operating in Italy, in 3rd International Freiberg Conference on IGCC and XtL Technologies, Dresden. Collodi, G. and Brkic, D. (2003), The experience of Snamprogettis four gasification projects for over 3000 MWth, in Gasification Technologies Conference, San Francisco, California. de Graaf, J. D., Koopmann, E. W. and Zuideveld, P. L. (1999) Shell Pernis Netherlands Refinery Residues Gasification Project, in Gasification Technologies Conference, San Francisco, California. Domenichini, R. (2009), Precombustion capture plants – IGCC with CCS, design and experience, in Workshop on Operating Flexibility of Power Plants with CCS, London. Dowd, R. A. (2000), Wabash River Coal Gasification Project: Final Technical Report, Wabash River Energy Ltd. ELCOGAS (2001), IGCC Puertollano – A Clean Coal Gasification Power Plant, based on Final Report of Thermie program. ELCOGAS (2010), First CO2 captured in the pilot plant of ELCOGAS Puertollano IGCC power plant, Press Release, 14 September 2010. EPRI (2007a), Integrated Gasification Combined Cycle (IGCC) Design Considerations for High Availability – Volume 1: Lessons from Existing Operations, Palo Alto, California, 1012226. EPRI (2007b), Integrated Gasification Combined Cycle (IGCC) Designs for High Availability – Volume 2: RAM Modeling of Standard Designs, Palo Alto, California, 1014871. Eurlings, J. and Ploeg, J. (1999), Process performance of the SCGP at Buggenum IGCC, in Gasification Technologies Conference, San Francisco, California. Gra¨bner, M., von Morstein, O., Rappold, D., Gu¨nster, W., Beysel, G. and Meyer, B. (2010), Constructability study on a German reference IGCC power plant with and without CO2-capture for hard coal and lignite, Energy Conversion and Management, 51, 948–958. Hannemann, F. (2002), V94.2 Buggenum experience and improved concepts for syngas application, in GTC 2002: Economics, Performance and Reliability, San Francisco, California.
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Hannemann, F. (2010), Gasification: status, outlook and technical developments, in 4th International Freiberg Conference on IGCC and XtL Technologies, Dresden. Hatayama, M. (2006), Clean fuel production technologies from high sulfur content crude oil of Middle East, in CPS–JPEC Joint Seminar on Refining Technology, Japan. Higman, C. (2005), Reliability of IGCC power plants, in Gasification Technologies Conference, San Francisco, California. Higman, C. (2008), COORIVA study, report # AP1008: RAM-Analyses of Integrated Gasification Combined Cycle (IGCC)-power plants, in Constructability study of a reference IGCC power plant for 2015 for lignite and hard coal including CO2 capture (in German), IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. Holt, N. (2006), Gasification Technology Status – December 2006, EPRI, Palo Alto, California. Hornick, M. (2007), Polk Power Station Unit 1, Mulberry, Florida, Power, 151, 48– 53. Ishibashi, Y. (2009), Second year operation results of CCP´s Nakoso 250 MW airblown IGCC demonstration plant, in Gasification Technologies Conference, San Francisco, California. Ishibashi, Y. and Shinada, O. (2008), First year operation results of CCPs Nakoso 250 MW air-blown IGCC demonstration plant, in Gasification Technologies Conference, Washington, DC. Kanaar, M. (2002) Operation and performance update Nuon Power Buggenum. Gasification Technologies Conference. San Francisco, CA. Kanaar, M. and von Dongen, A. (2006), Co-gasification at Buggenum power plant, in 7th European Gasification Conference, Barcelona. McDaniel, J. E. (2000), Polk Power Station IGCC 4th year of commercial operation, in Gasification Technology Conference, San Francisco, California. McDaniel, J. E. (2001), Polk Power Station IGCC 5th year of commercial operation, in Gasification Technology Conference. San Francisco, California. McDaniel, J. E. (2002), Tampa Electric Polk Power Station Integrated Gasification Combined Cycle Project: Final Technical Report, US, DOE, NETL. McDaniel, J. E. (2007), Tampa Electric Company: Polk Power Station 250 MW IGCC: Part 3 – Availability, Compact Course Gasification, IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. McDaniel, J. E. (2008), Tampa Electric Company: Polk Power Station 250 MW IGCC: Part 3 – Availability, Compact Course Gasification, IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. Me´ndez-Vigo, I. (2002), Elcogas Puertollano IGCC update, in Gasification Technologies Conference, San Francisco, California. NETL (2002), Pin˜on Pine IGCC Power Project: A DOE assessment. DOE/NETL2003/1183. Morgantown, West Virginia. Nykomb-Synergetics (2007), IGCC projects and plants summary. Available from http://www.nykomb-consulting.se/index.php?s=Projects [Accessed 06 June 2010]. Ono, T. (2003) NPRC Negishi IGCC startup and operation, in Gasification Technologies Conference, San Francisco, California.
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ORAP Databank, Product of Strategic Power Systems, Inc (SPS). Available at http://www.spsinc.com. Payonk, R. (2000), An overview of the past years activities for the Wabash River Repowering Project, in Gasification Technologies Conference, San Francisco, California. Payonk, R. (2008) Wabash River operations update, in Gasification Technologies Conference, Washington, DC. Pen˜a, F. G. (2005), Operating experience and current status of Puertollano IGCC power plant, in International Freiberg Conference on IGCC and XtL Technologies, Freiberg. Pen˜a, F. G. (2007), Puertollano IGCC power plant – operating experience and current developments, in 2nd International Freiberg Conference on IGCC and XtL, Freiberg. Pen˜a, F. G. (2008), Prenflo gasification/Puertollano, Compact Course Gasification, IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. Pen˜a, F. G. (2009), Prenflo gasification/Puertollano, Compact Course Gasification, IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. Pen˜a, F. G. (2010), R & D plan results and experiences in the Puertollano IGCC, in 4th International Freiberg Conference on IGCC & XtL, Dresden. Radtke, K. (2007), Hydrogen from coal, in 7th Annual Meeting of the Fuel Cell and Hydrogen Network NRW, Du¨sseldorf. Rigdon, R. and Schmoe, L. (2005), GE and Bechtel IGCC alliance, in Gasification Technologies Conference, San Francisco, California. Scheibner, B. and Wolters, C. (2002), Schumacher hot gas filter long-term operating experience in the NUON power Buggenum IGCC power plant, in Proceedings of the 5th International Symposium on Gas Cleaning at High Temperature, Morgantown, West Virginia. Schingnitz, M. (2008) Gaskombinat Schwarze Pumpe-Verfahren (GSP), in Schmalfeld, J. (ed.), Die Veredlung und Umwandlung von Kohle, Hamburg, DGMK, pp. 537–552. Schmalfeld, J. (ed.) (2008), Die Veredlung und Umwandlung von Kohle, Hamburg, DGMK. Sharp, C., Kubek, D., Kuper, D., Clark, M. and Didio, M. (2002), Recent SelexolTM operating experience with gasification including CO2 capture, in 5th European Gasification Conference, Noordwijk. Sutor, A., Confer, M. and Streer, W. (2009), RAM development for gasification and IGCC plants, in Gasification Technologies Conference, Colorado Springs, Colorado. Verwilligen, J. A. (2010), How to develop a maintenance programme for the new Magnum power plant – the Nuon Energy approach, VGB PowerTech, 26–32. von Morstein, O., Kuske, E., Rappold, D., Gu¨nster, W., Graeber, C., Beysel, G., Alekseev, A., Buschsieweke, F., Korbov, D., Riedl, K., Rainer, H. and Gra¨bner, M. (2009), COORIVA study, report # AP2001: improved IGCC power plant with and without CO2 capture (final report), Constructability Study of a Reference IGCC Power Plant for 2015 for Lignite and Hard Coal Including
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CO2 Capture (in German), IEC – Department of Energy Process Engineering and Chemical Engineering, Freiberg. Watanabe, T. (2010), Japanese Air Blown IGCC Project Progress, Clean Coal Power R&D Co. Ltd. Wolters, C. (2003), Operating experience at the ‘Willem-Alexander Centrale’, in Gasification Technologies Conference, San Francisco, California. Yamaguchi, M. (2004), First year of operational experience with the Negishi IGCC, in Gasification Technologies Conference, Washington, DC. Zuideveld, P. L. (2001), Overview of Shell gasification projects, in Gasification Technologies Conference, San Francisco, California. Zuideveld, P. L. (2002), Overview of Shell gasification projects, in 5th European Gasification Conference, Noordwijk.
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4 Environmental degradation of boiler components N . J . S I M M S , Cranfield University, UK
Abstract: During the operation of a boiler it is anticipated that there will be a range of interactions between the boiler components and their operating environments. These interactions are often classified in terms of deposition, erosion or corrosion, and can all degrade the operation of the boiler, either in terms of restricting heat transfer or reducing the potential lives of the boiler components. Understanding and modelling the mechanisms of such interactions enables their extent and effects to be reduced, e.g. by changing the operating environments or the materials used in the boilers. As a result, the efficiency of boiler operations can be improved and/or the lives of components in the boilers can be extended. This reduces the operating and maintenance costs associated with boilers and reduces the risks of costly unplanned outages. Key words: deposition, fireside corrosion, superheater corrosion, waterwall corrosion, coal, biomass, co-firing.
4.1
Introduction
The operating environments for components within power plants are a result of the fuels used, the power plant design and the component operating conditions. Most current large scale power plants contain a steam generating system that drives a steam turbine, either on its own (a Rankine cycle) or in combination with a gas turbine (as part of a combined cycle). For heat recovery steam generators located downstream from gas turbines in combined cycles, the environments are usually relatively benign as their gas stream inlet temperatures are generally limited by the gas turbine exhaust temperatures and the contaminant levels in these gases are low (to meet the strict limits placed on gas turbine operating conditions). However, power systems that rely entirely on heat exchange from a hot gas stream to a water/steam system can have much more aggressive environments around 145 © Woodhead Publishing Limited, 2011
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the fireside surfaces in their hot gas paths. In such systems, the hot gas streams can be produced by the combustion of a wide variety of fuels (solid, liquid or gaseous), which can contain a range of different impurities. As these hot combusted gas streams pass through the boilers and the various heat exchange surfaces, as well as transferring heat to the water/steam system they can interact to produce deposition, erosion and/or corrosion on the heat exchanger surfaces. Both corrosion and erosion damage to the fireside surfaces of the heat exchangers cause metal losses and so reduce component lives (though there is often a ‘corrosion allowance’ to enable design lives to be achieved). In order to avoid unexpected tube failures, which can result in costly plant outages, significant effort is devoted to non-destructive examinations of heat exchanger tubes during routine plant outages so that any affected tubing can be identified and replaced. The replacement tubing can be the same (if the component life is acceptable), or a better material may be used (if one exists), or protective measures may be required (such as coatings, coextruded tubes, bandages, etc.). Alternatively, the boiler operating conditions could be changed to reduce the damage rates or the compositions of the fuels used in the boiler restricted. Deposit formation usually has the effect of reducing heat transfer from the hot gas stream to the water/steam system, which in turn reduces boiler efficiency. In addition, such deposits are involved in some of the corrosion damage mechanisms that have been found in boiler environments (section 4.3). As a result, a range of techniques have been developed to try to remove deposits from the surfaces of heat exchangers during operations or periods of maintenance, including: mechanically hitting the tube surfaces; sootblowing using compressed air, steam or water jets; sonic waves; explosive charges; etc. This chapter describes the operating environments that are encountered in selected heat exchangers in solid-fuel-fired boilers. The complex mechanisms that govern the environmentally induced degradation of these components are outlined, together with available models of these processes. These environments and degradation processes can change during the life of a power plant as a result of alterations in the fuels used and/or operating conditions. Monitoring of the degradation of boiler tubes is described, together with preventative measures. Finally, the anticipated challenges in environmental degradation of boiler tubes is considered in terms of likely future developments for solid-fuel-fired power plants.
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4.2
147
Component operating environments
4.2.1 Superheaters/reheaters/waterwalls, etc. The hot gas paths of combustion systems contain a series of heat exchangers to generate high temperature/high pressure steam from water. Figure 4.1 shows a flow diagram for a typical water/steam system, with the series of heat exchangers being an economiser, evaporator and superheater before the steam enters the high pressure steam turbine. In this system, the steam is reheated before entering the intermediate pressure steam turbine. The highest steam temperatures in such a system are achieved in the final stages of the superheater and reheater. Figure 4.2 illustrates the layout of these heat exchangers around a conventional pulverised-fuel fired combustion power plant; this shows that the combustion zone is surrounded by waterwalls and that the hot gases from the combustion process then flow past the various superheater and reheater stages before going through the economiser. In such a system, the waterwalls are relatively cool (up to 400 8C) despite containing the fuel burners and gases of up to 1600 8C, but have high heat fluxes (up to ~0.4 MW/m2). The combustion gases have cooled to 1000–1200 8C by the time they pass through the superheaters and produce heat fluxes of ~0.2 MW/m2; the steam temperatures exiting the superheaters can be ~540– 620 8C depending on the age of the power plant. The combustion gases continue cooling through the superheaters, reheaters and economisers. The
4.1 Schematic flow diagram for a power plant steam/water system showing the main component parts.
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4.2 Schematic diagram of a pulverised fuel power plant showing the position of the main heat exchangers.
final stage of the reheaters is at similar steam and metal temperatures (but lower pressures) compared to the superheaters. Figure 4.3 shows a different power plant configuration based on a circulating fluidised bed combustion process, with this example being used to generate heat and power from biomass combustion (Henderson et al., 2002b). In this system, the combustion chamber is again surrounded by waterwalls, but superheaters are located in the gas pass after the cyclone. The flue gas temperature approaching the superheaters is ~860–880 8C. The steam system operates with the final superheater output at 480 8C/80 bars. Figure 4.4 shows a power plant configuration based on a grate-fired boiler, with this example based on a waste to energy process (Henderson et al., 2002a). In this system the combustion gases initially pass through a chamber surrounded by waterwalls, but the superheaters are located in the third gas pass to overcome environmental degradation issues associated with such fuels (section 4.3.3). In this system, these issues limit superheated steam temperatures to 360 8C (with pressures of ~ 33 bars) and such low steam conditions restrict the efficiencies of this type of power plant.
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4.3 Schematic diagram of a circulating fluidised bed (CFBC) biomassfired unit (adapted from Henderson et al., 2002b).
4.4 Schematic diagram of a waste-fired grate unit (adapted from Henderson et al., 2002a).
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Table 4.1
Example solid fuel compositions
Parameter Unit
UK coal
South South American African Wheat Palm Olive coal coal Willow Miscanthus straw nuts residue
Moisture
11.9
13.0
9.1
15.3
9.0
15.2
2.0
4.0
5.0
7.5
4.5
34.3 32.8 81.6 5.2 8.7 1.8 2.28 0.7
32.7 31.1 79.9 5.3 12.21 1.7 0.8 0.1
33.2 31.9 83.2 4.3 9.6 2.0 0.7 0.1
20.3 18.8 49 6.2 44 0.5 0.05 0.03
19.8 18.4 49 6.4 44 0.7 0.2 0.2
19.8 18.5 49 6.3 43 0.5 0.1 0.4
19.8 18.3 50.2 6.6 40 3.2 0.2 0.2
21.4 18.3 49 6.0 40 2.24 0.1 0.1
wt% ar Ash wt% dry CV (gross) MJ/kg CV (net) daf C wt% H daf O N S Cl
15
25
25
4.6 13.5
Source: Simms et al. (2007a).
4.2.2 Fuel effects A range of solid fuels can be combusted in boilers to provide the heat required, for example: . . .
coals, typically bituminous, sub-bituminous or lignite coals depending on the availability and cost of local and world-traded coals; biomass, traditionally residual agricultural by-products, but increasingly energy crops; waste products, derived from municipal solid wastes, sludges, etc.
These three general classes of solid fuels are described in detail in Chapter 1 of this book, which also describes the appropriate analytical methods for each of the fuels together with the characteristic properties of each type of fuel. Within boilers, the environments around the heat exchangers depend on the chemical compositions of the fuels used, as well as the operating conditions used in the boilers. Table 4.1 provides examples of each of these types of fuels in terms of their chemical composition. During the combustion processes, the fuels react with an oxidising gas stream, which is air in most current combustion plants, to produce a hot combusted gas stream. Figure 4.5 (Tomeczek and Palugniok, 2002) illustrates this process for pulverised coal combustion and shows the breakdown of the fuel in terms of the burnout of the combustible material and generation of ash particles and vapour phase species. As a result of the complex reactions of the inorganic elements present in the fuels in various different forms (Chapter 1), the minor and trace elements are partitioned between the coarse (bottom) ash, fly ash and gas/vapour phase. It is the fate
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Schematic representation of fuel combustion (adapted from Tomeczek and Palugniok, 2002).
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of these elements as they pass through the hot gas path of the power systems that is in a large part responsible for the environmental degradation of the heat exchangers. The particles produced can form deposits (section 4.3.1) or cause erosion damage (section 4.3.4); the vapour phase species can condense under particular conditions to become part of the deposits (section 4.3.1); the gases, deposits and heat exchanger surfaces can react to cause accelerated corrosion damage (section 4.3.3).
4.3
Degradation mechanisms and modelling
4.3.1 Deposition The deposits that form around the fireside surfaces of heat exchanger tubes are created from the particles and vapours that pass through a boiler by the action of a number of different mechanisms that can occur in parallel in the local environments, as shown in Fig. 4.6 (Simms et al., 2007b). For particles, the important potential deposition mechanisms are:
4.6 Schematic representation of interaction between superheater tube and its local environment (Simms et al., 2007b).
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direct inertial impaction: for larger particles, typically >10 μm on the upstream tube surfaces; thermophoresis: for smaller particles, typically <1 μm; Brownian/eddy diffusion: for particles <1 μm.
For vapours, the mechanisms are: . . .
vapour condensation on to the heat exchanger tubes; heterogeneous vapour condensation on to particles (which can then follow a particle deposition route dependent on the particle size); homogeneous vapour condensation into aerosols (which then follow particle deposition routes appropriate to smaller particles).
Thus, the total deposit growth rate can be thought of as a sum of the various possible deposition mechanisms (Zhou et al., 2007): dmðt; yÞ=dt ¼ Cðt; yÞ þ THðt; yÞ þ BEðt; yÞ þ Iðt; yÞ
½4:1
where m(t,θ) is the deposit weight, t is time and C(t,θ), TH(t,θ), BE(t,θ) and I(t,θ) represent the condensation, thermophoresis, Brownian and eddy diffusion and impaction rates, respectively, as a function of t and θ, the angle around the heat exchanger tube. Vapour condensation This is the mechanism that governs the transfer of vapour species from a hotter gas stream on to a cooler local surface. The flux of condensable species for the ith component in a gas stream may be calculated by the equation (Tomeczek et al., 2004): ? m d;i ¼ ðSh Di =dÞ½ðpi psi Þ=prg
½4:2
where Sh is the Sherwood number, Di is the diffusion coefficient of the ith component through the flue gases (m2/s), d is the tube diameter (m), pi is the pressure of the ith component in the flue gas stream (Pa), psi is the saturation pressure of the ith component at the surface temperature (Pa), p is the gas pressure (Pa) and ρg is the gas density (kg/m3). The saturation pressure can be calculated using the following equation (Tomeczek and Wacławiak, 2009): psi ¼ pn exp½Ai Bi =ðT þ Ci Þ
½4:3
where Ai, Bi and Ci are constants for the ith species, T is the surface temperature (K) and pn is 105 Pa. The constants Ai to Ci are given in Table 4.2 for key alkali sulphate and chloride compounds in coal and biomass combustion systems: NaCl, KCl, Na2SO4 and K2SO4. Figure 4.7
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Table 4.2 Constants required to calculate saturated vapour pressures of alkali sulphates and chlorides Component
Ai
Bi (K)
Ci (K)
Temperature range (K)
K2SO4 Na2SO4 KCl NaCl
18.08 15.03 11.01 11.68
39449 37452 17132 19315
0 0 –122.7 –82.6
1150–1800 1150–1800 1094–1680 1138–1738
Source: Tomezcek and Wacławiak (2009).
4.7 Effect of sulphur, chlorine and alkali metal availability on deposit compositions (adapted from Robinson et al., 1998).
illustrates the variation in the deposit chloride content as a function of the fuel sulphur, chlorine and available alkali levels (Robinson et al., 1998). Direct inertial impaction This is the mechanism by which larger particles (usually >10 μm) deposit on to the surfaces of heat exchangers. This deposition mechanism is particularly important for the upstream surfaces of heat exchanger tubes, with the larger particles not being able to follow the gas flow streamlines around the tubes (as a result of their having too high a momentum). Particles hitting the tube surfaces may either rebound from the surface or stick to it, depending on their state (solid or sticky) and that of the tube surface (e.g. with liquid or sticky deposits). One approach to calculating the deposition of particles by this mechanism
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is given by Zhou et al. (2007): Iðt; yÞ ¼ uf Cash;1 Z1 fstick s
½4:4
where uf is the bulk gas velocity (m/s), Cash,l is the mass concentration for large fly ash particles (kg/m3), ηl is the impaction efficiency (according to the impaction models developed by Rosner and Tandon, 1995; Wessel and Righi, 1988), fstick is the total sticking coefficient and s is the surface area (m2) at the angle section Δθ. An alternative approach to calculating this type of deposition process is given by Tomeczek et al. (2004): ? m ¼ P1 P2 Cp;s w
½4:5
? where m is the mass rate of deposition per unit surface area (kg/m2 s); P1 is the probability of the particles hitting the heat transfer surface (P1 = (Aduct – Afree)/Aduct), with Aduct and Afree being the cross-sectional areas of the duct and the free gas path respectively; P2 is the probability of the solid particles ? ? sticking to the surface (P2 = Smd), with md being flux of vapour condensation (kg/m2 s) and S = 6.2 (kg/m2 s)1; Cp,s is the concentration of the solid particles in the gas stream (kg/m3); and w is the gas velocity component perpendicular to the local point of the heat transfer surface (m/s). Direct inertial impaction can also occur on the downstream surfaces of tubes when particles that have passed by the tubes become caught up in turbulence and are then propelled towards a tube surface. Zhou et al. (2007) give an expression to calculate the particle sizes that may respond to the turbulent vortices: 0:5 ? dp m 0:5 rf m0f mf rp rf þ 1 ½4:6 ? where dp is the particle diameter (m); m is the macroscale of the turbulence (which may be taken as the tube diameter) (m); ρf is the gas density (kg/m3); u´f is the root mean square velocity of the gas in the wake of the probe (m/s); μf is the gas viscosity (Pa s); and ρp is the particle density (kg/m3). Other deposition processes Other deposition processes that were outlined above are believed to play a more minor role in deposit formation in boilers: .
Thermophoresis is a process that results in the transport of small particles through a gas along local temperature gradients (e.g. submicrometre particles from a hot gas stream to a cooled heat
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Power plant life management and performance improvement exchanger tube). Jacobsen and Brock (1965) report a model for this process. Brownian and eddy diffusion of submicrometre particles that produce deposition from turbulent gas streams. Wood (1981) provides a model for such deposition routes.
Deposit compositions The deposit compositions found on heat exchanger surfaces are determined by the balance of the deposition processes occurring around that component, and so can be boiler design specific in addition to depending on the composition of the fuels and the element partitioning that happens during the combustion processes. There are many reports of different deposit compositions available in the literature; these just emphasise the sensitivity of the various deposition processes to the fuels, combustion processes and local component operating environments. For superheater/reheater tubes in coal-fired systems, deposits can contain: . . . .
Si–Al–O compounds ○ derived from alumina–silicate minerals, ○ can fix Na, K if particle temperatures are high enough; Ca/Mg carbonates/sulphates/chlorides; Na/K sulphates/chlorides; Fe sulphates/chlorides/oxides/sulphides.
An example of the type of deposit that is often found on superheater tubes in coal-fired power plants is given in Fig. 4.8 (Nelson and Cain, 1960; Syrett, 1987). This emphasises the layered structure that develops in the deposit as it grows, with the shape and surface temperature of the deposit changing as it grows. For biomass-fired systems, a similar range of compounds can be found in deposits, but the balance between elements varies as a result of the different fuel compositions and element partitioning, as well as changes in the balance between the deposition processes due to differences in the particle size distributions produced and the availability of vapour phase species (especially derived from K). One classification of deposits found in biomass systems produces three main groupings (Livingston, 2010): . . .
high silica/high K/low Ca ashes with low melting points ○ from agricultural residues (straws, etc.); low silica/low K/high Ca ashes with high melting points ○ from woody materials; high K/high P with low melting points ○ from animal wastes, etc.
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4.8 Example of the structure of a deposit observed on superheater tube in a coal-fired boiler (adapted from Nelson and Cain, 1960; Syrett, 1987).
Co-firing coal and biomass produces deposits that are a blend of those expected from the fuels alone. However, there is a complex interaction due to the different size distributions of particles produced by the mixed fuels and vapour phase species. These change the balance of the deposition mechanisms and the resulting deposition fluxes and compositions depend on which specific fuels have been used. For waste-fired systems, there is a much wider range of fuel compositions and so a correspondingly wider range of potential deposit compositions. For many waste streams there is a particular concern over the levels of heavy metals in the wastes as high levels of some of these (e.g. Pb, Zn, Cd, Sn) can produce vapour phase species that condense as low melting point chlorides on the surfaces of superheater tubes (and then cause rapid corrosion damage; see section 4.3.4).
4.3.2 Oxidation During the course of their operation, the minimum chemical degradation that will be experienced by heat exchanger tubes in combustion systems is
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Table 4.3
Nominal compositions of heat exchanger tube materials
Alloy
Fe
Cr
Mo
Ni
1Cr T22 T23 T91 T92 X20CrMoV121 AISI347HFG
Bal 1 Bal 2.25 Bal 2.25 Bal 9 Bal 9 Bal 12 Bal 18
0.5 1.1 0.15 0.25 1 0.22 0.5 <0.4 0.22 1 <0.5 0.3 — 11
HR3C Alloy 625
Bal 25 <5 21
9
20 Bal
V
Nb
W
B
Others
0.05 1.6 0.002 0.08 <0.04 Al, 0.05 N 0.07 1.8 0.003 <0.04 Al, 0.05 N *
—
—
* *
*Nb + Ta = 86C or <1 % Nb + Ta = 0.4, 0.4 N *Nb + Ta = 3.7
oxidation. In this process, the tube materials react with oxygen (or oxygencontaining species) to generate a surface oxide scale. In its simplest form this can be represented as: M þ 12O2 < g >¼ MO
½4:7
where M represents a metal and MO represents a metal oxide. The continued oxidation of this metal depends on the transport of metal ions and/or oxidant species through the metal oxide scale. The most protective oxides that can grow are those that form dense, even scales and only permit a slow transport of metal and oxidant through them. The oxidation of metals has been thoroughly studied, so there are many textbooks that describe these processes well (e.g. Kofstad, 1988; Birks et al., 2006; Young, 2008). Heat exchanger materials are usually manufactured from various grades of steels, ranging from low alloy ferritic steels through ferritic–martensitic varieties to austenitic stainless steels; examples of nominal compositions are given in Table 4.3. In some situations (section 4.4) higher alloyed materials are used as coatings and nickel based alloys are also now being considered for use in future power plants with higher temperature steam systems (section 4.5). The oxides that form on low alloy steels (<~10–12 wt% Cr) under oxidising conditions are multilayered. At lower temperatures (below 570 8C for pure iron, but increasing with Cr content), the inner oxide layer is an inward growing spinel (Fe,Cr)3O4, the central layer magnetite (Fe3O4) and the outer layer haematite (Fe2O3). At higher temperatures, wustite (FeO) forms as an inner oxide layer, beneath magnetite and haematite layers. However, the formation of wustite permits much higher oxidation rates, which therefore provides one of the upper limits to the temperatures where such alloys can be used in practical systems (Stringer and Wright, 1995). On higher alloy materials (stainless steels, nickel based alloys and
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4.9 Growth rate constants for a range of different oxides (adapted from Meier, 1996).
coatings), oxide scale growth is dominated by the formation of chromia (Cr2O3), which can be a relatively slow growing protective oxide. In situations where this layer breaks down, various mixed oxide scales can also form (such as (Cr,Fe)2O3 and spinels (Fe,Ni,Cr)3O4) to produce multilayered scale structures or scales with internal oxidation beneath them (Bradford, 1987; Birks et al., 2006; Young, 2008). The relative growth rates of different oxides are illustrated in Fig. 4.9.
4.3.3 Fireside corrosion There are several different types of fireside corrosion that have been found in boilers over the years. Despite many extensive studies, the detailed mechanisms of these degradation processes have proved to be difficult to define fully (e.g. Syrett, 1987; Stringer and Wright, 1995). However, the key factors causing such damage have been identified, with mechanistic and/or empirically based models being developed to describe the effects of some of these corrosion processes under certain conditions.
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4.10 Typical structure of deposit/scale formed on a corroding waterwall surface (adapted from Stringer and Wright, 1995).
Waterwall corrosion In pulverised-fuel-fired boilers, it is sometimes found that there are areas of high metal wastage on the waterwalls. These are usually associated with flame impingement or a failure to establish a combustion zone towards the centre of the furnace. As a result, areas of the waterwall can experience significant exposure periods under reducing conditions, as well as the more usual oxidising regime. The surface scales produced during this type of corrosion are generally based on magnetite, but can contain FeS lamella close to the metal surface, with FeS islands, fly ash spheres and unburnt carbon particles closer to the scale surface; this is shown schematically in Fig. 4.10. As a result of the operating conditions of the waterwall, the metal surface temperatures may only be in the range 300–400 8C, but the deposit surface may be molten. Many causes of this form of damage have been suggested over the years (Syrett, 1987; Stringer and Wright, 1995), but it is now believed that periods in the oxidising and reducing environments coupled with the presence of sulphur (and possibly carbon) are responsible. A predictive model for this type of damage has been suggested by Davis (2010):
0:5 0:5 tr 6ACR M ¼ C to 6Kpo þ tr þ Kpr þ ½4:8 103 where Kpo and to are the rate constant and the time, respectively, under oxidising conditions; Kpr and tr are the rate constant and the time, respectively, under reducing conditions; C is a constant; and h i ACR ¼ ð14256%ClÞ6ðHFÞm 6e½Qu^=ðRTÞ P ½4:9 where %Cl is the % of chlorine in the coal (by weight); HF is the heat flux,
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Q is an activation energy, R is the gas constant, m and P are other constants and T is the metal surface temperature. Solutions to this type of corrosion damage include: . . . .
modifying the combustion environment (using improved burner designs, finer fuel particles or fuel mixes); providing an ‘air curtain’ in front of the waterwall to keep the environment oxidising; replacing tubes with co-extruded tubes (an outer highly corrosion resistant material outside a waterwall tube material) (Meadowcroft and Manning, 1983); coating the fireside surfaces of the tubes (section 4.4).
Another form of waterwall damage has been reported where the corroded tubes show deep parallel groves normal to the tube axis with spacings of the order of 1 mm (Stringer and Wright, 1995). These have been given a variety of descriptive names: e.g. circumferential cracking, horizontal cracking, elephant hiding, alligator-skin cracking. The suggested cause is a combination of (a) thermally induced alternating stresses, (b) the fireside corrosion environment in areas of particularly high heat fluxes and (c) a rippled magnetite layer on the inside (waterside surface) of the tubes. Solutions include oxygenated water treatment to reduce magnetite deposition on the waterside surfaces and methods for ensuring a more even heat flux distribution (e.g. by improved control of slag removal). Superheater/reheater corrosion There has been a long history of investigating the causes of fireside corrosion damage on superheater/reheater tubes in coal-fired power plants (described in detail by Stringer and Wright, 1995). However, it is now believed that the main cause of this form of accelerated damage is the presence of molten deposits on the surfaces of the tubes. These deposits can form as solid species and then become molten as a result of their reactions with other species in the deposit and the gas stream flowing around the tubes. In addition, corrosion products from initial reactions with the surfaces of the tube materials can also take part in further corrosion processes. Thus, in assessing this form of degradation it is necessary to consider both the immediate results of the deposition processes and the many potential further reactions that can take place. Compounds that have been identified as having the potential to form in deposits and cause fireside corrosion damage include: .
Sulphate deposits: ○ Pyro-sulphates, e.g. (Na,K)2S2O7
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4.11
Phase diagram for alkali pyro-sulphates (Lindberg et al., 2006).
4.12 Phase diagram for alkali–iron tri-sulphates (adapted from Cain and Nelson, 1961).
○ ○
Alkali–iron tri-sulphates, e.g. (Na,K)3Fe(SO4)3 Mixed sulphates, e.g. (K,Na,Fe)xSO4
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4.13 Phase diagram for alkali sulphates–chlorides (Lindberg et al., 2007). Note: temperatures in Kelvin.
. . .
Chloride deposits, with mixed compositions including Na, K, Fe, Ca, Mg, and other metal elements depending on the fuel used; Carbonates, with mixed compositions including Na, K, Fe, Ca, Mg, and other metal elements depending on the fuel used; Sulphate–chloride–carbonate ‘soup’ containing all the compounds above.
In considering the potential for such compounds to both form in deposits and cause corrosion damage, it is necessary to assess the melting points of the compounds and the conditions necessary for their formation. For example, both alkali pyro-sulphates and alkali–iron tri-sulphates need sufficient vapour pressures of SO3 around them for their stability to be maintained, with the alkali pyro-sulphates needing the higher levels. Figures 4.11 and 4.12 show phase diagrams of such sulphate species and illustrate that the lowest melting points in these systems are ~345 8C for the mixed alkali pyro-sulphates and ~560 8C for the mixed alkali–iron trisulphates. For biomass-fired systems that have deposits containing higher levels of
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4.14 Phase diagram for zinc chlorides–potassium chlorides (Hack and Jantzen, 2008).
potassium chlorides and sulphates, the alkali chloride–sulphate phase diagram is more useful. Figure 4.13 shows an example with the lowest melting point of a mixed alkali chloride–sulphate being ~517 8C. In addition, for biomass-fired systems the effects of other compounds, such as carbonates, need to be considered (Blomberg, 2008). For waste-fired systems, heavy metal chloride compounds need to be considered in detail, as these can cause deposits to have much lower melting points; potential mixtures contain combinations of alkali metals, Fe, Pb, Zn, Cd and Sn as oxides, chlorides, sulphates and carbonates. Figure 4.14 shows a ZnCl2–KCl phase diagram, which shows a minimum melting point of ~240 8C (Hack and Jantzen, 2008). With increasing metal temperatures such compounds can become unstable for a variety of different reasons, including: . . .
vapour condensation dewpoints being exceeded; insufficient SO3 being available to stabilise some sulphate phases (e.g. alkali pyro-sulphates, alkali–iron tri-sulphates), as the SO3/SO2 balance favours SO3 at lower temperatures; other phases becoming more stable with a change in temperature.
The result of this is a ‘bell-shaped’ curve in materials corrosion (Fig. 4.15). In this the lower limit is set by the melting point of a compound in the deposit and the increase in corrosion rate is dependent on the sensitivity of the corrosion reaction to temperature and the availability of reactants (deposition fluxes, gas partial pressures, etc.) in the local environment. The
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4.15 Characteristic bell-shaped curve for a fireside corrosion damage mechanism.
4.16 Effect of metal temperatures on corrosion rates in conventional pulverised-fuel-fired power systems (adapted from Natesan et al., 2003).
position of the peak (and downwards slope) depends on the cause(s) of the deposit instability (as listed above). Thus, the position of the peak and downwards slope of the bell-shaped curve can be influenced by the exposure environment, and so can be set differently in laboratory and plant exposures. Figure 4.16 shows multiple ‘bell-shaped’ corrosion peaks
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4.17 Relationship between the physical state of complex alkali sulphates and temperature compared to the corrosion behaviour of austenitic alloys (adapted from Stringer and Wright, 1995).
attributed to different compounds forming in deposits (Natesan et al., 2003). Figure 4.17 illustrates the relationship between molten, sticky and solid alkali–iron tri-sulphate compounds and a bell-shaped corrosion peak for austenitic stainless steels. The various alternative deposit compositions and their stabilities can cause different corrosion mechanisms. Figure 4.18 illustrates the different steps involved in sodium sulphate and chloride induced fireside corrosion on superheater tubes. Heat exchanger tube materials will respond in different ways to the aggressive deposits on their surfaces. For some lower alloyed materials, the fluxing reactions in the molten deposits result in rapid corrosion, whereas for more highly alloyed materials, the chromia scale formed is more protective and can provide some protection against such deposits. However, much higher levels of chromium are needed in the alloy to provide significant resistance against fireside corrosion than used in standard stainless steels (hence the high chromium contents of materials selected as potential protective coatings for heat exchanger surfaces). The relative corrosiveness of the superheater environments produced by one biomass (wheat straw) and coals is illustrated in Fig. 4.19. This figure shows the results of corrosion damage measurements carried out on the stainless steels AISI 347 and 347HFG exposed in combustion units as part of the EU COST 522 and 538 programmes, with the materials being exposed
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4.18 Examples of two of the potential corrosion mechanisms on a superheater in a pulverised-coal-fired power system. Reproduced with permission from Maney Publishing (NJ Simms, PJ Kilgallon and JE Oakey, ‘Fireside issues in advanced power generation systems’, Energy Materials: Materials Science and Engineering for Energy Systems, 2(3), 2007, 154–160. www.maney.co.uk/journals/ema www.ingentaconnect. com/content/maney/ema).
as parts of superheaters, reheaters and on cooled probes (e.g. Henderson et al., 2002a). The results show that the wheat straw induces a higher range of corrosion damage rates than the coal or co-fired conditions used (<10 % biomass by mass) at the same metal temperature. The range of corrosion damage rates observed reflects the known variability of wheat straw compositions and variations in local exposure conditions. The apparent temperature dependence of the corrosion damage rates is a result of a combination of the corrosion dependence outlined in Fig. 4.15 and a need to focus research activities on potentially realistic metal temperatures for component operation. This work provided aspirational targets for coal, biomass-and waste-fired power plants; these are summarised in Table 4.4 and illustrate the generally increasing aggressiveness of fireside corrosion that has been found for coal, biomass and waste fuels. As more data have been generated for fireside corrosion of superheater/ reheater tubes, there has been a desire to generate mathematical models to represent such forms of corrosion damage. Different approaches have been developed ranging from mechanistic modelling through empirical curve fitting to neural networks (Saunders et al., 2002). A relatively simple model
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Table 4.4 Aspirational targets for superheater/reheater tube lives, steam temperatures and degradation rates in coal, biomass and waste systems for EU COST 522/538 programmes
Fuel
Target lifetime (hours)
Desired maximum steam temperature (8C)
Metal temperature (8C)
Maximum acceptable corrosion rate (μm/1000 hours)
Coal Straw Wood Waste
100 000 20 000 40 000 40 000
650 580 580 500
680–700 610–630 610–630 530–550
20 100 50 50
Source: adapted from Henderson et al., 2002a.
4.19 Effect of metal temperatures on fireside corrosion of AISI 347 and 347HFG for combustion systems fired on wheat straw and coals.
was generated for UK coals (James and Pinder, 1997): Corrosion rate ¼ LE AðTg ÞB ðTm CÞD ðCl%fuel EÞ
½4:10
where A to E are constants, LE represents the leading edge of a tube bundle, Tg is the average gas temperature and Tm is an average metal temperature and Cl% fuel represents the average fuel Cl content. Alternative approaches have been reported by Larson and Montgomery (2006), Simms et al. (2007b), Simms and Fry (2010), Heikinheimo et al. (2008) and Linjewile et al. (2003). All the approaches have different benefits and limitations, and are still at various stages of development.
4.3.4 Erosion/wear Erosion is a damage mechanism that causes metal loss as a result of particles impacting on a surface (Finnie, 1995). Erosion damage has been found to
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vary with particle size, particle velocity, particle hardness, tube surface hardness and impact angle. Brittle and ductile damage regimes have been identified depending on the impact conditions. For heat exchanger tubing operating at higher temperatures, there can be an interaction with the oxide scales that form, resulting in what has been termed ‘erosion/corrosion’ damage (Stack et al., 1995). In this case the impacting particles can interact with either a surface oxide layer or the underlying alloy, depending on the exact exposure conditions. As a result, a number of different erosion/corrosion regimes have been identified, ranging from pure erosion, through oxidation enhanced erosion and erosion enhanced oxidation to modified oxidation depending on the impact conditions and temperatures. In pulverised fuel boilers, erosion damage can occur to the waterwalls, superheaters/reheaters and the economiser (Stringer, 1995; Foster et al., 2004), with fly ash particles either eroding the tube material directly or the surface oxide (for tube surfaces >425 8C). For example, erosion/corrosion damage can be found in the superheater/reheater platens under conditions in which deposit blockages have built up between some tubes in these platens and so have caused locally higher gas velocities elsewhere in the platens. Erosion/corrosion conditions have been a particular challenge in fluidised bed combustion systems, both for waterwalls and in-bed heat exchanger tubing (Stringer, 1995). As a result of many years of investigating cases of such damage, these damage modes can now usually be avoided by careful engineering design.
4.4
Quantification of damage and protective measures
4.4.1 Component and materials monitoring methods The environmental degradation of heat exchanger tubes has been studied in plants, pilot plants and laboratory tests. Each of these types of environment has its own particular benefits and limitation in terms of materials monitoring. In plant environments, the traditional monitoring of heat exchanger tube materials is carried out using a mixture of visual inspection, dimensional metrology and ultrasonic inspections. The data generated are used in combination with an assessment of the remaining life of the tubes to determine the risk of component failure before the next scheduled maintenance, and so whether any tubes need to be removed from service. Tubes removed from boilers (during plant maintenance or outages) can be destructively examined using standard laboratory techniques, including optical and electron microscopy, energy dispersive X-ray (EDX) analysis, X-
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ray diffraction, etc., to investigate their performance further (or determine the cause of a failure). However, the data generated from plant heat exchanger tubing can be difficult to interpret as a result of changing fuels (e.g. different coals or biomass, fuel preparation, etc.) and operating conditions (e.g. gas temperatures, metal temperatures, air/fuel ratios, etc.), as well as a lack of monitoring data. Specific materials evaluation programmes for extended periods (thousands or tens of thousands of hours) can be carried out in plants by several methods, including: .
. . .
Installing candidate materials within the heat exchangers, e.g. as short lengths within a superheater/reheater or as a small panel in a water wall (Henderson et al., 2002a). These pieces then need retrieving at appropriate plant outages. Installing materials in separate water/steam cooled loops within a boiler (Henderson et al., 2002a; Larson and Montgomery, 2006). These loops then need removing at appropriate plant outages. Exposing materials on cooled probes (using air, water or steam cooling) in the boiler environment (Henderson et al., 2002a). These can usually be retrieved during plant operation. Using on-line monitoring methods. These are at relatively early stages of development, but are being used by research activities to try to assess corrosion rates, e.g. by using electrochemical noise (ECN) and/or linear polarisation resistance (LPR) (Linjewile et al., 2003; McGhee, 2009), and deposition rates.
In all these cases, it is also necessary to arrange for appropriate gas and temperature monitoring around the materials. Following their removal from the plant the materials can be destructively examined to evaluate their performance. The use of pilot plants offers an alternative approach to full-scale plant exposures, with the advantages of still using real fuels, but allowing much easier and more extensive monitoring of the materials exposure conditions (Davis and Pinder, 2004). However, pilot plants are expensive to operate for extended periods and so this usually limits the lengths of the exposures to tens or perhaps hundreds of hours at most. Laboratory exposures can be carried out under much more controlled exposure conditions. However, these conditions are simulations of what happens in plants and are generally difficult to set up, with considerable care needed to relate them to the plant environments. The data generated have usually been reported in terms of mass change, but this has often proved to be misleading. Much more useful dimensional data can be used to generate datasets on metal losses, and this is now being increasingly generated following a draft EU standard method for corrosion testing (EC Project
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SMT4-CT95-2001, 2000). These tests are particularly valuable in enabling the effects of the different exposure variables to be separated out and their sensitivities determined (Saunders et al., 2002; Simms et al., 2007b). Thus, investigations of materials performance in plant, pilot plant and laboratory environments all have a role to play in determining the environmental degradation of heat exchanger materials, with each contributing to different aspects of generating the data required to understand the various processes involved.
4.4.2 Protective coatings Traditionally coatings have been used on the heat exchangers in power stations as one method of protecting the tubes from particularly challenging cases of environmentally induced degradation. A variation on the coating approach has been to use co-extruded tubes (Meadowcroft and Manning, 1983), with a highly alloyed material extruded around the outside of a standard low-alloy boiler tube material. As the thickness of the coatings that can be applied to heat exchanger tubing is limited (usually less than 2 mm), it is necessary for the coating materials to be highly resistant to the exposure environment in order for them to have a reasonable life (at least enabling the component to continue in service until the next scheduled major overhaul). Another consideration is the cost of using coatings, with prices currently of the order of 1000 euros per m2. To protect against fireside corrosion on waterwalls or superheaters/ reheaters, it has been found that highly alloyed coatings are required; in UK pulverised fuel systems, coatings of IN671 or Ni–50 %Cr have been used successfully, as have highly alloyed co-extruded tubes, such as AISI 310 (Syrett, 1987). However, such coatings are only applied to limited areas of tubing to protect against specific localised environments that have been found to have been causing accelerated damage. In contrast, waste-fired boilers have such aggressive conditions on heat exchanger surfaces that alloy 625 (Ni–20–23 % Cr–8–10 % Mo–<5 % Fe–3.15–4.15 % Nb + Ta) is often now applied to these tubes (to prevent otherwise frequent changes of the heat exchangers). For all these coating systems, it has been found that the coating quality is critical in providing adequate component life, as defects in the coatings can result in their rapid loss. The production of relatively smooth, defect-free coatings is easier in a production environment than in a power plant, though both types of locations have to be used. Weld overlay and high velocity oxygen fuel thermal spraying (HVOF) processes have both been used successfully in applying coatings to heat exchanger tubing.
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4.5
Future trends
There are currently many different pressures on the power generation industry that will influence how it develops in the near future, including: .
.
. . .
. .
A need to generate far more power worldwide to meet the needs of a growing world population and the expected economic development of many countries. IEA projections suggest 40 % more energy will be needed by 2030 (IEA, 2009). Concern over the emissions of greenhouse gases (GHG) leading to global warming and the varying desire to reduce such emissions around the world. Current EU policy is for a 20 % reduction (relative to 1990 levels) by 2020 (Dechamps, 2006) with a UK target of 80 % reduction set for 2050 (UK DECC, 2008). A desire to use more sustainable fuels, increasing the levels of biomass and waste used for power generation. Fuel supplies in terms of total availability, geographic distribution and costs. A wide range of alternative methods for generating power and heat, which are at different stages of development and with varying possibilities for successful application (including wind, wave, solar and nuclear power systems). Financial viability. Political policies and regulatory regimes that develop over time and vary with location.
As a result, there is currently considerable uncertainty about the nature of future power systems, but it is clear that they will have to be much more efficient than current systems and generate far less CO2 emissions (Farley, 2007, 2010). Thus, developments related to environmental degradation within solid fuel power systems can be grouped accordingly: .
Higher steam system temperatures. State-of-the-art pulverised fuel power plants have steam operating conditions of ~600/620 8C/240 bar (giving design efficiencies of ~46–47 %). These can be compared to many of the existing old power stations with steam conditions of ~540–560 8C/ 140–160 bar (and efficiencies of ~35–37 % with retrofitted gas cleaning systems included). European developments are heading towards higher steam operating conditions, with 650 8C and 700 8C systems, COST536 and THERMIE projects respectively (Blum and Vanstone, 2006). USA steam system developments have an even more ambitious target of 760 8C/350 bar (Shingledecker and Wright, 2006). Such plants will generate power at higher efficiencies (~60 %) and with less CO2 emitted per unit of electricity generated.
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Fireside metal temperatures are traditionally calculated to be ~30–50 8C hotter than steam temperatures, so the anticipated fireside metal temperatures are predicted to move from the 590–610 8C of existing older plants, to ~650 8C in the state-of-the-art plants and then to ~700– 810 8C in the various future systems that are being developed. From the description of fireside corrosion on superheaters/reheaters (section 4.3.3), it is anticipated that corrosion rates will increase with temperatures in line with the upward trend of the bell-shaped curve, until the reactive deposits become unstable. With increasing temperatures, this may become the life limiting degradation mechanism for some materials. For the highest steam temperatures, it will be necessary to use expensive nickel based alloys to provide sufficient creep lives. There has been little investigation of such materials in fireside corrosion conditions in the appropriate temperature range (though R&D activities are currently investigating this), but from lower temperature studies and comparison with gas turbine operating conditions, it is anticipated that similar bell-shaped curves will be observed, though with shifted peak temperatures and rates. The use of such materials would represent a major change in the construction of pulverised fuel power plants. Carbon capture systems. Three types of carbon capture systems are commonly suggested: pre-combustion capture, post-combustion capture and oxy-firing. Of these, the latter two are related to combustion plants. Post-combustion carbon capture systems are fitted downstream of the heat exchangers in a conventional type of boiler system (as a large addition to the gas cleaning systems). The efficiency penalty of such capture systems is such that viable power generation systems are going to need higher temperature steam systems to operate (see the issues outlined above). Oxy-firing systems are a distinctly different approach to carbon capture, in which the fuel is burnt in combination with oxygen and recycled flue gases. The result is the generation of gas compositions that are distinctly different from those generated in traditional combustion power plants (Fig. 4.20), though they are still oxidising with a similar O2 content compared to conventional flue gases. In this case the gas is dominated by CO2 (~60 vol %) and steam (~30 vol %) – hence the easier separation of the CO2 by condensing the water out of the fuel gas stream. However, the levels of the minor gas species (e.g. SOx and HCl) can be increased by the flue gas re-cycling, with the levels being boosted by up to 5 times depending on the configuration of the re-cycle. Again, the efficiency penalty of using this approach to CO2 generation and capture is such that higher steam temperature systems will be needed.
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4.20 Comparison between gas compositions produced in oxy-fired and air-fired pulverised fuel systems. Reproduced with permission from Maney Publishing (NJ Simms, PJ Kilgallon and JE Oakey, ‘Fireside issues in advanced power generation systems’, Energy Materials: Materials Science and Engineering for Energy Systems, 2(3), 2007, 154– 160. www.maney.co.uk/journals/ema www.ingentaconnect.com/ content/maney/ema).
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Depending on the assessments carried out (Simms et al., 2007a; Bordenet and Kluger, 2008), it is possible to anticipate significant changes to operating environments of the heat exchangers in these systems, none of which would make the conditions more benign, but these need to be checked in pilot and demonstration plants. Such work is the subject of ongoing research projects, as is the possible impact on the lives of the heat exchanger materials. Increased biomass levels in co-firing. Currently pulverised fuel power plants use up to ~10 % biomass (on an energy basis) to reduce net CO2 emissions. The use of higher levels of biomass would further reduce the net CO2 emissions, but there are many challenges associated with this, particularly in terms of fuel handling, storage, combustion (Chapter 1), as well as economic viability. In terms of the environmental degradation of heat exchangers, the main concern is for the superheaters/reheaters, where increasing the levels of biomass in a coal–biomass mix can increase the chance of aggressive chloride deposits forming and causing rapid corrosion damage (as observed in some biomass-fired power plants). To prevent this it is necessary to carefully assess the particular combinations of specific coals and biomass that could be used in order to minimise the risks involved. Improved modelling. The models that are currently available to predict
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environmental degradation in boilers are still somewhat limited, though these are now developing rapidly (Linjewile et al., 2003; Larson and Montgomery, 2006; Davis, 2010; Lant et al., 2010; Simms and Fry, 2010). The complex environments and interactions within the boilers have proved to be a longstanding challenge to researchers investigating such forms of degradation (Stringer and Wright, 1995). However, improved instrumentation and computing tools are now assisting with this. As a result, this is a topic that could greatly assist in the development of high efficiency advanced power plants that need to be built with confidence that the risks involved have been minimised. On-line corrosion/deposition monitoring. The development of on-line monitoring techniques for deposition and high temperature corrosion are ongoing R&D activities. However, the potential to use these techniques to investigate the performance of combustion plants along their hot gas paths offers valuable opportunities for detailed investigation of the operations of these power generation systems and for the validation of predictive models. In the future when such systems are more fully developed, it may even be possible to use them as part of a control strategy for new power plants to minimise the risks involved. Increased use of coatings. Some of the component operating conditions that are being proposed for future power plants many require materials that are far more resistant to the surrounding environment at metal operating temperatures of up to 800 8C. There may be issues with the cost of the materials that have sufficient creep strengths at such temperatures, and whether such materials will have adequate fireside corrosion resistance. As a result coatings may be required to protect these components from their aggressive operating environments.
Sources of further information and advice
Apart from the specific references provided throughout this chapter, further information is available from the following references, conference series and websites: . . . . .
Failures in pulverised coal boilers: French, 1993. General oxidation and corrosion: Birks et al., 2006; Young, 2008; Kofstad, 1988. Fouling, slagging and corrosion: Livingstone, 2009; Zhou et al., 2007; Tomeczek and Wacławiak, 2009. Erosion: Stringer, 1995; Finnie, 1995. Performance of materials in power plant environments: ○ Conference series: Materials for Advanced Power Engineering 1990, 1994, 1998, etc.
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4.7
European Federation of Corrosion Book Series: Numbers 14, 34, 47. Parsons Conference series: 1984, 1988, 1995, 1997, 2000, 2003, 2007.
References
Birks N, Meier G H and Pettit F S (2006), High-Temperature Oxidation of Metals, Cambridge University Press. Blomberg T (2008), ‘What are the right test conditions for the simulation of high temperature alkali corrosion in biomass combustion’, in Schu¨tze M and Quadakkers W J (eds), Novel Approaches to Improving High Temperature Corrosion Resistance, European Federation of Corrosion Publications Number 47, Woodhead Publishing, pp. 501–513. Blum R and Vanstone R (2006), ‘Materials development for boilers and steam turbines operating at 700 8C’, in Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Materials for Advanced Power Engineering 2006, Forschungszentrum Ju¨lich GmbH, pp. 41–60. Bordenet B and Kluger F (2008), ‘Thermodynamic modelling of the corrosive deposits in oxy-fuel fired boiler’, Materials Science Forum, 595–598, 261. Bradford S A (1987), ‘Fundamentals of corrosion in gases’, in Metals Handbook Ninth Edition, Volume 13: Corrosion, ASM International. Cain C and Nelson W (1961), ‘Corrosion of superheaters and reheaters of pulverized coal-fired boilers’, Trans ASME, J Engng Power, 83 (Series A), 468. Davis C (2010), ‘Impact of oxy-fuel operation on corrosion in coal fired power plant’, http://www.specialmetalsforum.com/uploads/docs/12754067434.EON EngineeringOxyfuelCorrosion.pdf [accessed 4 November 2010]. Davis C J and Pinder L W (2004), Fireside Corrosion of Boiler Materials – Effect of Co-Firing Biomass with Coal, UK Department of Trade and Industry, Report No. COAL R267 DTI/Pub URN 04/1795. Dechamps P (2006), ‘The EU research strategy towards zero emission fossil fuel power plants’, in Lecomte-Beckers J, Carton M, Schubert F and Ennis P J (eds), Materials for Advanced Power Engineering 2006, Forschungszentrum Ju¨lich GmbH, pp. 25–40. EC Project SMT4-CT95-2001 (2000), Draft Code of Practice for Discontinuous Corrosion Testing in High Temperature Gaseous Atmospheres, TESTCORR, ERA Technology, UK. Farley M (2007), ‘Clean coal Technologies for power generation’, in Strang A, Banks W M, McColvin, G M, Oakey J E and Vanstone R W (eds), Parsons 2007: Power Generation in an Era of Climate Change, IoM Communications, pp. 335– 342. Farley M (2010), Overview of Capture Technologies for Pulverised Coal-Oxyfuel and Post Combustion Capture, Doosan Power Systems. Available from: http://www.specialmetalsforum.com/uploads/docs/12754049682.DoosanMF NamtecHarrogate2010.pdf [accessed 4 November 2010]. Finnie I (1995), ‘Some reflections on the past and future of erosion’, Wear, 186–187, 1–10. Foster D J, Livingston W R, Wells J, Williamson J, Gibb W H and Bailey D (2004),
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Stringer J and Wright I G (1995), ‘Current limitations of high temperature alloys in practical applications’, Oxidation of Metals, 44, 265–308. Stringer J (1995), ‘Practical experience with wastage at elevated temperatures in coal combustion systems’, Wear, 186–187, 11–27. Syrett B C (1987), ‘Corrosion in Fossil Fuel Power Plant’, Metals Handbook Ninth Edition, Volume 13: Corrosion, ASM International. Tomeczek J and Palugniok H (2002), ‘Kinetics of mineral matter transformation during coal combustion’, Fuel, 81, 1251–1258. Tomeczek J and Wacławiak K (2009), `Two-dimensional modelling of deposits formation on platen superheaters in pulverized coal boilers', Fuel, 88, 1466–1471. Tomeczek J, Palugniok H and Ochman J (2004), ‘Modelling of deposits formation on heating tubes in pulverized coal boilers’, Fuel, 83, 213–221. UK Department of Energy and Climate Change (2008), Climate Change Act 2008. Wessel R A and Righi J (1988), ‘Generalized correlations for inertial impaction of particles on a circular cylinder’, Aerosol Science and Technology, 9, 29–60. Wood N B (1981), ‘The mass transfer of particles and acid vapour to cooled surfaces’, Journal of the Institute of Energy, 76, 76–90. Young, D (2008), High Temperature Oxidation and Corrosion of Metals, Elsevier. Zhou H, Jensen P A and Frandsen F J (2007), ‘Dynamic mechanistic model of superheater growth and shedding in a biomass fired grate boiler’, Fuel, 86, 1519– 1533.
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5 Creep in boiler materials: mechanisms, measurement and modelling V . S K L E N I Cˇ K A a n d L . K L O C , Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Czech Republic
Abstract: Creep is one of the most important factors determining the useful life of boiler structures in power plants. The chapter first discusses the physical basis of the creep processes pointing out the key role of microstructural evolution and stability of creep-resistant materials. Then creep measurement methods are reviewed and the results of some experiments to simulate the effects of the operating environment are described. Finally, the various methods of predictive modelling of the creep processes and their advantages and weaknesses are discussed. Key words: creep deformation and damage mechanisms; microstructural changes and stability, creep measurement methods, creep life; creep predictive models.
5.1
Introduction
Creep is defined as plastic deformation occurring in a material subjected to a constant applied stress and/or constant load (Evans and Wilshire, 1985; Cˇadek, 1988; Abe, 2008). In practice, this flow becomes especially important at elevated temperatures, typically above ~ 0.4Tm, where Tm is the absolute melting temperature of the material, because diffusional processes then occur fairly rapidly. During the last two decades, great progress has been made in design and developing boiler creep-resistant materials of high strength and corrosion resistance necessary for an improvement of the thermal efficiency and the reduction of emissions of environmentally hazardous gases of new thermal power generation plants (Hald, 2006; Mayer and Masuyama, 2008; Viswanathan et al., 2009). In the field of thermal power generation, the maximum allowable temperature was about 565 8C for conventional low-alloy ferritic steels. Progress in recent years has led to the design and development of high-strength 9–12 % chromium steels 180 © Woodhead Publishing Limited, 2011
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capable of operating in ultra super critical (USC) power plants at metal temperatures approaching up to 650 8C. The creep strength of austenitic creep-resistant steels has been enhanced to enable operation up to temperatures of 675–700 8C through development of high chromium, high nickel steels. Creep behaviour and degradation of creep properties of creep-resistant boiler materials are phenomena of major practical relevance, often limiting the lives of boilers designed to operate for long periods under stress at elevated temperatures. Thus successful plant operators should rely on detailed boiler material understanding and component quality checks to ensure safe and reliable operation of their plant. Since life expectancy is based on the ability of the material to retain its high temperature strength for a period longer than that of the projected design life, methods of creep property assessment based on physical changes in the material are necessary. The understanding of creep mechanisms and characteristics of boiler steels will help plant operators and service providers to operate the plant more efficiently and reduce forced outage periods and associated risks. This chapter attempts to highlight the creep problem areas just in this respect. Physical models in the prediction of the creep–microstructure property relationship are reviewed and the proposed approaches are illustrated by our recent experimental results on selected advanced creepresistant boiler materials.
5.2
Creep deformation and damage mechanisms in boiler materials
5.2.1 Creep deformation mechanisms Creep also occurs at stresses below the yield stress of the particular material. The rate of the creep process is strongly temperature dependent. Thus, the thermally activated processes control the deformation. The creep strain and creep damage of the material become important, if the material is loaded at elevated temperatures and should withstand the conditions for a long time. This is a typical situation in thermal power plants and similar high temperature facilities. The most important mechanism of plastic deformation in metals and alloys is dislocation glide, but activities on grain boundaries play an important role, too. The possible creep deformation mechanisms are: . . . .
conservative motion of dislocations (glide), non-conservative motion of dislocations (climb), grain boundary sliding, stress directed diffusion of matter between grain boundaries,
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the last mechanism representing the creep damage development, the topic of the second part of this section. Dislocation movement Though the stress is not high enough for free dislocation glide, the dislocations can be released from the obstacles by thermal activated processes and then glide until they meet another obstacle. The strain rate e_ is described by the Orowan equation e_ ¼ brm v
½5:1
where b is the length of the Burgers vector, ρm is the density of mobile dislocations and v is their mean velocity, involving the waiting time at obstacles. Nevertheless, both ρm and v variables are difficult to evaluate. The more effective that obstacles are for dislocation glide embedded in the microstructure, the better is the creep strength that is obtained. There are three types of obstacles: 1. 2. 3.
dislocations nodes and other dislocations (work strengthening), particles of second phase (precipitation strengthening) and solute atoms (solid solution strengthening).
During deformation, dislocation sources are activated and the dislocation density increases, causing immobilisation of dislocation segments and work strengthening. Nevertheless, the opposite process of dynamic recovery, that is decrease in the dislocation density, is running at elevated and high temperatures. It is generally accepted that these two processes tend towards some kind of dynamic equilibrium, in which the dislocation structure remains stable and the creep rate constant, leading to the so-called steady state creep. The steady state creep rate defined by the equilibrium between the rate of generation of dislocations and the rate of recovery is a variable of high importance for models of dislocation substructure development. Since the other processes (precipitation and solid solution strengthening, damage development) may proceed during this state, it does not necesarilly correspond to the constant creep rate. In materials like boiler steels, where the precipitation strengthening plays an important role, it is not possible to figure out if and when the steady state is reached from the creep curve. Setting the steady state and the minimum creep rate equal to each other is a frequently found mistake. In most cases, the development of the dislocation substructure leads to subgrain formation. Various composite models were developed to describe the creep behaviour of the structure (Blum et al., 1989; Vogler and Blum,
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1990; Henes et al., 1999; Blum, 2008). In general, dislocation structure development is controlled mainly by dislocation climb, which is controlled by self-diffusion of the base metal atoms. Pipe diffusion along dislocation cores is more important at lower temperatures, while at temperatures close to the melting point lattice diffusion prevails. There are several mechanisms that show how dislocations can overtake particles of second phase. If the particles are coherent with the matrix, particle cutting by dislocations can occur (Kocks et al., 1975). Another possibility is the Orowan looping, where the dislocations pass through the gap between particles, leaving a loop around the particle. The loop increases the effective diameter of the particle for the next dislocations moving on the same slip plane, and can annihilate by climbing around the particle. Dislocation segments can also climb to another slip plane around the particles. A model of the processes was presented by Ansell and Weertman (1959). However, the model is not capable of explaining apparent activation energy of creep much higher than that of self-diffusion, which is frequently observed with structural materials. To account for this, the concept of back stress was introduced by Davies et al. (1969) and many others. In this case, where the disordered non-coherent interface between the particle and the matrix is attractive to dislocations, the dislocation detachment from the particles is the main controlling process. This mechanism was analysed by Ro¨sler and Arzt (1990). Since the phase structure of the material is not in thermodynamic equilibrium, it is evolving in time at high temperatures. Particle coarsening, dissolving of metastable phases and precipitation of new ones are the most important processes. The kinetics is controlled mainly by solute atoms diffusing along dislocation cores, grain boundaries and through a lattice, depending on conditions. Though there is a strong interaction between the dislocation substructure and particle substructure, the kinetics of their development can be quite different. The shape of the creep curve results from the combination of these different processes. Thus the analysis of the creep curve shape without microstructural studies cannot resolve the different processes, and the extrapolation methods based on such simple analysis cannot be reliable. During the processes described above, both local and general concentrations of solute atoms may change, also effecting the solid solution strengthening. Kimura et al. (1996) have introduced the concept of inherent creep strength, corresponding to the state of material close to the thermodynamic equilibrium, where all temporal strengthening mechanisms are exhausted.
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5.1 Dependence of the contribution of grain boundary sliding to total creep strain on the minimum creep rate for various materials.
Grain boundary sliding When shear stress is applied to the grain boundary, tangential displacement of the adjacent grains can occur causing a strain. This process is known as grain boundary sliding and plays an important role in the creep deformation of polycrystalline materials (Langdon, 2006). The relative contribution of grain boundary sliding can reach values of 30 % or more and is growing with decreasing strain rates and grain sizes, as shown in Fig. 5.1. Since the accelerated laboratory experiments run with higher creep rates than those that occur in industrial practice, the role of the grain boundary sliding in creep is usually underestimated. When the grain boundary sliding occurs as an independent deformation mechanism (Rachinger sliding), it exerts large local strain concentrations at grain boundary irregularities (ledges, precipitates) and on triple points. These stresses may be relaxed by intragranular dislocation slip, localised plastic deformation at triple points (fold formation), grain boundary migration or by formation of grain boundary cavities and/or wedge cracks at triple points (Cˇadek, 1988). Thus the grain boundary sliding process may be crucial for the grain boundary cavity nucleation and growth and for the intergranular creep damage development (Sklenicˇka, 1998). Since the atomic structure of grain boundaries is generally complex and not well known, the mechanism of the grain boundary sliding is also not known. High temperature microstructural evolution and the associated creep property prediction could well be the next major application of
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molecular dynamics (MD) modelling, which goes back to first principles associated with interatomic potentials (Faulkner, 2008). Diffusional creep Normal stress on grain boundaries changes the energetic state of vacancies adjacent to those boundaries, provoking diffusional flux of vacancies between boundaries with different orientations to the applied stress axis. Clearly, the matter is transported in the opposite direction, causing deformation of individual grains (Fig. 5.2). The diffusional flux can go through the grain interior or through grain boundaries itself depending on conditions, mainly temperature. The former case is called Nabarro–Herring creep while the latter is Coble creep (Cˇadek, 1988). The rate of the diffusional creep e_ d can be expressed as Os pDB dB DL þ e_ d ¼ B 2 ½5:2 d kT d where B is the numeric constant derived from the grain shape, Ω is atomic volume, σ is applied stress, d is grain size, k is the Boltzmann constant, T is absolute temperature, DL is the lattice diffusion coefficient, DB is the grain boundary diffusion coefficient and δB is the effective width of the grain boundary for diffusion. The first term in parentheses represents the contribution of Nabarro–Herring creep and the second term the contribution of Coble creep. In contrast to the previous deformation mechanism, the diffusional creep theory is well established and clear, but the role of the process in the creep of structural materials is disputed and not clearly resolved. The strain rate of the diffusional creep depends strongly on grain size; this feature is used to identify the process experimentally. This was successful in some model materials (Kloc et al., 1999), but in structural materials it is not possible to
5.2 Diffusional creep mechanism.
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change grain size considerably without changing many other structural parameters. In any case, diffusional creep should be considered important mainly with fine grained materials. Diffusional creep deformation must be accommodated by grain boundary sliding (Lifshitz sliding), which has a slightly different character from that described above. In this case, no stress concentrations on triple points are generated, so the diffusional creep does not cause damage nucleation and can be considered harmless for the material. Transitions of creep deformation mechanisms The dislocation movement mechanism prevails under conditions of accelerated laboratory creep tests and thus it is the subject of main interest. Nevertheless, industrial application conditions lead to much lower strain rates, where deformation mechanisms based on grain boundaries (sliding, diffusional creep) tend to become more important. Qualitative change in creep behaviour, mainly the stress and temperature dependencies of the creep rate, can occur. In this case, predictions of the creep life based on extrapolations may fail (see section 5.5). Such a change in creep behaviour has been found with the P91 type of steel at approximately 100 MPa and 600 8C by Kloc and Sklenicˇka (2004) and by Haney et al. (2009). The results are based on a change in slope of the minimum creep rate dependence on stress and were disputed as nonconvincing due to hesitation if the minimum creep rate was reached.
5.3 Time to fracture dependence on stress for P91 steel compared to extrapolation by the SHC committee. Taken from Kimura et al. (2010).
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Recently, the results of extremely long-term creep tests were published by Kimura et al. (2010), providing direct proof of the change also at time-torupture dependendence on stress (see Fig. 5.3, reproduced from the cited work). Revision of current design codes for the creep life of the steel is required on that basis.
5.2.2 Creep damage mechanisms and fracture During the deformation at high temperatures a number of micromechanistic processes occur that damage the structural integrity of the boiler materials and are ultimately responsible for failure (Evans, 1984; Riedel, 1987; Cˇadek, 1988; Dyson, 1988; Maruyama, 2008; Viswanathan and Tilley, 2008). Here we concentrate on intergranular cavitation damage, which constitutes one of the most complex topics, because intergranular cavitation and fracture in general are a consequence of the joint action of creep deformation and diffusional processes, and/or environmental corrosion (Beere, 1981; Sklenicˇka, 1998). The creep cavitation depends strongly on material and loading conditions; in some creep-resistant materials cavities already appear in early stages of loading, whereas other cavitate in an observable amount only towards the end of their creep life tf so that the life fraction t/tf at which cavitation damage can be detected is variable. Nevertheless, some features of cavity accumulation in the fracture state are common to all situations and form a basis of the lower bound of creep life prediction (Sklenicˇka, 1999). Consequently, estimation of the degree of cavitation should be the starting point of any remanent life assessment route. In general, intergranular creep fracture involves the nucleation of intergranular cavities, their growth and linking to form microcracks and the eventual propagation of a main crack to final fracture (Fig. 5.4). For a long loading time the development of intergranular creep damage takes place in mutually independent microvolumes. Independently of the mechanisms of cavity nucleation and especially growth that are operating, sooner or later the coalescence of cavities starts to take place. In the course
5.4 Scheme of intergranular creep damage development.
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of coalescence, which already represents an advanced stage of intergranular damage, the local stress concentrations are not steady any more; depending on their stress sensitivities the operating processes speed up. The inhomogeneity of spatial distribution of isolated cavities leads to inhomogeneity of more advanced forms of creep damage, such as the coalesced cavities mentioned above and later grain boundary microcracks. Both of these types of intergranular damage appear first in the nearest vicinity of the free surface and their density predetermines the place where the last stage of the fracture process will start. This stage is characterised by extensive crack formation due to long-range coalescence of damage also inside the creeping body; these cracks may join surface cracks. The final fracture then takes place by a relatively quick propagation of one of the long surface cracks – the ‘main’ crack. We are still some distance from fully understanding the mechanisms by which materials fracture in creep. One of the main uncertainties is associated with quantitative characterisation of the critical extent of cavitation necessary to promote the main crack growth. In the past few decades various methods have been proposed for the qualitative and quantitative assessment of creep cavitation, such as metallography, ultrasonics, small angle neutron scattering, densitometry, hardness measurements, electrical resistivity and so on (Viswanathan and Tilley, 2008). So far, mainly metallographic methods, e.g. the replicas technique (Neubauer and Wedel, 1983), are widely and successfully used for the inspection of service exposed components of high temperature plant operating within the creep range. For a more detailed investigation of intergranular damage progress, it is necessary to know not only the total cavity volume (VV)T but also the separate contributions of the individual forms of cavitation. The total volume of intergranular cavities is given by ðVV ÞT ¼ ðVV ÞC þ ðVV ÞCS þ ðVV ÞCR ;
½5:3
where (VV)C, (VV)CS and (VV)CR represent volume fractions of isolated cavities, areas of cavity interlinkage (coalescence) and microcracks, respectively. Neubauer and Wedel (1983) and Neubauer (1984) proposed the characterisation to categorise creep damage and remaining life on a largely qualitative basis depending on the extent of cavities and microcracks observed during plant shutdowns. These authors considered that creep cavitation damage could be classified on the basis of a five-category scale, ranging from no detectable cavitation damage to severe damage in the form of macrocracks located at many grain boundaries. Assignment of the appropriate observed damage rating was determined on the basis of
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comparison of replicas prepared from service exposed components, with reference sets of micrographs illustrating the various levels of damage for the material in question. However, the field replication data forming the basis of the Neubauer–Wedel recommendations have shown no consistent trends in cavitation evolution with operating time or with calculated creep life fraction consumed (Viswanathan and Tilley, 2008). More field data are needed before a clear correlation can be established between replica results and cumulative creep damage. Subsequently, a more quantitative approach to creep damage accumulation has been developed to predict remaining life. Two damage parameters have been proposed to quantify cavitation damage (Cane and Shamma, 1984; Riedel, 1989). Both parameters are based on constrained cavity growth (Dyson, 1988), although their micromechanical backgrounds are different. The A parameter is a mechanistic measure for the continuum damage parameter ω in the classical Kachanov–Rabotnov treatment (Rabotnov, 1969). The A* parameter, which has been proposed by Eggeler et al. (1990) and Eggeler (1991), represents an alternative approach to creep damage quantification for creep life analysis. Its micromechanical background is an extension of the A parameter theory. Finally, the characterisation of the damage state by means of the A parameter and another parameter obtained by area counting rather than line counting has been proposed by Stamm and von Estorff (1993). A statistical model based on a Poisson–Voronoi tessellation (Stoyan et al., 1987) for the simulation of the austenitic grain structure is applied to derive an expression relating to the experimentally determined crack density parameter. It should be noted that the modelling and the analysis of creep damage development as described above concentrate on the average effects and ignore the non-uniform local accumulation of creep damage. Thus, the modelling of creep cavitation development by the phenomenological damage parameters should be complemented by a more detailed quantitative metallographic analysis based on stochastic geometry and stereology. The final stage of intergranular creep fracture occurs due to the critical decrease of the grain boundary cohesion caused by the accumulation of cavitation damage in the whole specimen volume. For this critical extent of cavitation the term ‘ultimate stage of damage’ has been proposed (Sklenicˇka, 1998). It can be defined as a state of damage at which the threshold probability of damage interaction over distances comparable with the specimen cross-section dimensions is attained. The results of the measurements of quantitative damage parameters obtained for various creep-resistant materials under a broad range of creep testing conditions (Sklenicˇka, 1998) are summarised in Fig. 5.5. From
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5.5 Fracture values of the areal fraction of cavities related to cavitated grain boundaries, A~c for various creep-resistant materials (d is grain size).
Fig. 5.5 it follows that the ‘ultimate state of damage’ (fracture state) can be characterised by the critical value of the areal fraction of cavities in cavitated grain boundaries A~c ¼ Y ¼ 0:270:4
½5:4
Hence it can be seen that this value depends neither on creep conditions (temperature, stress) nor on the material investigated. At the same time, the value A~c ¼ Y represents a universal geometric criterion of intergranular creep fracture. To explain theoretically the critical value of A~c ¼ Y the bond percolation theory (Hammersley and Welsh, 1980; Essam, 1989) can be used. It should be pointed out that the percolation model assumes spontaneous propagation of a main crack as soon as a sufficiently cavitated fracture path is formed across the specimen cross-section (Sklenicˇka, 1998). However, a gradual propagation of the main crack due to successive growth and interlinkage of cavities nucleated ahead of its tip requires a considerably lower value of A~c (Sklenicˇka et al., 1990). Since the percolation theory can be applied only in the case of a crack passing continuously without stopping through numerous coalesced cavity clusters distributed along its propagation direction, the limited applicability of the percolation theory in the case of fatigue failure can be expected (Sklenicˇka et al., 1993). Creep damage mechanics has been developed as a continuum approach to the analysis of creep fracture. The damage is measured by the scalar parameter ω, which varies from zero (no damage) to one (fracture). One
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interpretation of the Kachanov parameter ω is the decrease in effective load bearing area due to cavitation. In a metallographic assessment, this would correspond to the area of transverse grain boundaries occupied by cavities. Considerable success has been obtained in creep life prediction by assuming that the damage parameter ω is directly analogous to the A parameter measure of creep cavitation (Furtado and LeMay, 1997). However, one should not lose sight of the fact that, where ω is a continuum variable, A describes a mechanism that is discrete in both space and time. To treat this problem, some attention needs to be given to the difference between continuum and discrete structural representation of a material. Brear et al. (1998) have proposed modification to the A and A* parameters in order to allow quantification of damage inhomogeneity, by comparison with the binominal statistics expected under a uniform damage hypothesis. The main uncertainties used to understand creep fracture fully are associated with (i) identification of the processes responsible for cavity nucleation and (ii) quantitative characterisation of the critical extent of cavitation necessary to promote the main crack growth. At present, it is generally believed that cavities are nucleated due to grain boundary sliding at geometrical irregularities (suitable types of inclusion particles are frequently considered) on grain boundaries, where high local stress concentrations could develop. The association of cavities with secondphase particles (presumably M23C6 carbides) in grain boundaries in creepresistant steels has been frequently observed. The nature of cavity nucleation sites is of considerable technological interest. If cavities do nucleate only on non-adherent grain boundary inclusions in steels, then by control of residuals and proper heat treatment one expects to improve dramatically the creep properties. At present, cavity growth processes are understood much better than the process of cavity nucleation. There are two limiting kinds of cavity growth in creep: unconstrained and constrained cavity growth (Dyson, 1988). The unconstrained cavity growth takes place if cavities are present at all grain boundaries. On the other hand, the constrained cavity growth takes place if cavities are present only on isolated grain boundaries. Cavities can grow by various combinations of such processes as grain boundary diffusion, surface diffusion and power-law creep. However, for creep-resistant materials under typical service conditions, pure diffusional growth is probably less likely to occur due to the constraints applied by the less compliant material surrounding cavitated regions.
5.3
Measurement methods
Creep properties of materials are obviously tested in pure tension at constant temperature and constant load (or constant stress, see below).
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Besides this standard testing procedure, many non-standard techniques have been developed for some special purposes.
5.3.1 Standard creep testing methods The principle of creep testing is very simple: the specimen of material is held at an elevated and stabilised testing temperature and loaded by tensile force, while the strain is recorded in time. It should be noted that the so-called ‘creep rupture tests’ without strain measurements are obsolete, since reliable systems for strain measurement and recording are available, enabling much more information to be obtained from the test. The specimens used to have a cylindrical gauge section with heads for machine grips and possibly some collars for an extensometer attachment. In some cases, flat specimens with a rectangular cross-section are also used. The force is usually applied by weight through a lever system, providing excellent stability and reliability. Another possibility is using a standard tensile machine controlled by loop-back from the stress sensor. This solution appears more flexible if the stress changes are demanded, but the gravity based system provides better stability. Another disadvantage of the stress controlled tensile machine occurs in the case of accidental blackout, which may happen during long-term testing. In this case, the machine is not able to adopt the tensile train shrinkage as the temperature goes down, resulting in severe overloading. This may destroy the specimen or the stress sensor; at least, the microstructure of the material is changed, which makes the test unusable. In contrast, the gravity driven system maintains a constant load in such a case, enabling the test to be continued. For the sake of test results comparability, the testing equipment and procedures are ruled by national and international standards, listing demands for the temperature accuracy and stability, load stability, strain measurement precision and so on. The most important creep testing standards to date are ASTM E139 (2006) and ISO 204 (2009). The conventional creep tests occur in two ways: as constant load tests and constant stress tests. In the constant load tests, the applied force is constant throughout the whole test. Since the cross-section of the specimen reduces as the elongation proceeds, the applied stress grows slowly. The constant stress tests account for the effect of maintaining the stress constant by lowering the applied force that depends on the strain. This is usually achieved by a special mechanism on the loading lever; in the case of a stress controlled tensile machine this is simply due to the controlling software. The constant stress tests are based on the assumption that the volume of the material remains
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constant, leading to the applied force expressed as F¼
S0 s expðeÞ
½5:5
where F is the applied force, S0 is the initial cross-section, σ is stress and ε is true strain, defined as ε = ln(l/l0). The assumption is not strictly fulfilled due to cavity formation in advanced stages of the creep process, but is generally accepted. The constant load tests correspond to the application conditions, while the constant stress tests are important as a data source for finite element modelling of more complex stress situations. For some materials, the difference between the two types of test is very small, while for other materials the difference is large, depending on the creep plasticity of the material under testing.
5.3.2 Stress relaxation testing Conventional creep testing suffers from two main weaknesses: (i) it takes a long time to obtain results and (ii) it generates results with wide scatter. Therefore, other testing techniques have been proposed, such as the stress relaxation technique (SRT). Stress relaxation tests may provide an important information on the creep behaviour of the tested material (De Bruycker et al., 2009). By keeping the strain constant and letting the stress relax over time at a certain temperature, the minimum and/or steady creep rate could be attained. However, it should be noted that the stress relaxation tests are the short-term ones (usually less than one week of testing) compared to standard creep tests. If there are slow time dependent microstructural processes like coarsening of precipitates, which would occur within years, the relaxation test cannot measure the minimum creep rate that would occur after a few years. The application of the stress relaxation tests to provide an independent measure of embrittlement of material has been proposed by Woodford (2004).
5.3.3 Stress concentrations and multiaxial stress testing In real high temperature components, the pure tensile loading is rather exceptional and design stress concentrators occur frequently. Creep tests using the same equipment as described above, but notched specimens instead of smooth ones, are employed to address the creep behaviour of the material under such circumstances. The V-shaped notches are used to mimic the stress concentrators, while blunt or semicircular notches are used to evaluate the response on multiaxial loading. These tests are also standardised in ISO 204.
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Creep can be also studied on tubes loaded by internal pressure, which is the type of loading in boiler structures, but these experiments meet serious problems concerning loading medium and safety in the case of fracture, as well as with precision strain measurement (Yaguchi et al., 2009). Yet another possibility is a tubular specimen loaded by combined tension and torsion (Kowalewski, 1995).
5.3.4 Creep fracture growth measurements In this type of testing, the stress situation in front of the creep fracture tip is simulated. Though it is similar to the notched specimen described above, a higher degree of stress concentration and triaxiality is demanded, which cannot be satisfied by the axisymmetric shape of the specimen. For that reason, flat specimens with a sharp notch on one or both sides are used, but the use of standardised ‘C’ specimens for fracture toughness measurements seems to be the best. The experiments concentrate on the final stages of the creep process, when the main crack is growing through the material, so it is sometimes called an ‘accelerated creep test’. Though this may bring important information concerning the final stage of the creep, the material did not undergo the whole creep process and its microstructure may be quite different from that in long-term creep. Thus, this type of test cannot replace the standard creep testing experiments. Nevertheless, it seems worthwhile to do such tests on service exposed materials.
5.3.5 Creep testing under non-steady conditions Since the loading conditions in service are not as constant as that of laboratory experiments, the influence of stress and/or temperature changes on the creep life of the material is very interesting for engineering practice (Berger et al., 2008; Sklenicˇka et al., 2008). Such experiments may be conducted on slightly modified standard creep testing equipment. It is important to compare the results to that of standard smooth creep tests. It is also important to distinguish the direct effect of non-steady conditions on the microstructure and its resistance to the creep strain from side effects caused mainly by thermal stresses generated during temperature changes. In laboratory experiments, the thermal stresses are much smaller than those occurring in plant parts, since the volume of the material is much lower and the cooling or heating rates are lower.
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5.3.6 Measurements of very low creep strains Demand for safe creep lives of industrial parts is frequently 30 or 40 years, leading to a mean strain rate in the order of 1012 s1. Conventional mostly short-term creep tests are not able to measure creep properties relevant for such conditions. Special creep testing techniques have been developed to measure very small creep strains, enabling creep effects to be measured under conditions corresponding to the industrial service of the material. The techniques employ mechanical springs, which are useful for creating loading modes that connect relatively large displacements to small strains within a material’s elastic region. The same loading modes and specimen geometry can be used to measure very small creep strains. Thin plate bending, long rod torsion or helicoid spring deformation provide ‘magnification’ of strain for relatively easy measurement (Kloc and Marecˇek, 2009). Of course, the methods allow investigation of the early stages of the creep process only due to time limitations. For boiler materials, the helicoid spring specimen technique is most convenient, since the helicoid specimens can be directly machined from the boiler tubes. However, these techniques require a special experimental facility. Experiments reveal changes in creep deformation mechanisms, which are able to make common extrapolation methods useless (Kimura et al., 2010).
5.3.7 Special techniques for residual creep life assessment These techniques were developed to provide creep tests of material taken from power plants during periodic inspection without a strong intervention into the inspected structure. The methods enable a detailed evaluation of local properties in very small volumes and seem to be promising for the investigations of these properties in different weld zones. Two creep testing techniques are considered to meet these demands: small punch creep tests (Baik et al., 1983; Parker and James, 1994; Dobesˇ and Milicˇka, 2001) and impression creep tests (Hyde et al., 2009). Small punch creep tests In this technique, miniaturised discs are detracted from the surface of a tested part using on-site sawing or electric discharge sampling equipments. Then the disc is mounted on a supporting ring and penetrated by a small punch, usually a ceramic ball, at constant force and elevated temperature. The deflection in the centre of the specimen is recorded. The test is finished by a complete perforation of the disc, which corresponds to the creep fracture. Since the stress and strain state is rather complex compared to the
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conventional tensile test, it is not easy to compare the results of a small punch test to that of the standard test. Another problem is that the microstructural state of the surface layer is not representative of the bulk material. Nevertheless, the technique is considered as a prospective one and probably will be standardised soon. Impression creep In this technique, a hard rectangular indenter is impressed into the material heated to the required temperature, and the movement of the indenter is recorded in time. It is even possible to make such tests directly on the plant site, but more frequently it is done on small plate specimens cut from the part surface. The technique possesses similar disadvantages to that of the small punch technique, but the compression stresses prevail. Thus the processes important for tertiary stage creep, like cavity growth and coalescence, cannot be revealed. This technique will probably also be standardised soon.
5.3.8 Measurement of the synergistic effects of other degradation processes Besides the creep, industrial components like boilers are subjected to other degradation processes like thermo-mechanical fatigue, high-cycle fatigue, and corrosion and oxidation. It is important to know the effect of all these concurrent processes. Standard creep tests can be modified to study these synergistic effects. Many experiments have been done on the interaction of creep and high cycle fatigue by superimposing vibrations on the loading system of the creep machine (Vasˇ ina et al., 1995), but this interaction is more important for the turbine blade materials than for boiler materials. The creep machine could be equipped with a special chamber for implementation of various corrosive gases on the specimen during the creep test; these experiments are technically challenging since it is not easy to maintain precision of the loading train and accuracy of the strain measurement system during long-term exposure to a corrosive atmosphere. On the other hand, corrosion is one of the main degradation processes in boiler materials, mainly when the biomass is used as a fuel for the plant, so knowledge of its influence on creep behaviour of the material is of very high importance. As pointed out earlier in section 5.3.5, the creep experiments with temperature changes are not capable of simulating thermo-mechanical fatigue under real industrial conditions due to the low material volume and relatively slow changes of temperature. Thus the experimental investigations of interaction between creep and thermo-mechanical fatigue need massive
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specimens and/or extremely high rates of heating and cooling, which are hard to achieve under laboratory conditions (Berger et al., 2008).
5.4
Effect of operating environment
To assess the behaviour of a complex boiler component when subjected to creep loading it is important to determine the creep properties of the material under conditions representative of those likely to be seen in service. One of the techniques widely used for life assessment of high temperature components involves creep tests or stress rupture tests under accelerated conditions. Acceleration is achieved by increasing the stress or temperature or both (Woodford, 1993; Viswanathan and Tilley, 2008, p. 637; Wilshire and Bache, 2009, p. 44). It should be noted that the designed lives of power plant components are based on material properties that are usually those of virgin materials together with specified operating conditions. However, the actual material properties may differ due to ageing and stress variation, temperature variation and the combined effects of these variables. Periodic applications of excessive loads and intermittent heating are normally ignored when the material is evaluated. Thus, it is important to realise that the efforts to obtain high quality databases of relevant creep properties for design at high temperatures will be costly and time consuming. Accordingly, it will be necessary to make the most appropriate use of sophisticated testing methods where possible. Also, modelling methods (section 5.5) that predict material performance in creep will enable testing to be targeted on the most critical conditions, thereby reducing the overall amount and cost of creep tests. The first part of this section deals with the creep property deterioration of selected advanced ferritic creep-resistant steels during long-term isothermal ageing, which can be related to microstructure changes arising from such exposure. The second part deals with the effect of intermittent heating on creep behaviour to simulate non-steady loading conditions in service.
5.4.1 Creep behaviour of advanced boiler steels after longterm isothermal ageing In the creep literature the importance of correlating microstructural changes occurring in materials during service exposure with their resulting creep properties has been emphasised, since it is those changes in microstructure that are directly responsible for the observed changes in material strength due to temperature exposure (Sklenicˇka et al., 2003; Viswanathan and Tilley, 2008; Abe, 2009; Fujiyama et al., 2009; Kimura et al., 2009; Panait et al., 2010).
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5.6 Stress dependences of the times to fracture for P91, P92 and P23 steels.
In order to accelerate some microstructural changes and thus to simulate degradation processes in long-term service, isothermal ageing at 650 8C for
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10 000 h was applied to P91, P92 and P23 steels in their as-received states. The accelerated tensile creep tests were performed at a temperature of 600 8C on all steels in both the as-received state and after long-term isothermal ageing, in an effort to obtain a more complete description of the role of microstructural stability in high temperature creep of these steels (Sklenicˇka et al., 2010). For all steels under investigation, the double logarithmic plots of the time to fracture tf as a function of the applied stress are shown in Fig. 5.6. It is clear from these plots that the creep lifetimes of steels in their as-received states are considerably longer than that in the aged state. While the difference for steel P91 is independent of the applied stress, for P92 and P23 steels these differences considerably decrease with increasing applied stress (for P92 steel there is a tendency at the higher stresses towards no effect of the state of the steel on the lifetime). From the above results, it can be suggested that observed differences in the creep behaviour may be explained by taking account of the change in the microstructure during isothermal ageing and creep. Based on experimental investigations the P91 and P92 steels (Figs 5.7 and 5.8) exhibited the following microstructure components. The prior austenite grain boundaries are decorated by a network of M23C6 particles. Particles of this phase are also observed on subgrain boundaries as well as inside subgrains. Some coarsening of these chromium carbides was observed after ageing. Further, there are tiny precipitates of minor phases (MX), located both on grain boundaries and on former martensitic laths. Careful examination also revealed very small particles inside the former martensitic laths. By comparison with the as-received states, however, the subgrain structure evolution in the aged state exhibits growth of the subgrain size. More intensive growth of subgrains was found in the gauge length of the crept specimens, which is supported by the influence of the applied stress and/or strain and is more pronounced than the subgrain growth in the specimens’ heads subjected only to the influence of temperature. Similar results of the quantitative measurement of subgrain size in P91 steel in the as-received state and after long-term ageing and creep have been reported by Panait et al. (2010). Consequently, a decrease of the creep resistance of the aged P91 and P92 steels in comparison with their as-received states may be explained to some extent by taking account of the change of dislocation substructure during isothermal ageing. At present it is believed (Seung et al., 2006; Yoshizawa et al., 2009; Panait et al., 2010) that the long-term creep strength loss of 9%Cr creep-resistant steels is caused mainly by precipitation of the Laves phase. The solid solution strengthening effect by Mo atoms in P91 steel (the chemical composition of P91 steel does not contain W) was found to be weak so that precipitation of the Laves phase should not strongly affect the creep resistance of this steel. By contrast, the creep
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5.7 TEM micrographs showing the microstructure of P91 steel after ageing at 650 8C for 10 000 h and creep at 600 8C and 125 MPa: (a) head of creep specimen, (b) gauge length, (c) and (d) precipitates in P91 (TEM extraction carbon replica).
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5.8 TEM micrographs showing the microstructure of P92 steel after ageing at 650 PC for 10 000 h and creep at 600 8C and 150 MPa: (a) head of creep specimen, (b) gauge length, (c) and (d) precipitates in P92 (TEM extraction carbon replica).
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strength loss of the aged P92 is probably due to significant precipitation of the Laves phase during ageing and creep exposure, which reduces the amount of W and Mo atoms in solid solution. Thus it can be concluded that the creep behaviour and the creep strength in P91 and P92 steels are controlled by the coexistence of dislocation substructure and precipitates, the latter mainly acting as subgrain stabilisers and effective obstacles to the movement of mobile dislocations through the matrix. Significant microstructure changes in P23 steel were observed after longterm annealing, resulting in the recovery of the bainitic lath structure and a decrease in dislocation density reducing dislocation strengthening (Fig. 5.9). The coarsening of M23C6 carbides, connected with a higher consumption of alloying elements, causes the lowering of solid solution strengthening as well as precipitation strengthening. Similarly, the formation of M6C abates the solid solution strengthening owing to the decline in alloying elements in the matrix. Summarising, the significant creep strength drop of P23 steel after long-term annealing can be explained by the decrease in dislocation and precipitation strengthening, and solid solution strengthening due to instability of the microstructure at high temperature. Modelling of long-term precipitate stability should include predictions of phase stability, nucleation rates, growth rates and coarsening rates for precipitate phases as functions of chemical composition, temperature and stress (Hald, 2008). The CALPHAD type calculations (Hald, 1999) provide useful guidelines for the evaluation of the stable phases in the steel under investigation based only on the chemical composition. These thermodynamic calculations may be fundamental to the predictions of long-term stability of precipitate particles. Series of equilibrium phase diagrams, which show the calculated equilibrium mole fractions of coexisting phases as a function of temperature have been produced, giving useful information for the possibility to accelerate some microstructural changes and thus to simulate long-term service conditions by isothermal ageing. Figure 5.10 shows such a diagram for P91, P92 and P23 steels.
5.4.2 Effect of intermittent heating on creep behaviour As mentioned earlier, critical boiler components are often subjected to complicated load and temperature histories. The closest laboratory simulation involves both temperature and stress cycles. For example, the startup and shutdown cycles can be well simulated by temperature cycling with or without hold times. In order to consider the above effects of nonsteady loading, one needs to design experimental methods for simulating the operating (regime) state, which has many cycles of heating/cooling and nonsteady stressing. Such an approach will be illustrated by our recent
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5.9 TEM micrographs showing the microstructure of P23 steel (a) in the as-received state, (b) after ageing at 650 8C for 10 000 h, (c) after creep at 650 8C and 125 MPa (gauge length) and (d) after ageing at 650 8C for 10 000 h and creep at 600 8C and 125 MPa (gauge length).
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5.10 Calculated equilibrium mole fractions of coexisting phases as a function of temperature for P91, P92 and P23 steels.
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5.11 Scheme of intermittent loading and heating during creep exposure.
5.12 Creep curves of P91 steel with different loading histories at 125 MPa: (a) standard (monotonic creep), (b) cyclic creep curve (uncorrected) and (c) corrected cyclic creep curve. © Woodhead Publishing Limited, 2011
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experimental results on advanced 9–12%Cr ferritic–martensitic steels P91, P92 and E911 (Sklenicˇka et al., 2008). Constant load tensile creep tests were carried out at 600 8C. The effect of temperature variation during the creep test has been investigated by intermittent heating under constant load (Fig. 5.11). The creep specimens in the first cycle were moved from room temperature to the testing temperature, then creep loaded (the holding time at the testing temperature was about 144 h) and finally the specimens were cooled down to room temperature. The same type of temperature cycle was repeated after 24 h without removal of load during the whole creep test up to the final fracture of the specimen. Simultaneously, for comparison purposes, the standard creep tests at the same creep testing temperature were performed. Two elongations ε versus time t creep curves for steel P91 with different loading histories under the same level of the applied uniaxial stress σ of 125 MPa are shown in Fig. 5.12. The first curve (a) represents the standard creep test, which was run at constant load and temperature to the final fracture of the crept specimen. The second curve (b) in Fig. 5.12 shows the creep behaviour of the specimen with intermittent heating, resulting in the fracture of the specimen after 67 cycles. A clearer comparison of these two creep curves follows from the second part of Fig. 5.12, where the creep curve for non-monotonic (cyclic) creep encompasses the time at the testing temperature only. The results of cycled and standard creep tests of stress for P92 and E911 are shown in Fig. 5.13 for the same testing temperature and an applied stress of 150 MPa. As shown in these figures, the cycled specimens failed after 77 cycles (steel P92) and 107 cycles (steel E911), respectively. As demonstrated by the figures, no detrimental effect of intermittent heating on the creep behaviour of the specimen during the cycled tests was found. The cycled specimens exhibit a little longer time to fracture than those of the specimens crept at constant temperature, as can be seen from Table 5.1. A possible explanation for this difference may lie in a serial addition of holding and cooling period for cycled specimens overestimates the time to fracture. In fact, the time to fracture for the specimen of P91 steel subjected to standard loading at 125 MPa is about 9100 h whereas the test piece endurance for a cycled specimen is ~ 9700 h when derived from holding periods only. Thus, it is clear from these experiments that the creep resistance of the investigated steels does not deteriorate by intermittent heating. Microstructural investigation revealed that there are no substantial differences regarding the types and chemical compositions of precipitating phases for specimens after monotonic or cyclic creep in all three steels examined (Sklenicˇka et al., 2008). Subgrain coarsening during creep exposures seems to be more intensive in the specimens crept under a
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Table 5.1 Time to fracture for standard creep tests, tf , cyclic creep, tfc , and corrected time for cyclic tests, tfcor, at temperature 600 8C and various applied stresses Steel
σ (MPa)
tf (h)
tfc (h)
tfcor (h)
P91 P92 E911
125 150 150
9111 8773 13647
11350 13088 18120
9742 11420 15552
5.13 Creep curves at 150 MPa with different loading histories for P92 and E911 steels: (a) standard (monotonic creep) and (b) cyclic curves (uncorrected).
constant testing temperature in comparison with the intermittently heated specimens, as indicated by Figs 5.14, 5.15 and 5.16. However, to complete the characterisation of the effect of interruptions on subgrain size, quantitative data on the size distribution of subgrains (Orlova´ et al., 1998) are strongly needed. Such a quantitative study is in progress at
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5.14 TEM micrographs of P91 steel subjected to monotonic creep at 125 MPa: (a) head location, (b) gauge length location and cyclic creep at 125 MPa, (c) head location, (d) gauge length location.
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5.15 TEM micrographs of P92 steel subjected to creep at 150 MPa: (a) monotonic creep, (b) cyclic creep.
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5.16 TEM micrographs of E911 steel subjected to creep at 150 MPa: (a) monotonic creep, (b) cyclic creep.
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present. The results here suggest that the observed difference in the subgrain microstructure may be explained by possible additional precipitation reactions during intermittent heating. New precipitate particles on subgrain boundaries and dislocations may block the migration of subgrain boundaries or the movement of mobile dislocations, which contributes to the microstructural stability and consequently to the creep strength of the investigated steels. On examining the creep fracture surfaces of the investigated steels it was found that the failure is always of the ductile transgranular dimple mode regardless of the creep loading history. Finally, the temperature changes consisting of cooling the specimen from the testing temperature of 600 8C down to room temperature and heating up again to the testing temperature may, in a real high temperature component, cause some additional strain generated by the thermal stresses created during the temperature changes. However, such a phenomenon seems to have no effect on the creep laboratory behaviour of the specimen. Provided that the microstructures of specimens subjected to both monotonic and cyclic loading are relatively similar there is no reason to expect substantially different creep behaviour for monotonic or cycled creep specimens.
5.5
Predictive modelling
Since it is not possible to measure the creep life directly under the application conditions due to the long times needed, various methods are used to predict creep behaviour from accelerated tests. The higher stress, higher temperature or enhanced stress concentrations are used to accelerate the creep process.
5.5.1 Creep processes If material is loaded at elevated temperatures, many mutually interconnected processes are running in parallel. There are: (i) deformation processes, that is dislocations movement, grain boundary sliding and diffusional transport of matter; (ii) development of the microstructure, that is changes in the dislocation network, precipitation or dissolution of secondary phases, changes in their size and spatial distributions, and changes in local concentrations of elements in solid solution; and (iii) development of creep damage, that is creep cavities and microcracks. It is very difficult to separate individual effects of these various processes, so the complete physical description based on fundamental laws has not been available until now. All interactions in the rather complicated microstructure of the structural metallic materials are not explored enough to formulate exactly all the laws for the microstructure development kinetics. Nevertheless, the physically based description of the creep process
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is useful when looking for the weaknesses of the current phenomenological methods of residual creep life assessment.
5.5.2 Physical description From the physical point of view, the creep process can be formally described by a set of differential equations: e_ ¼ e_ s; T; si ; dj ½5:6 s_i ¼ s_i ðs; T; sk ; dl Þ
½5:7
d_j ¼ d_j ðs; T; sk ; dl Þ
½5:8
where the creep strain rate e_ is the function of the applied stress σ, temperature T, structure of the material expressed by the parameters si and the accumulated damage expressed by the parameters dj. Equations [5.7] and [5.8] represent the kinetics of the microstructure and the damage development. In principle, integration of the set of equations should provide a complete prediction of creep behaviour of the material. It is worth noting that the result may be strongly dependent on the initial state of the material si(0), parameters dj(0) and on the integration path (that is loading history), which is generally complicated under industrial service conditions. Unfortunately, there is still no clear idea about all the parameters that are necessary for the description of the creep process, their mutual influences and their kinetics. Moreover, most si and dj parameters should have a form of distribution functions rather than simple variables. Though a lot of work has been done in this direction (Blum, 2008), only rough approximations for materials of simple microstructure are available. Simplified physical descriptions along with phenomenological and empirical approaches are then used for predictive modelling of the creep process. Some models try to predict the whole creep curve, while others deal with some parameters only, of which the time to fracture tf is the most important one.
5.5.3 Simplified physically based models The models treat the microstructural and damage development separately and incorporate them into the creep deformation process using some special variables. Such a model was proposed by Dyson and McLean (1998) and the most recent, tailored for the boiler structural steels, by Magnusson and
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Sandstro¨m (2007). The latter model uses the modified Norton equation QC e_ ¼ AN exp ð s sb Þ n ½5:9 RT where AN is a constant, n the Norton exponent, QC the activation energy of creep, R the gas constant and σb = σb(t, T, si, dj) is the back stress, expressing the strengthening and softening effects of the microstructure and the damage development. Thermodynamic calculations are used to evaluate the development of microstructural elements. This approach should be considered the most prospective due to its physical basis. Nevertheless, the results obtained by extrapolation should be taken with care in many cases.
5.5.4 Phenomenological models based on creep curve selfsimilarity In the vast majority of cases, the creep curves can be described as a sigmoidal curve of three stages. Moreover, with boiler steels the secondary stage is reduced to a narrow minimum of the creep rate, so the curves consist of overlapping primary (decelerating) and tertiary (accelerating) stages. The models are based on fitting some function to measured creep curves, deriving the dependencies of their parameters on applied stress and temperature, and extrapolating these dependencies. The Θ-projection model (Evans et al., 1992) and logistic creep strain prediction (LCSP) (Holmstro¨m et al., 2007) can be used as an example. In Θ-projection, exponential functions are used to describe both primary and tertiary stages e ¼ y1 ð1 expðy2 tÞÞ þ y3 ðexpðy4 tÞ 1Þ
½5:10
and bilinear dependence on temperature and stress is assumed for all θi parameters yi ¼ Wi1 þ Wi2 T þ Wi3 s þ Wi4 Ts
½5:11
Thus, the sixteen Wij parameters provide a ‘creep curves map’ for particular material under virtually any conditions. The LCSP method uses a different equation for the creep curve description: ( 1=p ) logðtf Þ þ C e ¼ exp 1 x0 ½5:12 logðtÞ þ C where C, p and x0 are parameters and tf is the time to fracture. Note that x0
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must be negative to obtain the normal shape of the creep curve. Parameter dependence on the conditions is assumed in the form p ¼ p1 þ p2 logðsÞ þ
p3 T
x0 ¼ x1 þ x2 logðsÞ þ
x3 T
½5:13 ½5:14
and C is a constant. A table of pi, xi and C values for selected ferritic steels was presented by Holmstro¨m (2010). Since LCSP uses the time to fracture as a parameter, it can be calculated from the model if all other variables are known. In other cases, it can be combined with a tf predicting model (see the next subsection) for extrapolation (Holmstro¨m and Auerkari, 2009). Models of this type are frequently used as a basis for finite element calculations of a complex stress situation in particular parts of the power plant. It is important to realise that the models are built to describe a creep curve with static loading, that is constant temperature and constant stress or load. The models are not capable of describing transient effects in creep under variable loading, so their use in calculations of any dynamic loading situations is inadequate.
5.5.5 Time–temperature parameters For the boiler structures, the creep strain is not so important and creep life is then fully described by the time to fracture tf. Simple models are then used to extrapolate tf from creep rupture tests without taking creep deformation into account. Most of these models use the concept of time–temperature parameters (TTPs). This approach assumes that the creep processes can be scaled by applied temperature, since most of them are controlled by diffusion, which in turn obeys Arrhenius type kinetics. The Arrhenius relation for the time tf can be written in the form Q tf ¼ B exp ½5:15 RT where B is a pre-exponential factor, Q is activation energy for creep and R is the universal gas constant. Two different TTPs were derived from this relation under different assumptions. Sherby and Dorn (1952) assumed that the stress dependence of tf is concentrated in the B factor, while Q is a material constant. The socalled Sherby–Dorn parameter S is thus defined as SðsÞ ¼ logðsÞ
MQ RT
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where M = log(e) transforms the natural logarithm to a decimal one. In contrast, Larson and Miller (1952) assumed the B factor to be a material constant and Q stress dependent. The Larson–Miller parameter L is defined as L ¼ T½C þ logðtf Þ
½5:17
where the constant C = log(B) and L(σ) = MQ(σ)/R. The S or more frequently L parameters used to be plotted against log(σ) to derive the maximal applied stress for a given temperature and creep life. The most recent model of this kind was recently presented by Wilshire et al. (2009). In the model, the Arrhenius type kinetics is modified by a powerlaw creep equation and the applied stress is normalised by the ultimate tensile strength at a given temperature σT, the value of which can be easily measured. The main equation of the model is then s Qc u ¼ exp k1 tf exp ½5:18 sT RT where Qc is assumed to be 300 kJ/mol for ferritic boiler steels and k1 and u are parameters, obtainable from the ln[-ln(σ/σT)] versus ln(tf)-Qc(RT) plot. A common disadvantage of the TTP parameters method is a very strong assumption on the kinetics of all involved processes, which need not be fulfilled in all cases. The steel TAF650 can be used as an example of TTP method failure. The steel exhibits very good creep resistance with short-term tests at higher temperatures, while with long-term tests the creep performance is much worse than that predicted. The creep properties are degraded by precipitation of the Z-phase at lower temperatures. TTP models assume that the precipitation is accelerated by higher temperatures along with other processes, but in fact the precipitation of the Z-phase does not occur at all, since the Z-phase is no longer the equilibrium phase.
5.5.6 Creep models summary Simple phenomenological models are good tools for engineering practice, but their limitations and reliability must be realised, mainly when applied on new materials. The physically based models are contemporary rough estimations only, but only they are capable of dealing with dynamic loading conditions and unexpected changes in microstructure and controlling mechanisms. Thus, the application of different models is very likely to achieve more precise and reliable predictions.
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5.6
Sources of further information and advice
The European COST programmes, which have been running continuously since 1983, have provided a good forum for the development and evaluation of creep-resistant materials. Boiler projects have been mostly focused on tubes, pipes and water walls. Most of the achievements in COST programmes have been published in the series of the COST conferences on Materials for Advanced Power Engineering held in Liege, Belgium, every four years. In addition to the COST actions there are a number of parallel European activities, which also make an important contribution to the qualification of new high temperature materials (e.g. Komet 650, the PIPPE project, the VGB 158 and the Thermie AD 700). The ECCC (the European Creep Collaborative Committee) itself generated creep data packages for the analysis of creep properties of base materials and weldments from the COST and other programmes. In the USA, EPRI has been holding international conferences on Advanced Material Technology for Fossil Power Plant every three years. Similar efforts have been running in Japan (Masuyama, 2006). The general specifications for creep-resistant steels in Japan established JIS (Japanese Industrial Standards) and in METI (Ministry of Economy, Trade and Industry) Codes and JSME (Japan Society of Mechanical Engineers) Codes for power applications (Masuyama, 2008).
5.7
References
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Woodford D A (2004), ‘Creep strength evaluation of serviced and rejuvenated T91 using the stress relaxation method’, in Advances in Materials Technology for Fossil Power Plants, ASM, Chicago, Illinois, pp. 1101–1115. Yaguchi M, Ogata T and Saskai T (2009), ‘Creep strength of high chromium steels welded parts under multiaxial stress conditions’, in Shibli I A and Holsdsworth S R (eds), Creep and Fracture in High Temperature Components–Design and Life Assessment Issues, DEStech Publication, Inc., Zurich, pp. 215–226. Yoshizawa M, Igarashi M, Moriguchi K, Useda A, Armaki H G and Maruyama K (2009), ‘Effect of precipitates on long-term creep deformation properties of P92 and P122 type advanced ferritic steels for USC power plants’, Mater. Sci. Engng A, 510–511, 162–168.
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6 Microstructural degradation in boiler steels: materials developments, properties and assessment J . D O B R Z A N´ S K I , Institute for Ferrous Metallurgy, Poland and A . H E R N A S a n d G . M O S K A L , Silesian University of Technology, Poland
Abstract: This chapter presents a brief description of types, properties and application of ferritic–bainitic (pearlitic), martensitic and austenitic steels designed for use in pressure elements of power boilers. It also describes assessment methodologies for material condition and residual life determination of the elements operated under creep conditions, as well as microstructural degradation processes and properties of creepresistant steels under the influence of long-term exposure to high temperature, with specific reference to the precipitation sequences of carbide phases in ferritic–bainitic and martensitic steels including the new types, e.g. VM12, and in austenitic steels. Key words: power plant steel, residual durability, life-time assessments, degradation of microstructure, modeling of precipitation process.
6.1
Introduction
Creep-resistant steel is integral for the engineering and construction of power plants and is investigated in three areas of scientific research: . . .
design of the chemical composition and optimization of specific characteristics; steel selection and forging of the structural elements; operations, in terms of degradation, material integrity, and long-term durability.
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threshold operating temperatures and pressures. The following properties are required of materials under such conditions, i.e. under conditions that promote creep: .
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suitable mechanical properties at room temperature (plasticity, tensile strength, notch impact strength, elongation) as well as at high temperatures (R1,0 creep limit, RZ/100000 or RZ/200000 creep strength, thermal/mechanical fatigue strength with satisfactory plasticity and crack resistance); excellent microstructural stability and properties stemming from the desire for a low probability of failure and to an increase in strength during long-term operation with the concurrent exposure to temperature and strain; good physical and chemical properties, such as a high coefficient of the heat capacity with minimal linear thermal expansion, resistance to oxidation in the presence of water vapor, and resistance to corrosion as a result of exposure to fumes; good mechanical properties related primarily to bending, welding, and heat treatment processes; an additional criterion for the selection of steel being the price of the resulting products.
Although the operational temperature of the steam is not the only factor that contributes to degradation and instability in structural elements, it is the most important factor. For years, this temperature has not exceeded 540 8C under pressures of up to ca. 20 MPa in conventional boilers. The steady trend towards increased efficiency, reduced emissions of harmful pollutants into the atmosphere (mainly CO2, NOx, and SO2), and fuel conservation has led to significant increases in the operational steam temperatures (565–620 8C) and pressures (24–30 MPa). This increase is noted in the operation of new boilers already in use, as well as in boilers that have been built to withstand supercritical conditions (which are 40 % more efficient). The most advanced energy technology processes employ ultrasupercritical boilers in which the temperature of steam reaches 650–700 8C at pressures of 30–35 MPa, providing conversion efficiencies on the order of 45–50 %. These technologies do, however, require the use of materials with high functional specifications, and this requirement simultaneously drives progress in materials engineering.
6.2
The development of steel for power engineering
Steady determined efforts to increase the efficiency of power units by raising steam temperatures with concurrent significant pressure increases has
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motivated the search for engineering solutions that fulfill the ever increasing demands on steel used in boilers’ critical elements: the elements that operate under extreme temperature and stress conditions. The first high steam parameter boilers, built in the 1960s in the USA, Japan, and Ukraine, among other countries, used 18Cr–8Ni, 18Cr–10Ni(Ti), and 18Cr–13Ni (Mo) austenitic steels. The ensuing technological and maintenance problems halted further development of these materials. Power engineering around the world continued to use conventional solutions with maximum steam temperatures of 540 8C and pressures on the order of 18.6 MPa, which provided power unit efficiencies of up to 33 % (‘Report 2030’, 2008; Bernasin´ski, 2008; Najgebauer and Patrycy, 2008). Mild alloy steels processed in the 1930s and 1940s successfully met the material requirements, mainly short-term creep rupture strength. Among these alloy steels were the easily weldable ferritic–pearlitic steels 15Mo3 and 13CrMo4-5, the ferritic– bainitic steel 10CrMo9-10, as well as the less easily weldable bainitic–ferritic steel 14MoV6-3 and the high-alloy martensitic steel X20CrMoV12.1. These steels were used for coils, superheater chambers, and temperature regulators, as well as for pipework that operated under creep conditions. Since the 1970s, testing has focused on the development of steel processing methods that yield steels with properties superior to X20CrMoV12.1 steel. Testing programs have included the European COST 501 and 502, American EPRI-1403, and Japanese R&D along with the subsequent COST 522 and 536 programs, which resulted in the development of new bainitic and martensitic steels with properties that allow for use under conditions in which temperatures reach 620–650 8C, and at pressures of 30 MPa. The chemical composition modifications achieved provide greater microstructural and property stability during long-term use under creep conditions. The relatively good technology-oriented properties have been achieved by lowering the carbon content, by introducing tungsten and cobalt, and by introducing microadditions of boron, nitrogen, and niobium. The novel materials currently used to construct supercritical boilers may be categorized as follows: .
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Mild-alloy martensitic steels based on 2.25Cr–Mo steel, with chemical compositions that have been improved via the introduction of V, Ti, Nb, B, W (T/P23, T/P24). Membrane walls composed of these steels withstand operating temperatures up to 550 8C. High-alloy martensitic steels based on 9%CrMoV steel, with chemical compositions that have been improved via the introduction of W, Nb, N, B, Cu (T/P91, T/P92, E911, PB2). These steels withstand temperatures up to 600 8C and are mainly used in superheater chambers, steam pipes, and superheater coils.
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C 0.04–0.10 0.05–0.10 0.08–0.12 0.09–0.13 0.07–0.13 0.07–0.14 0.08–0.18
Cr 1.90–2.60 2.20–2.60 8–9.5 8.5–9.5 8.5–9.5 10–12.5 10–13
Mo 0.05–0.30 0.90–1.10 0.85–1.05 0.9–1.1 0.3–0.6 0.25–0.6 max. 0.8
V 0.20–0.30 0.20–0.30 0.18–0.25 0.18–0.25 0.15–0.25 0.15–0.3 0.18–0.3
Nb 0.02–0.08 — 0.06–0.1 0.06–0.1 0.04–0.09 0.04–0.1 0.03–0.06
W 1.45–1.75 — — 0.9–1.1 1.5–2 1.5–2.5 1–1.8
N max. 0.030 max. 0.010 0.03–0.07 0.05–0.09 0.03–0.07 0.04–0.1 0.03–0.09
Chemical composition
Sources: Husemann (2005); Viswanathan et al., (2005b); Masuyama (1998, 2004).
T/P23 7CrMoVTiB10-10T/P24 X10CrMoVNb9-lT/P91 XI lCrMoWV Nb9-1-1 E911 X10CrWMoNb9-2 T/P 92 HCM12A CT/P 122 XI2CrCoWVNb 12-2-2 VM12
Grade
B max. 0.006 max. 0.007 — 0.001–0.006 0.01–0.06 max. 0.05 0.001–0.01
Other Ti 0.05–0.10 — — — — Cu up to 1.7 Co 0.5–2
Table 6.1 Selected bainitic and martensitic steels intended for thick wall components and coils, where the T-tube is intended for thinwalled tubes and coils and P are pipes (Si=0.6 max, Mn=0.8 max, Ni = 0.6 max)
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6.1 Comparison of creep strength of bainitic steels with martensitic steel 9Cr–Mo–VNNb type (Prager, 2006).
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High-alloy steels based on 12 %CrMoV steels, with chemical compositions that have been improved via the introduction of W, Co, Nb, N, B, Cu (HCM12(A), EM12M, VM12, MARB2 (developed by NIMS)). These steels withstand temperatures up to 620 8C, and are mainly used for superheater coils and steam pipes. Steels and nickel superalloys with an austenitic matrix and an improved chemical composition and microstructure. These steels are used in the temperature range 560–700 8C, mainly for superheater elements, pipes, and turbine regulator valves.
Table 6.1 shows the chemical compositions of selected bainitic and martensitic steels (Husemann, 2005; Viswanathan et al., 2005b; Masuyama, 1998, 2004), while Fig. 6.1 compares their creep rupture strengths (Prager, 2006). Among the austenitic steels, standard 18-8, 18-10-Ti, and 16-13-Mo steels (TP304H, TP321H, and TP316H, respectively) have been used for superheater coils. The construction of supercritical boilers presently uses the austenitic steels TP347HFG, SUPER304H, HR3C, and Sanicro 25, in which precipitation strengthening and age-hardening are achieved through additions of Nb and N, which permit operation at steam temperatures of 560–620 8C. The chemical compositions of the austenitic steels are shown in Table 6.2 (Husemann, 2005; Viswanathan, et al,. 2005a, 2005b; Iseda et al.,
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Cu — — 2.50—3.50 — 2.99 —
Nb — 86C 0.30–0.60 0.20–0.60 0.44 0.2
Ti 0.5 — — — — 0.1
N — — 0.05–0.12 0.15–0.35 0.24 —
Sources: Husemann (2005); Viswanathan et al. (2005a, 2005b); Iseda et al. (2007); Blum and Vanstone (2006).
Mo — — — — — 1.5
EN 12952 620 8C 620 8C 645 8C 670 8C 700 8C —
Ni 10.0 9.0–13.0 7.5–10.0 17.0–23.0 25.5 25
C 0.08 0.04–0.10 0.07–0.13 0.04–0.10 max. 0.08 0.15
TP321H TP347HFG Super 304 H HR3C Sanicro25 NF709
Cr 18.0 17.0–20.0 17.0–19.0 24.0–26.0 22.6 25
Calculated temperature as per:
Chemical composition
Grade
Table 6.2 Nominal chemical compositions of the most popular austenitic steels used in the construction of power boilers (Si=0.6 max, Mn=0.8 max, Ni=0.6 max)
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6.2 Allowable stress at creep strength Rz/t for 100 000 h of martensitic and austenitic steels (Benedick et al., 1998).
6.3 Creep rupture strength (for 100 000 h) for martensitic, austenitic steel, and nickel superalloys (Viswanathan et al., 2005a).
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2007), and their creep rupture strengths are given in Fig. 6.2 (Benedick et al., 1998). The short-term operational outlook for current USC boiler projects indicates that only nickel superalloys ensure that the requisite high functionality of the pressure components’ critical elements may be met (Fig. 6.3) (Viswanathan et al., 2005a). Recent tests conducted in various countries show that the material strength of the reference material alloys HR6W, CR30A, Inconel 617, the refined Japanese alloy CCA617, Inconel 625, and Inconel 740 was conferred by the intermetallic phase γ´. Technological developments in chemical composition have sought to improve the creep-resistant properties relative to those of reference steels (Smith and Sizek, 2000; Husemann, 2005; Shingledecker et al., 2005; Viswanathan et al. 2005a; Blum and Vanstone, 2006; ‘INCONEL alloy 625’; Rautio and Bruce). Following creep rupture strength, the second most important factor for determining the suitability of a material for high-temperature use in the critical elements of supercritical boilers is the material’s resistance to hightemperature corrosion, oxidation in the presence of water vapor, and resistance to spallation of the oxide coating. Oxidation resistance mainly hinges on the concentration of chromium in the steel (Fig. 6.4) (Hahn et al., 2009); the composition of the coating changes from Fe2O3/Fe3O4 through Fe3O4/(Fe,Cr)3O4 /Cr2O3 to a layer of the pure oxide Cr2O3. Tests indicate (Husemann, 2005; Viswanathan et al., 2005a) that at
6.4 Oxidation resistance comparision of (9–12)%Cr steels and austenitic TP347 steel in pure steam (Hahn et al., 2009).
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temperatures of 620–650 8C, the greatest propensity for spallation is exhibited by ferritic/martensitic steels T23, T91, and T92, whereas austenitic steels Super 304H and TP347H have a lower propensity for spallation, which can be improved further via grain refinement (FG) and shot peening (blasting), which definitively increases the adhesion strength of the Cr2O3 oxide layer (Grabke et al., 1998). The greatest oxidation resistance with minimal spallation is provided by the nickel superalloys Nimonic 263, Inconel 617, 625, CCA617, and Inconel 740.
6.3
Methodology for assessing the state of a material and determining the residual durability of the operational elements under creep conditions
6.3.1 Assessment of pressure component elements made from a ferritic-matrix steel Measurements of material properties during creep tests and long-term operation are used to determine the durability of power plant elements. Creep tests measure a material’s durability in terms of the time to rupture in a creep test, or the time limit for safe operation prior to failure. For practical purposes an objective failure criterion must be developed. The projected lifetime of an element is determined based on the material durability measured in a creep test, including consideration for the oxidation resistance, although it is not identical to it. Elements are required to maintain durability between 100 000 and 200 000 hours of use; that is (e.g Polish Standards PN-75/H-84024, 1975) sop kmax ¼
Rz=100000 =T X
½6:1
where X is the safety coefficient. The following factors impact the durability of pressure devices: . . . .
material factors and factors involved in the production of the element (chemical composition, microstructure, properties, structural and technological factors – production quality); installation-related factors; usage-related factors (e.g. temperature and pressure variability, number of boiler launches and shutdowns, aggressiveness of the environment); maintenance, diagnostics, repair, and upgrade factors.
The complex interplay between these effects implies that the severity of the degradation processes outlined in section 6.4 may vary considerably during long-term usage. Monitoring the state of the material and of the component
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permits a prediction to be made of the timescale for extended safe operation. Each structural object and component must be treated individually. To this end, the material assessment procedures and the criteria for assessing the degree of material degradation are subject to constant refinement, which allows a determination of an objective time period for safe operation based on non-destructive testing and measurements, as well as destructive testing. From a practical point of view, a user is less interested in the durability as defined by the time to failure, as by the time for safe operation (available durability), which is limited by the critical degree of degradation or component damage. The degree of degradation f is defined most often as the relationship between the actual operation (or creep) time t to the time of element disintegration under the same operational conditions (or in a creep test) tr, i.e. f = t/tr. However, the degree of degradation (ft = Σfn) is treated as the sum of all degradation processes under all operational conditions (temperature and pressure). As a result, the residual life RCL is defined by the following relationship (based on Viswanathan, 1994, and Regis and D’Angelo, 1989): RCL ¼ ð1 ft ÞtSR
½6:2
where tSR is the time to disintegration in conditions of prolonged operation. In practice, residual life generally composes 0.6–0.8 of the total life, where the numeral 1.0 represents the state of failure. The general methodology for assessing the material and element state is shown in Fig. 6.5 (Hernas and Dobrzan´ski, 2003). Research in this area requires the synthesis of both theoretical and practical knowledge, as well as engineering experimentation. Thus, to increase the confidence in the material (or component) assessment, an array of testing methods is used, depending on the type of element and the nature of its function. The regions of a component that experience the most strain must be selected for testing. The life of a structural component is assumed to be determined by its weakest point or by the most degraded region. Therefore, the result of a material’s condition is not simply determined by the average value of the results obtained from the criteria used to assess the degree of material degradation. The methods used for assessing residual life may be divided into two groups (Cane, 1986; Dobrzan´ski and Milin´ski; 1992; Kautz, 1997): . .
methods that rely on non-destructive or destructive (after-operation) testing; computational methods that collect the controlled operational parameters, standard material data, durability measurements (accumulation of creep and fatigue depletion), and numerical assistance.
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6.5 Ways to approach the evaluation of the state of a material or component (Hernas and Dobrzan´ski, 2003).
To increase the objectivity and reliability of residual life assessment, engineering employs methods from both groups. Non-destructive testing involves measuring geometric features (defectoscopic, endoscopic, and metallographic tests) using response techniques and hardness measurements. Non-destructive physical methods are also used to measure magnetic properties. Metallographic tests are conducted to assess objectively a material’s microstructural state (Si) and the state of internal defects (ω): e_ ¼ fðs; T; Si ; oÞ
½6:3
Discussion is ongoing assuming that the creep rate, e_ , defined by equation (6.3), is a function of pressure σ, temperature T, microstructural state Si, and the state of the internal damage ω resulting from the nucleation and development of pores (Cane, 1986; Dobrzan´ski and Milin´ski, 1992), where Si = S1, S2, S3 . . . , Sn defines the microstructural properties of, for example, the grain size, subgrain size, type of precipitate, and their role, morphology, and location.
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6.6 Four classes of internal defects during long-term operation (Neubauer and Wedel, 1984).
Experiments in analyzing the degradation processes of pressure components, working under creep conditions, have allowed the general acceptance of four classes of internal defects, which are assigned the degree of degradation f (Fig. 6.6). These classes outline the appropriate behavior connected with diagnostics and the time of expected further operation (Neubauer and Wedel, 1984). Apart from microstructure tests, tests of the fundamental mechanical and LCF properties, and non-destructive tests of the material creep, before and after use, provide more trustworthy characterizations of the predicted life and residual life. Because the test procedures are time intensive, accelerated creep tests are conducted with a time to rupture of at most 10 000 hours, at operational pressures, and at test temperatures that are elevated relative to operational temperatures. The results obtained are extrapolated to operational conditions. Computational methods for assessing residual life are widely welcomed by power plant users. Computational models are used to assess chambers, pipes, turbine rotors, and bodies in accordance with the requirements for national or European standards, e.g. EN-12952-4/2000, TRD-508, or ASME Code Case N-47. They predict the probable effect of the operation process on the degree of material degradation of operational elements for timescales in excess of the computed time of 100 000 hours. Using measurements of actual operational parameters (temperatures, pressures → operating stresses resulting from the component’s geometric features) as
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well as trustworthy estimates for the creep rupture strength (for 100 000 and 200 000 hours), one may determine: . .
the time limit for operation of the element, and compare it with the current time of operation; the predicted component wall thickness at 200 000 hours, and compare this value with the measured thickness of the component wall.
The computational results obtained and the results of the material state assessment permit operators to predict: . . .
the time limit for prolonged safe operation under current or deteriorating conditions; the time limits for conducting complex diagnostic tests; component replacement times and conditions.
The computational results are, however, vulnerable to errors stemming from the data that characterize the material in its original state, although not from data that characterize the material in a specific state after operation. The residual life assessment methodology described has been validated in Poland at all power plants with operational times between 100 000 and 300 000 hours. A team of specialists from the Institute for Ferrous Metallurgy in Gliwice, the Silesian University of Technology, and the RAFAKO SA Boiler Factory has designed models for the microstructural changes that occur in the conventional mild-alloy steels and in X20CrMoV12.1, and they have devised a means of classifying and assessing the state of the materials. The material assessment relies on assigning the actual microstructure degradation class using partial assessments of the changes in areas of pearlite/bainite/martensite, advancing precipitation processes, internal defects, and the appropriate degree of material degradation. Knowing the degree of material degradation allows an estimation to be made of a safe timeframe for further operation of the element being tested, under the required operational parameters. The procedure relies on the correlations between the state of microstructural degradation, internal defects, and the creep test results before and after operation. An example assessment method for the material state of an element composed of the steel 10CrMo910 exposed to creep conditions is shown in Fig. 6.7 (Dobrzan´ski, 2002). The material reaches a critical state when the internal defects arise. In particular: .
Degradation degrees of f = 0.2–0.6 are ascribed to a material without creep defects, but in which the microstructure has decomposed in pearlite/bainite/martensite areas and in which precipitation processes have developed on the grain boundaries and within grains.
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6.7 Assessment method for the material state of an element composed of the steel 10CrMo910 and exposed to creep conditions (Dobrzan´ski, 2002).
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6.8 A scheme of microstructure change of 12Cr–1Mo–V steel during long-term services without internal damage (Dobrzan´ski, 2005).
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Degradation values of f = 0.6–0.8 are ascribed to elements containing isolated oriented pores (classes A and B in Fig. 6.6), as well as displaying advanced spheroidization of carbides within grains or as chains on the grain boundaries. Degradation values of f = 0.8–1.0 indicate an advanced state of defect in the form of macro- and microfissures (classes C and D in Fig. 6.6), as well as the presence of precipitates that form continuous networks on the boundaries and deposits of coagulated precipitates that are evenly distributed within the grains.
It should be emphasized that the method for assessing the degree of material degradation described here is used to determine the safe operation
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6.9 Microstructural patterns of changes in the microstructure image of 12Cr–1Mo–V steel after operation under creep condition: (a) class of microstructure – 0; (b) class of microstructure – 1; (c) class of microstructure – 2; (d) class of microstructure – 3 (continues on next page).
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6.9
(continued)
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6.10 A scheme of microstructure change of 0.5Cr–0.5 Mo–0, +925V steel during long-term services without internal damage (Dobrzan´ski, 2003).
lifetime of the critical elements used in pressure and turbine areas of the boiler, which operate under the most extreme temperature and pressure conditions, i.e. the superheater elements (chambers, coils), temperature regulators, or main and communicating steam pipes. Many years of research carried out by the Institute for Ferrous Metallurgy, Silesian University of Technology, and RAFAKO S.A. have made it possible to prepare the diagrams of structural changes during longterm operation, under creep conditions, in connection with exhaustion degree in terms of strain resulting from creep. The diagrams include materials after operation without internal defects. The diagrams of structural changes generated include 12Cr–1Mo–V with a tempered martensite structure and 0.5Cr–0.5Mo–0.25V with a mixed bainite
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6.11 Microstructural patterns of changes in the microstructure image of 0.5Cr–0.5Mo–0.25V steel after operation under creep conditions: (a) class 0 (160–180 HV10); (b) class 1 (150–160 HV10); (c) class 2 (140– 150 HV10); (d) class 3 (120–140 HV10).
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6.12 Classification of creep damage in 12Cr–1Mo–V steel depending on the degree of exhaustion (Dobrzan´ski, 2005).
and ferrite structure (Figs 6.8 and 6.10). Micrographs showing successive stages of changes in the microstructure of those steels after operation under creep conditions are shown in Fig. 6.9 for 12Cr–1Mo–V steel and Fig. 6.11 for 0.5Cr–0.5Mo–0.25V steel, respectively. Separate diagrams of microstructural changes were prepared for the steels tested after operation under creep conditions and having internal defects. The diagrams of development of internal defects depending on the exhaustion degree are also shown in Fig. 6.12 for 12Cr–1Mo–V steel and Fig. 6.14 for 0.5Cr–0.5Mo–0.25V steel. The selected patterns of structures including successive development stages of internal defects in those steels after long-term operation under creep conditions are shown in further figures presenting images of structures including successive development stages of internal defects in those steels (Fig. 6.13 for 12Cr–1Mo–V steel and Fig. 6.15 for 0.5Cr–0.5Mo–0.25V steel).
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6.13 Microstructural patterns including internal defects for 12Cr–1Mo– V steel after operation under creep conditions: (a) void nucleation at the interphase boundary (ferrite – inclusion) – class 4-A/1; (b) development of chains of voids along the boundaries of former austenite grains – class 5a/B2; (c) transgranular cracks across one grain – class 5b-B/4; (d) transgranular cracks including boundaries of a few former austenite grains – class 6-C/l (continues on next page).
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6.13
(continued)
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6.14 Classification of creep damage in 0.5Cr–0.5Mo–0.25V steel depending on the degree of exhaustion (Dobrzan´ski, 2003).
6.3.2 Assessment of pressure component elements made of austenitic steel In Europe, fewer experimental assessments are made of components made of austenitic steel subjected to long-term operation under creep conditions. The processes and mechanisms of deterioration in such steels differ somewhat from those in steels with a ferritic matrix. However, the general principles governing assessment of residual life are similar. These are: . .
qualitative and quantitative assessment of the microstructural state using light and scanning microscopy techniques, as well as physical– chemical methods; a method for qualitative and quantitative assessment of creep defects;
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6.15 Microstructural patterns including internal defects for 0.5Cr– 0.5Mo–0.25V steel after operation under creep conditions: (a) single voids along boundaries of ferrite grains not uniformly arranged in lowalloy bainitic–ferritic steel 0.5Cr–0.5Mo–0.25V – class 4-A/1; (b) development of voids along boundaries of ferrite grains. Voids oriented perpendicular or at an angle of 458 to the direction of the main stress – class 5a/B1; (c) development of internal defects – coalescence of voids along boundaries of ferrite grains – class 5b/B3; (d) transgranular cracks – class 6-C.
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248 . .
Power plant life management and performance improvement an assessment of the mechanical properties with particular consideration forassessing residual life based on accelerated creep tests; an analysis of the oxidation processes.
The testing methods aim to develop objective criteria for assessing the state of a material by correlating the microstructural state (the type, quantity, and location of the phases that are present), the state of the creep defects, and the residual creep rupture strength as determined by laboratory creep tests of the material after operation. A statistical sample size permits development of appropriate model states and models for use in assessment procedures for the degradation of steel, as is done for ferritic-matrix steel. Every country or energy concern that is engaged in this issue has a preferred approach that aims to increase assessment objectivity and simplify the methodology, with preferences for non-destructive tests. At the same time, the goal is also to also reduce the time and money spent on testing (EPRI, 2002). The test results and the knowledge gained from assessing elements, including those made of ferritic steel, allow the formulation of basic criteria for assessing the degree of degradation in austenitic steels. The criteria for assessing the degree of degradation for austenitic steels are: . . . . .
the state of advanced precipitation processes, in particular, the role, size, and location of the sigma phase, and particularly whether the grain boundaries have been filled with this phase; advanced creep defects; residual life in accelerated creep tests; the thickness of the oxide coating (mainly internal); changes in the magnetic properties that require correlations with the results of other methods.
For example, in the 321H steel, the fundamental degradation criterion will be an assessment of the degree of sigma phase precipitation at grain boundaries and the state of internal defects.
6.4
Characteristics of microstructure and property degradation processes
Creep-resistant materials subjected to the long-term effects of increased temperatures and pressures, as well as to corrosive environments, undergo a process of degradation that primarily involves: . . .
microstructural changes; formation and development of pores and microfissures; decreased functional properties (mechanical and physical–chemical).
Analyses conducted in this period are both comparative and practical in
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nature, seeking to define the links between the microstructural state of the material and the properties that determine the durability and reliability of the structural element under specific operational conditions. From the materials science perspective, degradation processes are viewed as problems of stability (or instability) in the microstructure and in the material’s properties. Microstructural degradation is described as an array of structural and physical–chemical factors that undergo changes under the action of extended periods of increased temperature (and pressures), e.g., during creep tests, isothermal annealing, and most of all during operation. These conditions lead to the creation of pores and microfissures, and eventually to the deterioration of the structural element (Herans, 1992; Scarlin et al., 1994; Vanstone, 1998; Zielin´ska-Lipiec et al.; 1998). The general criteria for the degradation (instability) of microstructures are: . . . .
the degree of substructure changes, including changes in dislocation density and the progression of the recovery and recrystallization processes (subgrain size), both of which are vital; the transformations of carbides and the precipitation of intermetallic phases; for e.g., Laves, Z, σ; changes in phase morphology (in decomposition, shape, size, and the distance between particles); the degree of matrix depletion in alloying elements, mainly in Cr, Mo, W.
Regardless of the chromium concentration in the steel, the above factors affect corrosion resistance, including the adhesion of oxide coatings, decrease in strength and increased susceptibility to crack development. The specific applications for creep-resistant steels within specific temperature, pressure, and environmental conditions determine the stability of these processes as a function of chemical composition, base type, and microstructure. As with ferritic steels, the residual life RCL in austenitic steels may be determined. In austenitic steels, the quantitative fraction of grain boundaries, characterized by the presence of defects in the form of pores, is considered. The formula defining residual durability is 2 RCL ¼t 0; 0:6 1 ½6:4 fa where t is the operation time and fa is the share of the grain boundary surface with defects in the form of pores (Needham, 1985).
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6.4.1 Degradation of microstructure and mild-alloy steel properties The basic types of steel developed more than 50 years ago, 1Cr–0.5Mo, 2.25Cr–1Mo, and (0.5–1.5)Cr–1Mo–0.25V, as well as new steels, such as T23 and T24, are used after normalization and high tempering. Depending on the speed of quenching (wall thickness) they may show different microstructures, i.e.: . . . .
ferritic–pearlitic; ferritic–pearlitic–bainitic; ferritic–bainitic; bainitic or bainitic–martensitic (e.g. T24).
Thorough testing indicates that the original structures impart an array of essential properties that determine long-term operation. For example, 0.5Cr–1Mo steel consists predominantly of bainite, which yields a high initial creep rupture strength. However, this steel quickly loses its strength and plasticity during long-term operation as a result of the reduced matrix stability relative to the initial structural mixture of ferrite and pearlite. In their original state, ferritic creep-resistant steels, both mild-alloy and high-alloy, have an increased dislocation concentration on the whole, which, after extended exposures to creep conditions (or operation), asymptotes to a lower stable value (Zielin´ska-Lipiec, et al., 1998). Within the established creep range, the average dislocation concentration value does not change, but the configuration does change as a result of recovery and the polygonization processes. The measurable indicators of substructure stability are the subgrain size, the stability of which has a positive influence on the stability of creep resistance (Vanstone, 1998). The results of current tests indicate that changes in creep rupture strength in high-alloy martensitic and austenitic steels are a function of carbide transformation and intermetallic phase precipitation (e.g. Laves, Z, σ). Depending on the heat and mechanical loads, the transformation and precipitation occur in parallel, through the dissolution and nucleation of new phases. Changes in morphology of carbides and precipitates occur as a result of processes of coalescence, coagulation, and precipitation on the grain boundaries. The sequence of these changes in mild-alloy steels is as follows:
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ð1Cr 0:5MoÞM3 C þ ðMo2 CÞ ) M23 C6 þ M7 C3 þ M6 C ) M6 C þ M23 C6 ð10CrMo910ÞM3 C þ Mo2 C þ ðM23 C6 Þ ) M6 C þ M23 C6 þ ðM2 CÞ þ ðM7 C3 Þ
½6:5
½6:6
Examples of the structural changes for the initial and subsequent stages of the precipitation processes for 0.5Cr–0.5Mo–0.25V, observed by SEM, are shown in Fig. 6.11. The first stage of structural changes is characterized by a slight decomposition of bainite/pearlite areas. That is accompanied by coagulation of M3C precipitates in those areas. The carbide coagulation process is accompanied by commencement of M23C6 carbide precipitation along ferrite grain boundaries. Simultaneously, within the ferrite grains the growth of very fine MC carbide precipitates occurs along with M3C carbides, as shown in Fig. 6.11 (class 1). The long-term influence of temperature and stress results in the dissolution of one type of carbides while other types of carbides are formed at the same time or in formation of new carbides by in situ transformation. The next stage of structural change is significant decomposition of bainite/ pearlite areas. Within bainite/pearlite areas there are coagulated M3C precipitates of varied size (some of quite large size). The M23C6 carbide precipitates forming chains can be seen along ferrite grain boundaries. Simultaneously, MC carbide precipitates are still observed within the ferrite grains. Examples of structural images characteristic for this stage of changes observed with SEM are also shown in Fig. 6.11 (class 2). The final structural change is ferrite with uniformly arranged MC and M6C precipitates within grains and chains of large mainly M23C6 precipitates along grain boundaries. The main phase component of material precipitation is M6C and M23C6 carbides with little evidence of other types of carbides such as MC, M7C3, and M3C, depending on operation conditions. An example of this final structure observed with SEM is shown in Fig. 6.11 (class 3). The phase composition and morphology of carbides, especially the size of the grain boundary, decidedly reflect the changes in creep rupture strength (Figs 6.16 and 6.17). The appearance of the carbide M6C does not disqualify the material’s usefulness for long-term operation, despite opinions originally articulated in the literature during the 1970s (Herans, 1992; Dobrzan´ski, 1995; Hernas and Dobrzan´ski, 1995; Pasternak et al, 1997).
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6.16 Influence of the phase composition of carbides, temperature, and operating time on creep resistance of 1Cr–0.5Mo type steel (Dobrzan´ski, 1995).
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6.17 Influence of the carbide type in the ferrite on the durability of 1Cr– 0.5Mo steel after long time operations (Dobrzan´ski, 1995).
6.4.2 Degradation of the microstructure and properties of high-alloy martensitic steels of the 9–12 %Cr group After air quenching, these steels have a uniform martensitic structure that, after subsequent high tempering, forms elongated alpha phase blocks with M23C6 or MC carbides on grain and block boundaries. Depending on the concentrations of other additives (Mo, W, V, Nb, N, B), carbide transformations in these steels are similar to the transformations in mildalloy steel but occur at temperatures of 540–650 8C. There is also the tendancy for intermetallic phase precipitation of the Laves and Z type (Hernas and Maciejny, 1998). The upper suitable temperature limit for long-term operation of martensitic steels depends on the occurrence of processes that decrease the creep strength and plasticity, such as those described by Zielinska-Lipiec et al., 2009a:, namely: . . .
breakdown of martensite and the precipitation of intermetallic phases; coarsening and migration of carbides and intermetallic phases to grain and block boundaries; polygonization of the alpha matrix and increase in subgrain size.
The literature contains many descriptions of the degradation processes of X20CrMoV12.1 steel (Battaini et al., 1990; Ennis et al., 1998; Strang and Vodarek, 1998; Rodak et al., 2003; Zheng-Fei and Zhen-Guo, 2004). This martensitic steel was extensively used in power engineering because the composition of this steel was designed for significantly longer than 200 000 hours, but the steel was withdrawn from use. This steel was mainly used in subcritical superheater chambers and coils, and at steam temperatures of up
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to 540 8C. It was characterized by a relatively high microstructural stability and functional properties. Microstructural changes during long-term operation depended not so much on the operational time as on the actual operational temperature, in particular, on how often the optimal temperature was exceeded. As a result, its structure showed an intensification of precipitation processes on the grain boundaries of former austenite and blocks of martensite, which favor the precipitation of the Laves phase (Fe, Cr, Si)2Mo, the development of polygonization processes, and the disappearance of the martensite’s block structure. The initial stages of the precipitation processes are characterized by gradual decomposition of tempered martensite. It is shown by partial disappearance of martensite laths, precipitation in a form of chains along grain boundaries of former austenite, small growth of subgrains, and the number and size of M23C6 carbides. The subsequent changes include further growth of subgrains and the growth of carbide size by their coarsening and coalescence. The last stage of changes occurring due to long-term creep is the ferrite and carbide structure, with further coarsening and coalescence and simultaneous growth of ferrite grains after decomposition of martensite and formation of the Laves phase. Figure 6.9 (class 3) shows structural images characterized by significant and almost total disappearance of tempered martensite, consisting of ferrite and M23C6 carbide. An increase in the steam temperature and pressure in modernized subcritical boilers and in new supercritical boilers has required the use of steels with significantly higher creep rupture strengths and greater microstructural stabilities. These requirements can be met by (9–12 %)Cr– Mo–V steel containing tungsten or cobalt additives, and with microadditions of Nb, N, and B, which create stable microdispersive phases, such as MX. However, creep test results to date for both P92 and VM12 steels have not been completely satisfactory, and have narrowed the temperature range in which they may be used for long-term operation to guard against precipitation processes. In the 9%CrMo(W)V steel group, the precipitation processes occur as follows: 9%CrMoðWÞVð550590o CÞM23 C6 þ MðC; XÞ ) MC þ M23 C6 þ Laves phase
½6:7
Over the temperature range from 600 8C to 650 8C, the VM12 (X12CrCoWVNb12-2-2) steel shows an initially rapid increase in M23C6 and MX carbide particles, followed by a gradual stabilization of their sizes. Over extended exposure times or high creep temperatures, dissolution of the MX phase and precipitation of the ZCr(Nb,V)N or Laves (Fe, Cr)2(Mo, W) phases ensues on the grain boundaries with M23C6 carbides (Zielinska-
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6.18 Example of the VM12 microstructure in the initial state and after service with a strong effect of degradation (Zielin´ska–Lipiec et al., 2009c); (a) microstructure of VM12 steel in the delivery conditions – elongated subgrains with precipitated M23C6 carbides; (b) carbonitride type M(C,N) within the subgrains of tempered martensite; (c) microstructure of VM12 steel after creeping at 625 8C for 17 478 hours – subgrain structure; (d) precipitates of M(C,N) within subgrains (continues on next page).
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6.18
(continued)
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6.19 Changes in the dislocation density in P92 steel after creep (Zielin´ska-Lipiec et al., 1998).
6.20 Influence of the carbide type in the ferrite on the durability of martensitic steel after long time operation (Dobrzan´ski, 1995).
Lipiec et al., 2004, 2009b, 2009d). The precipitation process is as follows: VM12ð600650o CÞM23 C6 þ MX ) M23 C6 þ MX þ M6 C þ Z þ Laves phase
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Small domains of delta-ferrite that occur in the (9–12) %Cr steel structure are also conducive to the transformations that occur during creep. A large decrease in strength during long-term creep is caused by the precipitation of large grain boundary carbides up to about 1 μm in size (Fig. 6.18). A general diagram capturing the effects of the strength mechanisms on the decreased creep rupture strength of steel with tungsten is shown in Fig. 6.19. Of course it is necessary to note that these data concern the definitely shorter time exposure. In consequence, it is difficult to use them to verify long-term conditions. For these reasons, VM12 steel did not possess the anticipated high creep rupture strength, and its use was significantly limited (Danielsen and Hald, 2006; Hald, 2008, 2009; Cipolla, 2009). Schematically these phenomena are presented on Fig. 6.20.
6.4.3 Degradation of the microstructure and austenitic steel properties during long-term operation After solution heat treatment but prior to operational use, these steels have a homogeneous austenitic structure with annealing twins, with grain sizes on the order of ASTM 4-6 and 8-9 for fine grains. In steels stabilized by Ti and Nb after solution heat treatment, some isolated primary MC carbides are present in typical, regular shapes (Fig. 6.21). Among the austenitic steels used in power engineering, many reports have described the degradation processes of TP321H and 316 steels, from which superheater and reheater coils are made (for example, in the USA they have functioned in excess of 200 000 hours) (Swindeman et al., 1986, Swindeman, 1998; Soo Woo, 2002; Bouchard et al., 2004). Long-term operation intensifies the dissolution/precipitation of secondary phases, namely carbides and intermetallic phases, typically sigma: Austeniteðþferrite; MCprimaryÞ ) M23 C6 þ MC þ sphase
½6:9
The occurrence of these phases and their distribution (with the exception of the TiC carbide) determines the degree of material degradation. In steels containing molybdenum, e.g. 316, the Laves phase Fe2Mo and η phase arise after long operation times under creep conditions at temperatures in the range 600–700 8C. Microstructure tests using light and scanning microscopies of coil samples from the Amos 3 power plant in the USA after more than 200 000 hours of operation in similar temperature conditions provided by EPRI showed varying degrees of structural degradation and in the thickness of the internal and external oxide coating, depending on the location in the boiler. The extent of precipitation processes was compatible with the creep defects. With regular decomposition of the particles in the matrix, no defects were noted.
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6.21 Microstructure of austenitic steels with isolated primary MC carbides (Hernas et al., 2009).
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Point and linear pores, as well as microfissures, were generated, mainly in the areas in which precipitation networks occurred, especially σ phases along the grain boundaries. Changes in the microstructure were based on the varying extents of TiC and M23C6 carbide precipitation processes and the sigma phase, as was the case for decomposition in the matrix. Microstructural changes occurring under the constant effect of increased temperatures and strain conditions directly affect the mechanical properties and processes of nucleation and development of internal defects during creep tests and operation (Dobrzan´ski and Zielin´ski, 2002). Thus, aside from microstructural degradation, the main criterion for assessing the state of a material is the condition of its internal defects. Our understanding of these effects is based on a classification scheme for creep defects developed in the 1980s. It has been verified and used in assessing material states after operation. Degradation tests conducted on ferritic- and austenitic-matrix steels during long-term operation are aimed at developing objective criteria for assessing the state of a material and the component, and later to the determination of residual life and usefulness for extended safe operation.
6.5
Preparation of a classification system for material after operation
The specific degree of material exhaustion is a result of co-occurring structural changes and internal defects during long-term operation under creep. Therefore, the classification prepared, which has already been mentioned earlier, includes the following elements: (a) changes connected with decomposition of pearlite/bainite or tempered martensite, (b) changes connected with development of carbide precipitation processes, (c) changes connected with development of internal defects. The specific structural changes are assigned with the corresponding exhaustion degree. An exhaustion degree is assigned with the main structural class comprising classes of processes of component changes in the structure. Defining the assumptions in this way made it possible to prepare classification of material changes for 12Cr–1Mo–V and 0.5Cr– 0.5Mo–0.25V steels. The classifications of structural changes are accompanied by structural patterns that are illustrations of changes occurring in the alloy steel 12Cr–1Mo–V with tempered martensite structure and the low-alloy steel 0.5Cr–0.5Mo–0.25V with bainite/pearlite and ferrite structure. The classification method involved in assigning the exhaustion degree
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6.22 Classification method of the material condition on the basis of changes in the structure based on component processes in relation to the degree of exhaustion for steels operated under creep conditions.
with the main structural class comprising classes of processes of component changes in the structure in a graphical form is shown in Fig. 6.22. From a practical point of view, the assessment method shown in Fig. 6.23 (for 0.5Cr–0.5Mo–0.25V steel) and Fig. 6.24 (for 12Cr–1Mo–V steel) is of greater importance. On the basis of (a) the analysis of a structural condition, (b) the development degree of precipitation processes, (c) the phase composition of carbide precipitates, and (d) their share and development degree of damage processes, it is possible to assess the exhaustion degree. When a degree of material exhaustion is known, it is possible to determine a safe time of further operation time for the materials analysed for the operating parameters of further operation given by an operator.
6.6
Modeling degradation processes and their use
Modeling microstructural changes that occur in multicomponent systems under conditions of long-term exposure to temperatures and strains is a complex problem that is useful for assessing the state of a material and its residual durability. In practice, modeling is used to describe the kinetics of precipitation processes, which, beyond the generation and development of creep defects, also help determine the strength and durability of structural
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6.23 Assessment method of the material condition operated under creep conditions on the basis of assessment of the state of component processes for 0.5Cr–0.5Mo-0.25V steel.
elements. Models also address durability under creep conditions with uniaxial stress. The literature describes many models that characterize the kinetics of precipitation in steels of various types (Thornton et al., 2003; Rajek, 2005). Unfortunately most of such models describe phenomena that are connected with precipitation of the same phase type. In works that model precipitation processes, a different type of numerical and programming assistance is used; – inter alia DICTRA (diffusion-controlled transformation). These programs permit forecasting of equilibrium states in ‘true’ alloys, with the equilibration time being one of the most important variables. Defining the kinetic processes of phase transformations has become possible, particularly for those transformations controlled by processes of a diffuse nature. Usually the Thermo-Calc system is used for
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6.24 Assessment method of the material condition operated under creep conditions on the basis of assessment of the state of component process for 12Cr–1Mo–V steel.
thermodynamic computations. Other systems include: MatCalc, the Monte Carlo method, and the Ising model. Simulations of M23C6 carbide precipitation are fairly thoroughly described by Faulkner et al. (1998), using the classical theory of nucleation and the increased precipitation in grain boundary areas (Yin and Faulkner, 2003). This model has held up particularly well for the analysis of the precipitation of M23C6 carbides in Ni–Cr austenitic steels in the presence of only chromium (Carolan and Faulkner, 1988; Jiang and Faulkner, 1996a, 1996b; Faulkner et al., 1998; Goodwin et al., 2001). Simulating precipitation becomes more complicated when several secondary phases are present in the microstructure with varying chemical compositions. Adapted to multicomponent materials (Carolan and Faulkner, 1988;
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Jiang and Faulkner, 1996b), this model predicts that the precipitates form cap-shaped structures. The computational process runs in four phases: . . . .
I – determining the decomposition of the alloy components in the matrix, which reduces to defining the segregation of alloy components at the grain boundaries; II – determining concentrations in areas of potential precipitate nucleation at the grain boundaries, as a function of temperature, taking into account the concentration of the alloy components in those areas; III – defining an increase in precipitates arising in the second stage as a result of the diffusion processes from the grain core to the boundary, as a function of the heating temperature; IV – defining the time range in which all components that participate in the precipitation process are located in the vicinity of the grain boundary, and the proliferation of precipitates occurs concurrently with the dissolution of smaller precipitates.
It should be noted that, in this case, the results obtained depend on the average precipitate size. This model is helpful for analyzing the process of sigma phase precipitation. A further development of this model would entail considering the influence of the remaining alloy components on the kinetics of nucleation of, increase in, and dissolution of precipitates. The introduced modification assumes that each alloy addition that decreases the dissolution of chromium in ferrite (W, Al, Si, Mn, Co, Cu, Ni, Mo) causes an increase in the speed of the nucleation and proliferation of precipitates that are rich in chromium. The results of theoretical models for precipitation in new martensitic steels are presented inter alia in the words of Gustafson (Gustafson et al., 1999; Gustafson and Ha¨ttestrand, 2002), who implemented them to simulate the proliferation of carbonitride precipitates in P92 steel and, for example, in the works of Ghosh and Olson (2001), Bernhard et al., (2001), Knezevic et al. (2008), Schneider and Inden (2005), and Gaude-Fugarolas and de Carlan (2008). One method proposed using a modified Johnson–Mehl–Avrami equation, as presented in the works of Bhadehesia and others (Christian, 1975; Robson and Bhadeshia, 1996a, 1996b, 1997a, 1997b, 1997c; Brun et al. 1999). This work developed kinetic models for the carbide phase precipitates and Laves phases in martensitic steels used in power engineering, under conditions of long-term operation at high temperatures. In practice, this problem reduced to working out TTT-type (time, temperature, transformation) diagrams that characterized the processes of precipitation during tempering as a function of chemical composition and tempering temperature. The results of these tests may be used both for optimizing the chemical
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6.25 Model for microstructural evolution in ferritic–pearlitic/bainitic steels.
6.26 Models for microstructural evolution in martensitic steels.
6.27 Model for microstructural evolution in austenitic steels.
composition of steel and for predicting the behavior of materials during long-term creep or operation conditions. Practical analytical assessments of the material states durability predictions have been described by the more useful, simple models of the microstructural states that correspond to defined degrees of material degradation. An example of such a model is the model for microstructural evolution in ferritic–pearlitic, bainitic, martensitic, and austenitic steels, which may be used in algorithms for assessing structural degradation after long-term operation. Schematically these simple models, in the case of ferritic–pearlitic/bainitic, martensitic and austenitic steels, are presented in Figs 6.25 to 6.27. The objectivity of those models is based on results obtained for a large number of actual statistical data for materials after creep or after operation, which is verified in creep tests. The complex durability assessments for pressure components combine the methods that enable prediction of microstructural changes with the methodologies for assessing durability and residual durability under creep conditions. For example, Faulkner
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proposed a method that combined Monte Carlo methods with CDM (continuum creep damage mechanics) methods (Faulkner and Yin, 2003; Faulkner, 2006).
6.7 .
.
.
6.8
Conclusion An array of non-destructive methods (NDT), computational and modeling, is used for the objective assessment of a material’s state. Safety precautions assert that the least optimistic result should be considered for durability or residual durability. The method and measurement chosen for the assessment varies as a function of the type of component and the nature of its function. The degree of material degradation during long-term operation is mainly determined by the extent of creep defects. Regardless of the degree of structural degradation, these defects are critical and are decisive in assessing an element’s usefulness for further operation. Tests show that the first appearance of creep defects in steels with a ferritic matrix implies degradation of the material f to approximately 0.55–0.6, whereas in austenitic steels f falls within the range 0.6–0.7. The criteria for assessing the level of structural degradation are specific to a particular group of steels and to processes that determine the stability of functional properties. Therefore, mild-alloy steel decomposition is dominated by areas enriched in pearlite/bainite with spheroidization of carbide, whereas martensitic steel decomposition is dominated by the precipitation and disappearance of martensite with the appearance of subgrains. Degradation assessment in austenitic steels from group 18-10 arises from the precipitation processes, and is particularly dependent on the morphology and location of the σ phase.
References
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Bouchard, P J, et al. (2004), ‘Quantification of creep cavitation damage around a crack in a stainless steel pressure vessel’, Acta Materialia, 52, 1, 23–34. Brun, F, et al. (1999), ‘Theoretical design of ferritic creep-resistant steels using neural network, kinetic and thermodynamic models’, Materials Science and Technology, 15, 547–554. Cane, B J (1986), ‘Present status of predictive methods for remnant life assessments and future development’, Materials Forum, 9, 1–2. Carolan, R A and Faulkner, R G (1988), ‘Grain boundary precipitation of M23C6 in an austenitic steel’, Acta Metallurgica, 36, 257–266. Christian, J W (1975), Theory of Transformations in Metals and Alloy, 2nd edn, Part I, Pergamon Press. Oxford. Cipolla, L (2009), ‘Formation of Z-phase in a 12%Cr CrVNbN model, steel’ in Creep and Fracture in High Temperature Components, 2nd ECCC Creep Conference, DEStech Publications, pp. 863–876. Danielsen, H K and Hald, J (2006), ‘Behavior of Z-phase in 9–12% Cr Steels’, Energy Materials, 1, 49–57. Dobrzan´ski, J and Milin´ski, P (1992), ‘Possibilities of extending the period of operating time of pressure elements working in the increased temperature’, Lecture Note of IFM, vol. 38, no. 1–2 (in Polish). Dobrzan´ski, J (1995), ‘Analysis of structure and properties changes of the 1Cr– 0.5Mo type steel subjected to long-term creep as the basis for forecasting the life of the power industry equipment components’, PhD Thesis, Silesian University of Technology, Katowice (unpublished) (in Polish). Dobrzan´ski, J (2002) ‘Evaluation of the state of material in the diagnostics of elements of energy equipment’, in Proceedings of the VII National Energy Conference, Rydzyna (in Polish). Dobrzan´ski, J and Zielin´ski, A (2002),’Evaluation of operating life-time steel working above the border temperature based on shortened creep test’, in Proceedings of the IX Seminary on Materials Examinations for the Purposes of the Power Station and the Power Industry, Zakopane (in Polish). Dobrzan´ski, J (2003), ‘Processes of internal failure in low-alloy Cr–Mo steels working above the border temperature’, in Proceedings of X Seminary on Materials Examinations for the Purposes of the Power Station and the Power Industry, Zakopane (in Polish). Dobrzan´ski, J (2005), ‘The classification method and the technical condition evaluation of the critical elements material of power boilers in creep service made from the 12Cr–1Mo–V’, Journal of Materials Processing Technology, 164– 165, 785–794. Ennis, P J, Zielin´ska-Lipiec, A and Czyrska-Filemonowicz, A (1998), ‘Quantitative comparison of the microstructure of high chromium steels for advanced power stations’ in A Strang, J Cawley and G W Greenwood (eds), Microstructure of High Temperature Materials, No. 2, The Institute of Materials, London, p. 135. EPRI (2002) ‘Remaining life assessment of austenitic stainless steel superheater and reheater tubes’, EPRI, Palo Alto, California, 2002, Technical Report. 1004517. Faulkner, R G, et al. (1998), in A Strang et al. (eds), Microstructural Stability of Creep Resistant Alloys for High Temperature Plant Applications’, The Institute of Materials, London, pp. 431–443. Faulkner, R G and Yin, Y (2003), ‘CDM modelling and alloy development of
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ferritic/martensitic steels for ultra-super critical boiler applications’, Oak Ridge National Lab Report, USA, pp. 1–22. Faulkner R G (2006), ‘Continuum damage mechanics modelling based on simulations of microstructural evolution kinetics’, Materials Science and Technology, 22, 929–936. Gaude-Fugarolas, D and de Carlan, Y (2008), ‘Modelling precipitate distribution in reduced-activation steel’, Journal of Nuclear Materials, 374, 109–115. Ghosh, G and Olson, G B (2001), ‘Simulation of paraequilibrium growth in multicomponent systems’, Metallurgical and Materials Transactions, 32A, 455– 467. Goodwin, C C, Faulkner, R G and Spindler, M W (2001), ‘Modelling long-term precipitate growth in austenitic steels for boiler applications’, in Ageing Studies and Lifetime Extension of Materials, Kluwer Academic, Oxford, pp. 289–395. Grabke, H J, et al. (1998) ‘Effects of grain size, cold working, and surface finish on the metal-dusting resistance of steels’, Oxidation of Metals 50, 241–254. Gustafson, A and Ha¨ttestrand, M (2002), ‘Coarsening of precipitates in an advanced creep resistant 9% chromium steel–quantitative microscopy and simulations’, Materials Science Engineering A, 333, 279–286. Gustafson, A, Ho¨glund, L and A˚gren, J (1999), ‘Simulation of carbo-nitride coarsening in multicomponent Cr steels for high temperature applications’ in R Wiswanathan and J Nutting (eds), Conference Proceedings on Advanced Heat Resistant Steel for Power Generation, San Sebastian, Spain, pp. 270–276. Hahn, B, Spiegel, M and Bendick W (2009), ‘T23, T24 and VM 12–structure, properties, application’, in VGB Workshop: Material and Quality Assurance, Copenhagen, May 13–15, 2009, www.vgb.org/en/event_mqa_pres.html? dfid=23699. Hald, J (2008), ‘Microstructure and long-term creep properties of 9–12% Cr steels’, International Journal of Pressure Vessels Piping, 85, 30–37. Hald, J (2009), ‘Status of the martensitic creep resistant 9–12% Cr steels’, in Creep and Fracture in High Temperature Components, 2nd ECCC Creep Conference, DEStech Publications, p. 3. Herans, A (1992), ‘The issue of residual life-time from materials science point of view’, Material Engineering, vol. 4-5 (in Polish). Hernas, A and Dobrzan´ski, J (1995), ‘Correlation between phase composition and life-time of 1Cr–0.5Mo steels during long term service at elevated temperature’, Journal of Materials Processing Technology, 53, 101. Hernas, A and Dobrzan´ski, J (2003), Life-time and Damage of Boilers and Steam Turbines Elements, Part 2, Silesian University of Technology Press (in Polish). Hernas, A and Maciejny, A (1998), Creep-resistant Alloys, Ossoilneum Press, Wrocøaw (in Polish). Hernas, A, et al. (2009), ‘Austenitic steel and nickel superalloys used in the construction of supercritical boilers and waste plants’, in A Hernas (ed.) Materials and Technology for Constructions of Supercritical Boilers and Waste Plants, SITPH Press, p. 111 (in Polish). Husemann, R U (2005), ‘Advanced materials for AD700 boilers’, presentation presented to European Conference AD700–Advanced (700 8C) PF Power Plant: A Clean Coal European Technology, CESI Auditorium, Milano.
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‘INCONEL alloy 625’, Special Metals Corporation Products, www.specialmetals. com/products. Iseda, A, et al. (2007), ‘Creep properties and microstructure of SUPER304H, TP347HFG, HR3C steels, in Fifth International Conference on Advances in Materials Technology for Fossil Power Plants, Florida. Jiang, H and Faulkner, R G (1996a), ‘Modelling of grain boundary segregation, precipitation and precipitate-free zones of high strength aluminium alloys – II. Application of the models’, Acta Materialia, 44, 1865–1871. Jiang, H and Faulkner R G (1996b), ‘Modelling of grain boundary segregation, precipitation and precipitate-free zones of high strength aluminium alloys–I. The model’, Acta Materialia, 44, 1857–1864. Kautz, H R (1997), Ageing of Materials and Methods for the Assessments of Lifetimes, CAPE 97, Balkema, Roterd. Knezevic, V, et al. (2008), ‘Design of martensitic/ferritic heat-resistant steels for application at 650 8C with supporting thermodynamic modeling’, Materials Science and Engineering A, 477, 334–343. Masuyama, F (1998), ‘Steam plant material development in Japan’, in J LecomteBeckers et al. (eds), Materials for Advanced Power Engineering 1998, Part III, Ju¨lich GmbH, Forschungs-zentrum, Ju¨lich, p. 1807. Masuyama, F (2004), ‘Alloy development and materials issues with increasing steam temperature’, in Proceedings from the Fourth International Conference on Advances in Materials Technology for Fossil Power Plants, Hilton Head Island, USA, p. 35. Najgebauer, E and Patrycy, A (2008), ‘Commitment of the Polish power industry to the EC’, www.geoland.pl/ dodatki/energia/xxxv-energ.belch.html (in Polish). Needham, N G (1985), IIW Creep Committee. Neubauer, B and Wedel, U (1984), ‘NDT: replication avoids unnecessary replacement of power plant components’, Power Engineering, 88, 44. Pasternak, J, Hernas, A and Milin´ski, P (1997), ‘New martensitic steels for supercritical boilers’, Energetics, 10, 543 (in Polish). Polish Standards PN-75/H-84024 (1975), ‘Steel for the work at increased temperatures –types’ (in Polish). Prager, M (2006), ‘Creep strength-enhanced ferritic steels’, presentation presented to WRC/MPC Data, ASME SCII TG, Henderson, Nevada. Rajek, J (2005), ‘Computer simulation of precipitation kinetics in solid metals and application to the complex power plant steel CB8’, Graz University of Technology, Graz. Rautio, R and Bruce, S, ‘Sandvik Sanicro 25, a new material for ultrasupercritical coal fired boilers’, Sandvik Materials Technology. Regis V and D’Angelo, D (1989) ‘Advanced residual life methods for improved design and operation’. Nuclear Engineering and Design Journal, 116, 399–406. ‘Report 2030. The influence of proposed EU regulations within the scope of implementing European strategy of the power industry development free from CO2 emission . . . ’, System Research ‘EnergSys’ Sp. z o.o.2008, Warsaw, 2008 (in Polish). Robson, J D and Bhadeshia, H K D H (1996a), ‘A new model for simultaneous transformation kinetics in power plant steels’, in Proceedings of Creep ’96, on
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Viswanathan, R (1994), ‘Life assessment of superheater/reheater tubes in fossil boilers and others’, Journal of Pressure Vessel Technology, 116, 1–16. Viswanathan, R, et al. (2005a), ‘US program on materials technology for ultra supercritical coal power plants’, Journal of Materials Engineering and Performance, 14, 3, 281. Viswanathan, R, Purgert, R and Rao, U, (2005b), ‘Materials and technology for advanced coal power plants’, in Proceedings of the 1st International Conference on Super High Strength Steels, Rome. Yin, Y F and Faulkner, R G (2003), in A Strang, R D Conroy, W M Banks, M Blackler, J Leggett, G McColvin, S Simpson, M Smith, F Starr and R W Vanstone (eds), 6th International Charles Parsons Turbine Conference on Engineering Issues in Turbine Machinery, Power Plant and Renewables, Institute of Materials, London, pp. 525–535. Zheng-Fei, H and Zhen-Guo, Y (2004), ‘An investigation of the embrittlement in X20CrMoV12.1 power plant steel after long-term service exposure at elevated temperature’, Materials Science and Engineering A, 383, 2, 224–228. Zielin´ska-Lipiec, A, Czyrska-Filemonowicz, A and Ennis P J (1998), in Proceedings of International Conference on ‘Materials for Advanced Power Engineering and Other Applications’, Liege, Belgium. Zielinska-Lipiec, A, et al. (2004), ‘Microstructural development of VM12 steel caused by creep deformation at 6258C’, in Materials for Advanced Power Engineering 2006/ Proceedings of the 8th Liege Conference, Liege, Belgium, pp. 1077–1086. Zielinska-Lipiec, A, Kozieø, W and Czyrska-Filemonowicz, A (2009a), ‘Microstructure and properties of new martensitic 12% Cr steels for supercritical power plants’, Metallurgy, Metallurgical Engineering News, 76, 4, 269–275 (in Polish). Zielin´ska-Lipiec, A, Kozieø, W and Czyrska-Filemonowicz A (2009b), ‘TEM study of secondary precipitates influencing creep strength of martenistic VM12 steel’, in Microscopy Conference, Graz, Austria. Zielin´ska-Lipiec, A, Kozieø, W and Czyrska-Filemonowicz, A (2009c), ‘Changes of microstructure of 12%Cr martensitic steel lowering their creep resistance’, in A Hernas (ed.) Materials and Technology for Constructions of Supercritical Boilers and Waste Plants, SITPH Press, p. 102 (in Polish). Zielin´ska-Lipiec, A, et al. (2009d), ‘Effect of creep deformation at 625 8C on microstructure development of VM12 steel’, in 2nd ECCC Creep Conference ‘on Creep and Fracture in High Temperature Components: Design and Life Assessment issues, Zurich, Switzerland, pp. 619–627.
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7 Boiler steels, damage mechanisms, inspection and life assessment A . S H I B L I , European Technology Development, UK
Abstract: This chapter covers the general introduction to the materials used for power plant boilers and a brief history of their development. Materials ageing and component failure mechanisms are then outlined. This is followed by a somewhat more detailed coverage of the microstructural aspects of 9–12 %Cr martensitic steels, the formation of various precipitates during their production and operation at high temperature and their role in the creep strengthening of these alloys. As the use of high Cr steels is now approaching the mid-life stage, cracking and failure has been experienced in many plants worldwide. However, it has been realised that inspection and early stage creep damage detection in these steels is somewhat more challenging than the traditional low alloy steels and in view of this, inspection and monitoring techniques recently developed or being developed for the high Cr martensitic steels have been discussed. Key words: martensitic steels, 9Cr steels, microstructural deterioration, precipitates, NDE techniques, SFM/AFM, potential drop technique, integrity assessment, damage assessment.
7.1
Introduction
This chapter deals with three aspects of power plant boilers, all of which relate to materials issues. To start with, it traces the history of the use of boiler materials and their development over the last half a century. It then covers the metallurgy and microstructural issues that concern the performance of these steels at high temperature. More emphasis has been placed on the relatively newer 9–12 %Cr martensitic steels due to the increasing use of these not only in the new ultra-supercritical power plant for operation at temperatures of between 568 and 625 oC but also as replacement thick section materials in older plant for use at lower 272 © Woodhead Publishing Limited, 2011
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temperatures of 540 to 568 oC or in the heat recovery steam generators (HRSGs), which operate even at lower temperatures. Furthermore, these steels are now being seriously considered for the new Generation IV nuclear power plants and fast breeder reactors, which will operate at much higher temperatures compared with the conventional PWRs, which operate below the creep range. The use of high Cr martensitic steels first started in the UK and Japan in 1989. In the UK ASME P91 (9 %Cr martensitic steel) is used as a replacement material in the existing power plant with operating temperatures of up to 568 oC, while in Japan P91 is mainly used in new power plants at higher temperatures of up to 600 oC. Soon afterwards 9Cr martensitic steels (including the European steel P911) were used in new plants in Denmark and Germany. However, during the last 10 years or so the frequency of their use has significantly increased. Both with the increase in the frequency of their use and time in service the weldments in these steels have experienced cracking and failure in the Type IV position in a similar way to the traditional low alloy steels, and in some cases with even higher frequency than the low alloy steels [1]. It is now clear that these new steels are not as ‘ideal’ as initially thought and as they age they show their own ‘mid-life’ crisis. This means that methods must be found or developed to inspect and monitor them. However, it has been observed that with the use of traditional non-destructive examination (NDE) techniques (ultrasonic, replication, etc.) these steels do not show creep cavitation until about 70 to 80 % of life [2]. Thus by the time damage is discovered it can be too late to save the damaged components by repair or to carry out timely replacement. This has meant that new NDE methods have to be developed to detect such damage at an early stage, possibly at about 20 % life, so that remedial steps can be taken at the earliest opportunity. The development of such methods has been described in the last part of this chapter. This chapter also discusses cracking or failure mechanisms that affect power plant boilers. The main mechanisms such as creep, fatigue and creep– fatigue interaction have been discussed in some detail, but other mechanisms have also been briefly described or simply referred to. In a chapter of limited size it is not possible to do justice to these important topics but an effort has been made to introduce different aspects and throw some light on the new and latest developments. Additional references have been provided for the reader to enhance personal knowledge of the subject and to seek further information. Unfortunately, as happens with most of the latest developments, not all work has been published as yet so some references had to be made to unpublished reports.
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7.2
Boiler materials, metallurgy and microstructure
7.2.1 Materials for boiler tubing and headers Over the last half a century power generation boiler capacity has increased very significantly, with steam pressure rising by about 6 times to 25 MPa and steam temperature rising from 450 8C to 568 8C. Demand for increased efficiency in the early to mid 1980s was met with efforts to improve the design of the plant and the materials that can withstand higher temperatures and pressures. A steam temperature of up to 596 8C was achieved in the plant built in the early 1990s, while a temperature of 600 8C was achieved by around 2000. Equipment making up boilers using tubing and headers is of various types and designs. Pressurised components (pressure vessels) used at relatively low water and steam temperatures include: steam drums for natural and forced circulation type boilers, mixing vessels and distributing vessels for combined circulation type boilers (which combine the characteristics of forced circulation and once-through), as well as steam–water separators for startup and variable pressure operation of once-through type boilers. Carbon steels are used for drums, while 0.5Mo steels and SA-387 Grade 11 are used for steam–water separators, made of thick plates in the range of 50 to about 170 mm. Pressure vessels exposed to both boiler steam/water and combustion gas include economisers, furnace walls, superheaters, reheaters, etc. These components consist of small diameter steel tubes. In the 538 8C class boilers, mainly 1.25Cr0.5Mo (T11) and 2.25Cr1Mo (T22) are used, while 9Cr1Mo (T9, having superior corrosion resistance) was also adopted for some time. However, a disadvantage of T9 steel is its high susceptibility to weld cracking, and for this reason its application has been limited. Steels such as 2.25Cr1.6W (T23), 9Cr2Mo (STBA27) and 9Cr1MoVNb (T91), 9Cr0.5Mo1.8W (T92), 12Cr1Mo1W (HCM12) and 11Cr0.4Mo2W (HCM12A, T122) feature improved weldability, and these have come into widespread use instead of the conventional 9Cr1Mo (T9) bainitic steel. In 568 8C and above class boilers, low alloy steels and 9Cr steels have insufficient strength and corrosion resistance for the hot end of the superheater and reheater, and so 18Cr8Ni stainless steels (TP304H, TP321H, TP347H, TP347HFG, etc.) are used.
7.2.2 Materials for steam pipework Historically, the type of material used for the construction of steam pipework systems has been similar to that used for the superheater outlet headers of the boiler, i.e. low alloy ferritic steels. However, more recently 9 %Cr martensitic steels (P91 and P92) have been utilised in order to cope
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with the more demanding steam conditions of a modern power plant. The use of these higher strength steels is also considered more beneficial as it reduces the component wall thickness, which results in lower temperature gradients across the wall and thus reduces the damage and cracking due to creep–fatigue interaction. This is especially important as many power plants now operate in the plant cyclic or load following regime and are thus subject to creep–fatigue interaction. Of the conventional low alloy steels, 2.25Cr1Mo and 0.5CrMoV have been widely used for steam piping in a conventional boiler plant. Traditionally 0.5CrMoV steel has been employed in the UK and some other countries, whereas in Germany and elsewhere the martensitic 12 %Cr steel X20CrMoV11-1 was used. More recently the modified low alloy ferritic/bainitic steels, P23 and P24, have been developed as cost-effective alternatives to the 9 %Cr martensitic steels for use as piping and headers in the 500–550 8C steam range. The steels P91, P23 and P24 are also used for retrofit applications in place of the conventional low alloy ferritic steels. The advantage of the use of P23 and P24 steels is that the absolute temperature control in the welding and post-weld heat treatment of these steels is not as critical as for P91 and P92.
7.2.3 Metallurgy of low alloy ferritic steels The low alloy steels of relevance are those based on the chromium– molybdenum series with additions of elements such as vanadium, niobium, titanium and others. They are often designated as creep-resistant steels. Service design requirements also require other more basic properties such as elevated temperature proof strength (sometimes tensile strength) and impact/toughness resistance. The most commonly used ‘standard’ steel grades, which are used for tubes and pipes in steam boilers and heat exchangers, are shown in Table 7.1. These steel types are usually referred to as ‘ferritic’ steels and usually have a bainitic microstructure, which generally has the best creep resistance [3], although long term service/testing strength depends mainly on the fine precipitates within the bainitic microstructure [4]. Different alloying elements can have different effects on the microstructure. For example, chromium, apart from conferring oxidation resistance, retards the transformation on cooling from austenite such that when it is present in the range 9 to 12 % martensite becomes the predominant microstructure. Similarly, different heat treatment types and cooling rates can result in different microstructures [5], giving different properties to the same chemical composition alloy. All steels are in a meta-stable state after tempering and so extended time during service exposure at high temperatures will produce further metallurgical changes that can lead eventually to instability, usually beyond
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Table 7.1
Typical low alloy steels used for tubes and pipes
Common name
EN name
ASTM name
0.3 % or 0.5 %Mo 1 %CrMo 2.25 %Cr–1%Mo 0.5 %CrMoV
16Mo3 (for 0.3 %Mo steel only) 13CrMo4–5 10CrMo9–10 14MoV6–3
T1 (for 0.5 %Mo steel only) T11 T22 —
the design life. The most stable carbide formed during tempering in the low alloy steels (without vanadium) is Mo2C. It has low nucleation energy and forms as very fine needle-shaped particles, which grow very slowly at normal service temperatures due to a high coherency between the lattice structures of the carbide particle and the surrounding ‘ferrite’ matrix. For steels containing vanadium, another distinctive precipitate forms, namely V4C3, which has a plate-like appearance in the microstructure. Like Mo2C this is also very stable and confers high creep resistance.
7.2.4 Metallurgy of 9–12 %Cr martensitic steels The 9–12 %Cr martensitic steels are given a normalising and tempering heat treatment in order to develop the desired microstructure with optimum properties, i.e. tempered martensite, (Fig. 7.1). On cooling from the austenitisation temperature, a martensitic microstructure is obtained for a
7.1
P91 tempered martensite microstructure.
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wide range of practical cooling rates. However, a weaker ferrite phase can form on very slow cooling and in practice this can happen in very thick section components. The martensite has the typical lath structure within the prior austenite grain boundaries. The chemical compositions of these steels are shown in Table 7.2. The effect of some of the important alloying elements is shown below. Mo and W The microstructure of the 9 %Cr steels, with Mo, W, Nb and V as important alloying elements, consists of tempered martensite in which islands of δ ferrite form during cooling. The martensite has a typical lath structure within the prior austenite grain boundaries. The long term stress rupture behaviour of these steels is mainly affected by: . . . .
solid solution hardening by C, Cr, Mo and W; the type, size and distribution of fine precipitates; dislocation density within the sub-structure; the austenite grain size, although this has a somewhat lesser effect.
Upon cooling from normalising/austenitising and especially during tempering, carbide formation takes place. The more important carbides that form and their transformations are listed below: . . .
M3C ? M7C3 ? M23C6; MC ? M2C ? M6C; various carbonitrides.
The formation of various precipitates at different temperatures is shown in Table 7.3. It is now well known that Mo and W both belong to the same group as Cr and thus have a similar effect in high Cr steels; i.e. their principal contribution to enhancing the creep rupture strength is due to their solid solution strengthening. In addition, they contribute to rupture strength through precipitate hardening by the formation of carbides and carbonitrides. However, as the atomic weight of W is nearly twice the atomic weight of Mo, to achieve the same effect of strengthening twice as much W is needed. It is also known that the low diffusion rate of W leads to a stabilising effect on the carbides and carbonitrides. Tests in Japan and elsewhere in the 1960s and 1970s found that the maximum creep strength was obtained for an Mo equivalent (%Mo + 0.56%W) of 1.5. However, because of the increase in δ ferrite formation, a decrease in long term rupture strength was observed. The interesting thing, however, was that high strength was only obtained when Mo and W were present together. Both in Japan and in European COST programmes 1 %Mo and 1 %W were found
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max. 0.15 0.25–1.00 0.30–0.60 max. 0.025 max. 0.025 — — 8.00–10.00 0.90–1.10 — — — — — — — —
0.08–0.12 0.20–0.50 0.30–0.60 max. 0.020 max. 0.010 max. 0.40 — 8.00–9.50 0.85–1.05 — 0.18–0.25 0.06–0.10 max. 0.02 0.030–0.070 — max. 0.01 max. 0.01
P91 0.07–0.13 max. 0.50 0.30–0.60 max. 0.020 max. 0.010 max. 0.40 — 8.50–9.50 0.30–0.60 1.50–2.00 0.15–0.25 0.04–0.09 max. 0.02 0.030–0.070 0.001–0.006 max. 0.01 max. 0.01
P92 0.09–0.13 0.10–0.50 0.30–0.60 max. 0.020 max. 0.010 0.10–0.40 — 8.50–9.50 0.90–1.10 0.90–1.10 0.18–0.25 0.06–0.10 max. 0.040 0.050–0.090 0.0005–0.0050 — —
E911
0.17–0.23 max. 0.50 max. 1.0 max. 0.030 max. 0.030 0.30–0.80 — 10.0–12.5 0.80–1.20 — 0.25–0.35 — —
— —
— —
X20
max. 0.15 0.20–0.65 0.80–1.30 max. 0.030 max. 0.030 max. 0.30 — 8.5–10.5 1.70–2.30 — 0.20–0.40 0.30–0.55 —
EM12
* Compositions of T9, P91, P92 and P122 according to ASTM A213/A335; others according to European sources. ** Non-martensitic steel.
C Si Mn P S Ni Cu Cr Mo W V Nb Al N B Ti Zr
T9**
Table 7.2 Chemical compositions of 9–12 %Cr steels*
0.07–0.14 max. 0.50 0.25–0.60 max. 0.020 max. 0.010 max. 0.50 0.30–1.70 10.00–11.50 0.25–0.60 1.50–2.50 0.15–0.30 0.04–0.10 max. 0.020 0.040–0.100 0.0005–0.005 max. 0.01 max. 0.01
P122
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Tempering below 700 8C
Tempering above 700 8C creep exposure Long term creep exposure Creep exposure
M2X(Cr2N)
Secondary MX(VN)
Tempering (> 1.6 % Mo) long term creep exposure
Tempering
M 6X
AIN
Cr(V, Nb)N Z-phase Laves phase Fe2(Mo, W)
Solidification
Tempering
M23C6
Nb(C, N)
Formed during
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Grain and sub-grain boundaries
Grain and sub-grain boundaries, on M23C6 Grain and sub-grain boundaries on M23C6
Inside sub-grains, M laths and/or dislocations M23C6, Nb(C, N)
Prior austenite grain and M lath boundaries Austenite grain boundaries Inside sub-grain and M laths
Preferred precipitation
Dissolution of Nb(C, N), M2X and VN Dissolution above 650 8C, fast coarsening at 600 8C High coarsening rate, dissolution of M2X, Nb(C, N); decrease of Mo, W in SS High coarsening rate, reduction of N in SS
Restriction of grain growth Dissolution during creep at high temperatures High dimensional stability during creep
Medium growth rate during creep
Characterised by
Basic characteristics of minor phases in 9–12 %Cr steels [6]
Precipitate
Table 7.3
Poor plasticity
Decrease in SS, lowering PS, poor plasticity
Decrease in SS lowering PS, poor plasticity
Lowering PS
Precipitation strengthening
Precipitation strengthening
Precipitation strengthening (PS)
Main contribution to
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to be optimal. However, as Mo is a stronger δ ferrite former (which weakens the rupture strength and impact strength), the Mo content of 0.5 % and W of about 2 % are more common, as in the case of P92 and HCM12A/P122. It is also known that with increasing test/creep rupture duration, W and Mo are also increasingly incorporated to a certain extent into the M23C6 phase. The matrix therefore becomes depleted in these elements, which are not very effective for strengthening when present in precipitates compared with when they are dissolved in the matrix. In the typical 12 %Cr steel X20CrMoV11-1 is the Cr rich M23C6 carbide. The addition of V, Nb and N in P91 causes additional precipitation of MX particles, where M consists of V or Nb and X consists of N or C. This precipitate is most stable over the long term component exposure to creep temperature. The precipitation reactions during heat treatment have a significant effect on the long term creep strength of the material. The carbide and nitride particles precipitate on: . .
prior austenite grain boundaries; ferrite subgrain boundaries and inside the sub-grains.
V and Nb In 9Cr steels the solubility of the V rich MX phase (which is formed during tempering) is very high. MX is practically VN (with very small amounts of Nb and Cr) and these precipitates are very fine particles. Dissolved V and N at the austenitising temperature promote the precipitation of VN during tempering and mainly during creep exposure [7]. These particles have a very low coarsening rate during creep exposure. Nb forms coarse particles of niobium carbonitride Nb(C,N) and can lower the amount of N available for VN formation. However, its usefulness is that it serves as the nucleation site for VN precipitates. Nitrogen This was studied by Nippon Steel Corporation in 1993, who reported an increasing creep rupture strength with an increase in nitrogen content [8]. More recently the effect of N was studied by Wachter and Ennis [8], who found that in P92 the relative proportion of carbonitrides was higher and the precipitates were finer in the cast with a higher amount of nitrogen (0.49 % compared with 0.42 %).
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Al/N ratio Work carried out at VITKOVICE Research Institute in the Czech Republic recommends minimising the Al content of these steels. Their work showed that the lowering of the Al content from 0.03 % to 0.003 % resulted in an increase in the creep rupture strength of about 15 % [7]. This was attributed to the observation that aluminium nitride creates large sized precipitates often situated on grain boundaries, which impair creep ductility very significantly. The AlN precipitation can reduce the fine VN particle precipitation. The creep rupture strength in 100 000 hours of a high Al cast of 12Cr steel was observed to be lower by as much as 30 % compared with other casts with lower Al contents. This showed that coarse AlN particles cannot contribute to precipitation strengthening. Laves phase A further decisive factor in the creep resistance is the formation of the Laves phase. Precipitation of the Laves phase occurs during creep and the amount of Laves phase close to equilibrium at 600 oC/1112 oF can be expected in most Cr modified steels at about 10 000 to 20 000 hours. Although during this period the Laves phase can contribute to precipitation strengthening, it is known that this phase coarsens very rapidly and the precipitation strengthening due to this phase is lost very rapidly [7]. A significant amount of W and Mo, which are important for the solid solution strengthening of the matrix, is removed from the matrix when the Laves phase forms, thus reducing the creep strength.
7.2.5 Metallurgy of austenitic stainless steels The Type 300 stainless steels were used in a large number of plants in the UK and USA in the 1960s, since at that time the only practical way to raise power plant efficiencies to levels closer to 40 % was to increase operating temperatures and pressures. This obviously required improved creep strengths for superheater, reheater, and in some cases steam lines and valves. Furthermore, increased metal temperatures had reached the point at which fireside attack was becoming a serious issue. The most advanced of these plants in the USA was that at Eddystone B, where the initial design conditions were 650 /565 /565 C and 345 bar stop valve pressure. These levels were subsequently down-rated to an inlet temperature of 605 8C because of thermal fatigue. In the UK, the most advanced unit was Drakelow C, in which the steam temperatures were 593/ 566 8C and the pressure was 241 bar. Here, severe thermal ratcheting leading to continued distortion of the steam pipework occurred. Another super-
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critical plant in the USA, Philo 6, has operated successfully with steam conditions of 610 /565 /538 8C. It is noteworthy that the reheat temperatures in these plants were significantly below the high pressure (HP) steam conditions. This apparently resulted from the fact that the steam volumes at the lower pressures in the reheat system required larger diameter pipework. Fabrication limitations for larger diameter stainless steel lines, plus concern about thermal fatigue, resulted in the need to operate the reheat systems to run at lower temperatures. A number of the plants in the UK operated at superheater and reheater temperatures of 566 8C and needed to use the Type 300 austenitics for both. These materials were, and are, basically, of the 18/8 type, that is Fe18Cr8Ni. The main alloys that were used were: . . . .
Type Type Type Type
304: 316: 321: 347:
Fe18Cr8Ni Fe18Cr10Ni2.5Mo Fe18Cr10NiTi Fe18Cr10NiNb
All of these alloys had been developed primarily for use in aqueous corrosion environments and their good creep properties were somewhat fortuitous. There was also another alloy, Type 310, Fe25Cr20Ni, whose formulation was intended to give good oxidation resistance at temperatures up to 1000 8C and was only used as a cladding alloy to give resistance to fireside corrosion. Despite the high alloy content of these materials, superheater tubing was susceptible to fireside attack. A large programme of work was mounted by the then CEGB in the UK to overcome this problem, but the results were not completely successful and the coatings that were required added to the cost of tubing. Steam-side exfoliation of oxides was another problem. In the last coal-fired station built in the UK, Drax, the more advanced austenitic steel Esshete 1250 was used. This steel incorporated 6 % manganese in the alloy composition to improve weldability. The material was also used in Drakelow C. Most other countries took a more conservative approach and opted for steam temperatures of under 540 8C, which allowed them to use low alloy ferritic steels. At that time, the standard materials were T22 (2.25Cr–1Mo) in fossil fuel power plants. In the UK, however, T9 (9Cr–1Mo) was used in AGR (advanced gas-cooled reactor) units. Both of these alloys are of the ferritic/bainitic type but where fireside corrosion was likely to be a problem, X20, the 12 %Cr martensitic stainless steel, was specified.
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Damage mechanisms and component failure
7.3.1 Creep Creep resistance is dependent on microstructure – creep-resistant steels contain large numbers of very fine precipitates in order to impede the motion of dislocations. However, most of the precipitates coarsen with time, their distribution becomes less dense and thus they become less effective in resisting creep deformation. The role of grain boundaries is an important factor in the creep deformation process. At low temperatures, grain boundaries impede dislocation movement but once temperature conditions reach a creep-initiating level they provide the process with sources of atoms and vacancies that permit dislocation climb during creep. Grain boundaries at high temperature reverse their role of resisting deformation to aiding deformation. After sustained periods of time under creep conditions, cavities form on the grain boundaries. Initially the cavities are isolated but later cavities aligned on the same grain boundaries will link together to form microcracks; as the process continues, further the microcracks will combine to create macrocracks, leading to the failure of the material by intergranular fracture. By analysing the development of creep cavitation (and the microstructure) the extent of the creep damage in a component can be estimated. Figure 7.2 shows an example of creep cavities on grain boundaries. Figure 7.3 shows the evolution of creep damage from cavity formation to macrocracks and eventual creep rupture failure.
7.2 Creep cavities on grain boundaries.
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7.3 Creep deformation process evolution from cavity formation to macrocracks and creep rupture failure.
Boiler tubes usually fail due to either short or long term creep rupture failure. Short term failure is usually caused by significant overheating, usually due to steam starvation of a blocked tube. This failure is thin lip ductile failure. Long term creep failure is caused by a small amount of overheat and is thick lip brittle failure. In short term overheating insufficient or no internal working fluid causes the tube temperature to increase rapidly and remain so for some period, ranging from minutes to months, at temperatures well above the normal operating temperature. This elevated temperature leads to degradation of the microstructure and reduces the tensile properties, but most of the damage found normally results from accelerated creep deformation at the higher temperature. Under such conditions, the material is highly ductile and typically swells at the hottest area before developing a longitudinal split. The fracture appearance is of high ductility with a thin edge split, known as a ‘fish-mouth’. Long term overheating results from operation at temperatures above the normal design/operating temperature for protracted periods of time, typically several years. Temperatures are not as high as those found in short term overheating/steam starvation and the tube material does not show the same high ductility. However, creep damage still accumulates, resulting in accelerated microstructural evolution and eventually leading to
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failure. Due to the lower temperature, swelling and bulging are not evident, but the oxides adjacent to the overheated area will be thick due to the higher operating temperature, and wall thinning may result. The common appearance of creep damage in industrial plant is simple cavitation. As a creep crack forms, it acts as a stress concentrating feature generating a highly stressed area ahead of the crack tip. A relatively large strain can develop in this small area, leading to ductility exhaustion and the growth of the crack by link-up of cavities. This is the normal process for the growth of creep cracks and can result in their gradual growth through the component.
7.3.2 Fatigue Fatigue is a phenomenon that leads to failure due to variation or fluctuation in applied stress. Various types of fatigue failure have been experienced. From the perspective of fatigue life, there is high-cycle and low-cycle fatigue, and there have been many cases of failure of both types. Other types of fatigue are thermal, high-temperature, low-temperature, corrosion and impact fatigue. Fatigue failure can typically be prevented through appropriate design and/or improvement in the plant operating conditions. It can also be prevented by the use of higher fatigue resistance materials. A basic tenet of anti-fatigue design is to separate the aspects of life into crack initiation and crack propagation. Considering the conditions of usage applying to the product in question, the fatigue design life (low-cycle or high-cycle) and design criteria to be utilised must be selected. First, a stress frequency distribution for location under consideration must be obtained, and the fundamental S–N curve established. Next, considering the material, manufacturing and service conditions, the levels of influence of the various factors on crack initiation must be determined. The design curves must be calibrated and, using the fatigue damage rules, life to crack initiation estimated. Based on the results, assuming that consideration of possible crack propagation is warranted, or if cracking has already been detected, it will be necessary to integrate crack propagation life design. In such cases the applied stress and/or stress frequency must be initially determined. The stress intensity factor and the K value (together with its distribution) are obtained and, based on crack propagation rules, crack length growth curves are calculated. Life assessment is then undertaken in comparison with the time to reach the allowable crack length as separately defined. With respect to the various types of creep damage referred to earlier, design procedures for life to crack initiation and life during crack propagation are followed, after which product design and reliability assessment are carried out. Of great importance in this context is the fatigue data that are used; i.e. while employment of appropriate data is essential to life design, it must be kept in
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mind that the required data and influence factors differ with the product or object concerned. In terms of service conditions, temperature (which plays a major role with respect to fatigue behaviour and design conditions) is divided into room temperature and high temperature, and environmental conditions are determined. Separation into low- and high- cycle fatigue is also undertaken, and the types of applied stress are elucidated. Next, the applicable fatigue data and influence factors are considered with respect to the relevant service conditions. Few issues arise when existing fatigue data can be used, but when this is not the case, or when life estimation accuracy must be increased, it is necessary to obtain and prepare data by means of fatigue testing capable of replicating actual conditions.
7.3.3 Creep–fatigue interaction In machines and structures used at high temperatures, transition temperature gradients accompany startup/shutdown and fluctuations in service conditions. Repetition of such fluctuation can result in cyclic loading of thermal stress due to differences in thermal expansion, and the accumulation of fatigue damage. Furthermore, when such equipment or material is used in the creep regime, creep damage must be considered in addition to fatigue. Accordingly, the interaction between fatigue damage and creep damage is important from the standpoint of equipment/material reliability, and considerable research has been undertaken with respect to damage rules for application in design and life prediction. Investigation of creep–fatigue interaction has focused on the influence of frequency and holding time on strain control fatigue. The contribution of creep increases with lower cycles and longer holding time, with fatigue cracks and creep cavities linking in progress towards failure. On the other hand, in the case of fatigue testing with higher cycles and shorter temperature holding times, the fatigue mode governs, with failure starting in the vicinity of the surface and propagating within the grains. Much research has been performed on damage rules and life prediction with respect to the interaction of creep and fatigue, and there are four main methods used for the estimation of cumulative damage: (a) (b) (c) (d)
the the the the
linear damage summation method; frequency-modified strain range method; strain range partitioning method; ductility exhaustion method.
For the purpose of discussion, the linear damage summation method is taken as a representative approach. In this method, fatigue damage and creep damage are simply added together. Current design techniques are centred on the linear life fraction rule (which considers the life fraction to be
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7.4 Creep–fatigue design curves based on the linear damage summation method [10].
damage), and this is the basis of ASME Boiler and Pressure Vessel Code Section III Subsection NH. The linear damage summation method combines Miner’s rule, with respect to fatigue [9], and Robinson’s rule, with respect to creep, as expressed by the following formula [10]: D0 ¼ Sn=Nf þ S t=tr
½7:1
Here, n/Nf is the life fraction with respect to cyclic loading, n is the number of cycles in a given strain range, Nf is the simple fatigue life in that strain range, t/tr is the life fraction with respect to time-dependent creep, t is the time at a given stress and tr is the creep rupture life at that stress. Considering creep due to stress relaxation, division is undertaken into a number of segments where stress may be viewed as constant, with t/tr calculated for each segment, and then added together. D´ is the cumulative damage index, such that rupture is held to occur when D´ = 1. In this approach, assigning fatigue damage (n/Nf) to the horizontal axis and creep damage (t/tr) to the vertical axis, the sum of damage is represented by a straight line connecting the respective ‘1’ positions. In actuality, however, there are regions where the value of D´ is greater than 1 or less than 1, and measured values are represented using two lines. The design damage curve
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Table 7.4
Some examples of erosive wear in power generation
Component
Cause of erosive damage
Economiser tubes Turbine inlet valves and nozzles LP turbine blades
Fly ash from burning coal Spalled oxide (magnetite) from superheater Water droplets in wet steam
adopted in ASME Section III Subsection NH is presented in Fig. 7.4. [10], and it is important to remember that, in this figure, the design allowable values Nd and td (taking into account a safety factor) are substituted for Nf and tr. Other damage mechanisms experienced in power plant boilers are briefly described below.
7.3.4 Erosion Erosion is the wear produced on materials by the impact of solid particles or the impingement of liquid droplets. Erosion can result in continuous and possibly rapid metal loss and may produce smooth, scoured or rippled surfaces. As with other wear processes, erosion may combine with corrosive attack to increase the rate of damage accumulation, with the erosion process removing the passivating or protective films as they form. Some examples of erosive wear in power generation are shown in Table 7.4. Jetting damage is damage caused by a steam leak from an adjacent tube, i.e. the damage is secondary to the original leak. Jetting damage may be more widespread and is generally more costly to repair than the original failure. Once found, all local tubing should be inspected for evidence of thinned areas that could fail on return to service. Jetting damage failures should not be sent for failure analysis (although they often are). Small round, thin-edged holes should be treated with suspicion, as these are almost always jetting damage. Spalled oxide from the boiler can be carried by the steam flow into the turbine where it can cause erosion of the turbine blades and nozzle vanes. Typical damage areas are the leading edges of HP/IP blades and on the underside of blade shrouds. The onset of oxide exfoliation in the superheaters/reheaters is a cause for increased maintenance attention to inlet valves, valve seats and other components where reliable operation could be influenced by the build-up of spalled oxide.
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7.3.5 Microstructural degradation, particle coarsening and spheroidisation Figure 7.5 illustrates the step-by-step process of ageing and its effect on the microstructure of low alloy ferritic steels. Stage A is the normal microstructure of low alloy ferritic steels containing ferrite and lamellar pearlite. With time in service, spheroidisation begins and carbides precipitate on the grain boundaries (stage B). Stage C represents the intermediate stage of the spheroidisation process, and then in stage D the spheroidisation has completed. Stage E represents evenly dispersed carbides with few traces of prior ferritic/pearlitic structures and coarsening of carbide particles. Precipitation of carbides and coarsening of carbide precipitates are the main contributors to microstructural deterioration in this steel as a result of an increase in the severity of thermal ageing. Ageing adversely affects creep strength at the various stages of spheroidisation and carbide coarsening. Microstructural instability of this steel type in terms of cavity formation and cavity linkage under stress also leads to a generation of microcracks.
7.5 Effect of ageing on the microstructure of ferritic steels.
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With increasing ageing in service, a variety of secondary carbide precipitation and their coarsening with service time takes place, further reducing creep strength. This results in microstructural variation of components quite different from the initial optimised microstructure obtained in the normalised and tempered or annealed condition. Softening due to carbide coarsening thus appears to make the dominant contribution to creep, and therefore hardness change with service duration can be used as an important parameter for component integrity evaluation. Advanced ferritic/martensitic steels do not undergo the same spheroidisation and carbide coarsening processes that are observed during the ageing of low alloy ferritic and bainitic steels. Some particle coarsening does occur but on a much finer scale in the martensitic steels. The M23C6 carbides coarsen slowly during service but the very fine MX carbonitride creep-strengthening precipitates are highly resistant to coarsening. Other phases, such as Laves and Z-phase, appear and/or coarsen in some of the martensitic steel grades during long term high-temperature service. This is discussed later in this chapter.
7.3.6 Steam oxidation Oxidation due to superheated steam is a serious problem in boilers with high steam temperatures, which have become increasingly common in recent years. Damage resulting from steam oxidation includes tube blockage caused by the exfoliation and build-up of scale. This leads to creep damage, induced by overheating of affected tubes, as well as particle erosion in turbine side equipment downstream from the boiler due to dispersion of scale. The formation and exfoliation behaviour of steam oxidation scale is quite different for ferritic and austenitic steels. In the past, because the damage resulting from steam oxidation was not as severe in cases involving ferritic steels as compared to austenitic steels, ferritic-related research was relatively sparse. Recently, however, considerable research has been conducted on high-strength ferritic steels. As the difference in the coefficient of thermal expansion between ferritic steels and their scale is less than with austenitic steels, scale does not reach the exfoliation limit thickness until a considerable time has elapsed. Nevertheless, scale exfoliation has been confirmed for CrMo steels after long term usage of over 100 000 hours. Scale on CrMo steels is known to have a dual layer structure, with exfoliation occurring when the thickness reaches approximately 500 μm. CrMo scale is different from that affecting austenitic steels, in that all layers exfoliate simultaneously and the scale is hard.
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Inspection and monitoring of damage and integrity/life assessment issues in high chromium martensitic steels
7.4.1 Issues concerning the use of ASME P91 and P92 type high Cr martensitic steels The 9 %Cr martensitic steel P91 has now been in use in the power generation industry for over 20 years. Over this time there has been a number of incidents of cracking and failure in components made from this steel. Thick section component failures have generally been caused by Type IV cracking at welds (Fig. 7.6), while thin section components have failed due to higher than expected levels of steam oxidation, (Fig. 7.7) and an associated increase in the metal operating temperature. Many of the Type IV failures have been associated with materials that are relatively weak as a result of possible chemical composition effects (Al:N ratio) and/or incorrect heat treatment (e.g. over-tempering). Consequently, some plant operators (such as RWE npower in the UK) have focused their inspection activities on the P91 components that exhibit relatively low hardness values and, typically, this means that components with hardness of less than 200HV are kept under close observation. Although the 9–12 %Cr martensitic steels are expected to be less vulnerable to thermo mechanical fatigue damage (as a result of reduced component thickness and therefore lower through thickness temperature gradients), there is some evidence that the effect of creep–fatigue interaction could be more severe for the P9 weldments as a result of the low ductility associated with Type IV failure in this material [1]. Consequently, for cycling plant, there is a need to observe closely the casts with a tendency to show a larger drop in creep ductility with service duration. As a result of the increasing number of plant failures, interest in integrity/ life assessment and monitoring has become acute. This is especially so with regard to creep damage development because the traditional NDE methods of metallographic replication and early damage detection in these steels have been found to be less than satisfactory. Therefore, there is a need to study, develop and establish new methodologies and techniques for life assessment of the 9–12 % Cr martensitic steels. These aspects are discussed below. Although thick section P91 was first introduced in power plants just over 20 years ago, non-destructive integrity and creep life assessment technology has not yet been established for this steel, particularly for welded joints that are susceptible to Type IV failure. Indeed, most of the failures to date in P91 thick section components are known to have been Type IV failures. Now that the use of this steel is reaching the mid-life stage, there is a strong demand from power plant operators for the development of non-destructive
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7.6
Example of a ‘Type IV’ cracking in the fine grain region of P91 HAZ.
7.7
Cracking and exfoliation of internal oxide scale.
damage detection and life assessment technology for welded structures. One problem with Type IV failure in the welded high-temperature steel structures has been that the Type IV cracking usually initiates sub-surface and cannot be detected by more economical and time-saving procedures such as magnetic-particle testing (MT), metallographic replication, etc. The cracking visible by replication, etc., only emerges at the surfaces near the last leg of its journey, thus making the structures potentially unsafe. P91 steel
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has further problems in that, unlike the low alloy ferritic steels operating in the creep regime, the more easily visible changes such as spheroidisation and breakdown of the microstructure at the scale that can be seen under an optical or a scanning electron microscope does not occur. This then means that the costly transmission electron microscopy (TEM) has to be used to identify changes in the microstructure. Even then the changes may not be so obvious to a non-specialist, as it is mainly the changes in the precipitates and the relative dislocation density that adversely affect the life of these steels. Thus there has been a need to develop new tools and more sensitive methods for integrity and damage assessment in these steels. Recent systematic research and study, particularly in Europe and Japan, of the inspection and monitoring of component microstructural damage, cavitation, cracking and failure are giving indications of potential new inspection, monitoring and integrity assessment techniques [11] that may be used successfully on 9–12 % Cr steels. Thus, for example, the UK based European Technology Development (ETD), together with its industrial collaborators from Europe, Japan and North America, has been studying the development of tools and methodologies for early stage damage detection and life prediction as a part of its international multiclient project P91 Integrity [12]. The recent NDE developments are discussed in the following sections.
7.4.2. Microstructure based integrity assessment As stated earlier, the observation of microstructural damage is widely used in inspection, condition monitoring and support of life management in hightemperature low alloy ferritic steels. The damage evolution is materialdependent and requires confirmation from inspection data. For most low alloy steels, compilation of inspection data has been carried out to establish guidelines for this purpose. Useful experience of in-service damage is less easily available for the relatively newer P91 steel. It is well known that the microstructure and precipitation in the 9–12 %Cr martensitic steels are influenced by the chemical composition details and heat treatment parameters. For the determination of microstructural parameters specific methods are required. For monitoring and component assessment a detailed knowledge of the interaction between the load parameters and microstructure is necessary. The high creep resistance of these steels is based on their martensitic microstructure, carbide and nitride particles size and density, dislocation (sub-grain) structure and the Laves phase, which has a detrimental effect in the medium to long term. The appearance of the in-service Z-phase may indicate a loss in rupture strength [13]. In 9–12 %Cr steels the martensite needles are decorated with precipitates. In long term service the precipitates
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TEM extraction replica
TEM metal foil
Table 7.5 Methods for identification and quantitative description of microstructural parameters [14] Microstructural parameter
Method
Sub-grain size Mag. 25 000 :1 (3 x 3 micrographs ≈ 50 m2)
TEM Intersections of sub-grain boundaries with test lines at ± 458 to sub-grain elongation are counted TEM Intersections of dislocations at foil surfaces are counted
Dislocation density inside the sub-grains Mag. 100 000 :1 (10 micrographs = 4 m2) Mag. 50 000 :1 (5 micrographs = 8 m2) Characterisation of precipitates according to: . Phase type (classification) – Morpholgy and internal structure – Composition – Crystal structure .
Energy filtered TEM (EFTEM) TEM EDS (100 to 300 particles) Selected area diffraction (SAD) Image processing (IA)
Particle size, size distribution and number
at the needle boundaries grow and thus an incipient dissolution of the needle shaped microstructure can be observed under an electron microscope but not under an optical microscope, which is therefore not a suitable tool for evaluation of the material life exhaustion. In high-temperature service additional precipitation occurs in 9Cr steels. The addition of W in P92 inhibits the precipitation of the intermetallic Laves phase Fe2(Mo,W). The Z-phase, Cr(Nb,V)N, precipitates at service temperatures consuming MX particles, which give creep strength to these steels at high temperature and are stable over long service durations. The Zphase precipitates as large particles that do not contribute to the creep strength. The time–temperature regions in which Z-phase precipitation may occur in high Cr martensitic steels depend on chemical composition and the as-received microstructure after heat treatment [13]. For qualitative evaluation of the microstructure with respect to precipitation, dislocation density and sub-grain formation, the use of transmission electron microscopy and some specific methods to relate the changes in microstructure to life consumed will be necessary. This may involve the use of statistical analysis. As an example, the microstructural parameters and the methods of their quantification used in Germany are listed in Table 7.5. The sub-grain size and dislocation density can be determined by conventional TEM. For this, metal foils need to be prepared
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from the specimens by electropolishing and ion milling. The classification and quantification of precipitates can be carried out on extraction replicas using a qualified combination of energy filtered TEM (EFTEM), energy dispersive X-ray spectrometry (EDS), electron diffraction and image analysis applied at the same visual field [14].
7.4.3 Optical and scanning electron microscopy for cavitation measurement Work on this aspect has been discouraging as the microstructural changes do not appear to be significant enough to be noticed under an optical or even scanning electron microscope [2].
7.4.4 Atomic/scanning force microscopy for on-site cavitation damage assessment The conventional laboratory version of the atomic or scanning force microscope (SFM) has been generally used on biological small samples for laboratory studies. However, a portable version of this is now available and has been modified and used for on-site high-temperature plant creep damage assessment [12]. This portable version was used for studies on a P91 creep damaged pipe (Fig. 7.8) and heavily oxidised T91 tubes containing known levels of creep cavitation. Both studies showed that creep cavities only a few nanometres in size can be resolved by this instrument. A further advantage of using this instrument is that it can also measure the cavity depth and can
7.8 Portable SFM being used on a creep damaged P91 welded pipe.
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7.9 SFM images acquired before and after chemical etching for 2 minutes (scale on right represents depth). Top left: unetched sample showing cavities and path of the line scan. Top right: line scan showing cavity depth. Bottom left: etched sample showing cavities, etched surface and path of the line scan. Bottom right: line scan showing cavity depth and profile of the etched surface along the line scan.
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easily calculate cavitation damage in volumetric terms – a better measure of damage due to creep. The required sample preparation is the same as that for conventional metallographic replication, i.e. a finish of 1 μm. Indeed, etching of the sample is not always necessary, as the results of unetched samples were also reasonably satisfactory, although the resolution improved with etching (Fig. 7.9). Fig. 7.9 shows how clear images and analysis of cavitation damage can be obtained using the portable SFM. Various experiments have been carried out using this tool (known as ‘DASFM’, or damage assessment-SFM) [12]. As the performance of SFM can be sensitive to vibration, for the tube and pipe specimens the study included artificially introduced vibration to simulate the environment of an industrial plant. The results have been very encouraging. Indeed, the effect on thick section components, which will be less affected by vibrations in the neighbourhood, was very slight. A software especially developed for this purpose has been used to refine the image further. The system has now been tried in industrial plants on creep damage pressure vessels and turbine rotors with excellent results [12]. A special advantage of this tool is that DA-SFM is a mechanical device and can be used on all types of materials (metallic and non-metallic) and for detection and inspection of all types of damage (creep, fatigue, corrosion, etc.). The fact that the depth of cavities can also be studied (to a few nanometres) and their volume calculated means that DASFM can provide a better tool for more accurate life assessment. In the case of P91 type steels, it means that smaller cavities can be detected at a much earlier stage in life than has been possible so far. As the cavitation or other damage images can be recorded on a connected laptop computer, this means that unlike the replicas, which are usually taken to laboratory and studied later, the results from DA-SFM are available instantly for evaluation. Indeed, it is possible that the use of SFM in situ may render the making of replicas for assessing in-service creep damage a thing of the past, thus saving a great deal of cost and time to industry.
7.4.5 Ultrasonic noise method for damage monitoring Recently this technique was used by IHI in Japan [2] to monitor damage in P91 pressure vessel creep tests. In this method, the ultrasonic sensor was moved in a perpendicular direction across the seam weld to measure the creep damage distribution. The principle lies in the fact that when a steel plate is irradiated on the surface by ultrasonic waves, an echo is generated from the back surface of the plate. ‘Noise’ signals, observed between the incident pulse wave and the echo, are generated by aggregates of creep voids and microcracks. Thus, the ultrasonic ‘noise’ method can measure the damage under the surface by measuring the noise produced.
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7.4.6 Detection of creep damage by ultrasonic velocity change method In a recent investigation three ex-service T91 tubes were examined for creep damage using more sensitive techniques than the traditional ultrasonic (UT) methods [12]. The three samples are shown in Fig. 7.10. One of these tubes (tube A) had little or no creep damage and the other two had varying degrees of damage (tube C showing the most cavitation). All three tubes had a similar nominal wall thickness of about 5 mm. The literature suggests that creep damage causes reduced ultrasonic propagation speed. The results of UT tests are shown in Fig. 7.11.
7.4.7 Potential drop as a damage monitoring technique It is well known that the electrical potential drop (PD) method can be used to monitor cracks in components. Indeed, it is now a well-established practice in creep crack growth testing where either the direct current (DC) or alternating current (AC) potential drop method is used to monitor crack initiation and crack growth as a function of time. This has been done both for standard laboratory specimens and for large pressure vessel tests. In these tests, normally the input current to the test specimen is kept constant, using a constant current source, and output voltage over the region being monitored is measured continuously. In a cracked specimen or component, as the crack grows the path travelled by the output current increases, thus showing an increase in the measured PD across the crack. Similarly, as the microstructure changes, or creep cavitation develops, the resistance to the current flowing through a specimen or a component changes, resulting in an increase in the PD output. The output PD is usually normalised against a reference output PD across a reference point on the same or a similar testpiece through which the same current is flowing, so that any small changes in the current input due to a change in the specimen temperature or other effects can be normalised. The DC input current has the advantage that it flows deep into the specimen and therefore can give an average value of the damage through the wall thickness of a component. The AC current is more restricted to flowing along the specimen surface but has the advantage that the output is more sensitive to damage or cracking in the specimen and the source is more easily available. For P91 components, some exploratory PD tests have been carried out in Japan using AC and the results due to Masuyama [15] are shown in Fig. 7.12. They have shown a good correlation between creep life and the PD measurement. This technique has thus the potential to be used for future P91 and other materials integrity monitoring and life assessment.
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7.10 Tube specimens A, B and C.
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7.11 UT velocity change results. Propagation speed measurements for tube specimens A, B and C. The thick and thin lines in the centre show the respective mean values and standard deviations of the measurements along the lengths of the tubes. The regions of high outer surface scale and creep damage are indicated by the thick semi-circular lines. © Woodhead Publishing Limited, 2011
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7.12 Potential drop ratio versus the creep life fraction relationship [15].
7.4.8 Hardness monitoring as an integrity assessment tool Masuyama has recently demonstrated a method for creep life assessment of P91 welded components based on hardness measurements [16]. In this study, P91 base metal creep specimens were tested at five different stress/ temperature conditions and a large P91 cross-weld creep specimen was tested at a single temperature/stress condition. At each stress/temperature condition, base metal tests were terminated after different creep exposure times and hardness tests were performed on the testpiece gauge length (creep deformed) and on the grip end (no creep deformation). The cross-weld specimen test was interrupted periodically so that hardness measurements could be performed on the weldment and base metal regions of the testpiece gauge length. Masuyama showed that the hardness measurement based on the minimum value in the HAZ is a candidate non-destructive creep life assessment method for P91 welded joints.
7.5
Sources of further information and advice
Shibli I A and Robertson D G, ‘Review of the use of new high strength steels in conventional and HRSG boilers – R&D and plant experience’, ETD Report 1045-gsp-40, September 2006,
[email protected]. Nonaka I et al. ‘Full size internal pressure creep test for welded P91 hot reheat piping’, Special Issue, Int. J. Pressure Vessels and Piping, 84, Issues 1–2, January–February 2007, 88–96. Auerkari P et al. ‘Creep damage and expected creep life for welded 9–11 % Cr steels’,
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Special Issue, Int. J. Pressure Vessels and Piping, 84, Issues 1–2, JanuaryFebruary 2007, 69–74. Bendick W, Fuchs R, Hahn B, Heuser H, Jochum C and Vaillant J C and Weber J, ‘Application/capability and welding of modern heat resistant steels (T/P91; T/ P23; T/P24) for the maintenance and refurbishment of power station components’, in 2003 EPRI International Maintenance Conference, Chicago, 2003. Ennis P J, ‘Steam-side oxidation’, ETD Training Course on P/T91, London, September 2006,
[email protected]. Fleming A, Maskell R V, Buchanan L W and Wilson T, ‘Materials development for supercritical boiler and pipework’, in Proceedings of Conference in Materials for High Temperature Power Generation and Process Plant Applications, Materials Congress 1998, IOM Publications, London. Quadakkers W J and Ennis P J, ‘The oxidation behaviour of ferritic and austenitic steels in simulated power plant service environments’, in Materials for Advanced Power Engineering 1998, Proceedings of the 6th Liege Conference, Liege, Belgium, Vol. 5, Part 1, p. 123. Viswanathan R et al. ‘Boiler materials for ultra supercritical coal power plants – steam side oxidation’, J. Mater. Engng and Performance, 15 (3), June 2006. Matsumoto H, Nishimura N, Iwamoto K, Tominaga K and Kuroishi T, ‘MHI’s activity for life management of aged power boilers’, in Proceedings of Conference on Maintenance and Overhaul Management, 2–3 October. 2003, London, European Technology Development,
[email protected]. Nishimura N, Masuyama F, Nakatani H and Inada M, ‘Development of creep life monitoring method using alternating current potential drop method for fossilfuel power boilers’, Mitsubishi Juko Giho, 30, 1993, pp. 342–346 (translated from Japanese). Masuyama F, ‘Integrity and life assessment of P91 components’, in Proceedings of International Seminar on Industry and Research Experience in the Use of P/T91 in HRSGs/Boilers, 7–8 December, 2005, London, European Technology Development,
[email protected]. Masuyama F, Fujita M, Maruyama M and Endo M, ‘Development of creep life assessment technology for mod. 9Cr–1Mo steel base metal and weldments’. Shibli I A and Holdsworth R, ‘Creep and fracture in high temperature components – design and life assessment issues’, in ECCC Conference Proceedings, April 2009, DESTech Publishers, Lancaster, Pennsylvania, USA. Shibli I A, Holdsworth R, and Merckling G, ‘Creep and fracture in high temperature components – design and life assessment issues’, in ECCC Conference Proceedings, September. 2005, DESTech Publishers, Lancaster, Pennsylvania, USA.
7.6 1.
2.
References Shibli A., ‘Performance of P91 thick section welds under steady and cyclic loading conditions: power plant and research experience’, OMMI, www.ommi. co.uk Vol1, Issue 3, December 2002. Nonaka I, et al. ‘Full size internal pressure creep test for welded P91 hot reheat
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3.
4. 5.
6.
7.
8.
9. 10. 11. 12. 13.
14.
15. 16.
303
piping’, Special Issue, Int. J. Pressure Vessels and Piping, 84, Issues 1–2, January–February 2007, 88–96. Orr J, Beckitt F R and Fawkes G D, ‘The physical metallurgy of chromium molybdenum steels for fast reactor boilers’, in Proceedings of BNES Conference on Ferritic Steels for Fast Reactor Steam Generators, London, 1977. Murray J D, ‘Precipitation processes in steel’, ISI Special Report 64, 1959, pp. 285. Kushima H, Watanabe T, Murata M, Kamihira K, Tanaka H and Kimura K, ‘Metallographic atlas of 2.25Cr–1Mo steels and degradation due to long term service at elevated temperatures’, in Proceedings of ECCC Creep Conference, Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, 12–14 September, 2005, London, pp. 223. Jakobova A, Vodarek V, Hennhofer K and Foldyna V, ‘Microstructure and creep properties of P91 steel and weldments’, in Materials for Advanced Power Engineering 1998, Proceedings of the 6th Liege Conference, Liege, Belgium Vol. 5, Part 1, pp. 373–382. Kubon Z, Foldyna V and Vodarek V, ‘Optimised chemical composition of 9– 12 %Cr steels with respect to maximum creep resistance’, in Materials for Advanced Power Engineering 1998, Proceedings of the 6th Liege Conference, Liege, Belgium, Vol. 5, Part 1, p. 375. Wachter O and Ennis P J, ‘Investigation of the properties of the 9 %Cr steel of the type 9Cr–0.5Mo–1.8W–V–Nb with respect to its application as a pipework and boiler steel operating at elevated temperatures’, PhD Thesis, Research Centre Julich, Germany, March 1995. Masuyama F, in e-Lifing–ETD’s Compendium of Lifing Procedure for Power Plant and HRSGs, 2010,
[email protected]. ASME Boiler and Pressure Vessel Code, Section III, Subsection NH, 2007. Brear J M ‘Generic guidelines for component life assessment’; SPRINT SP 249 Project, ERA, December 1994, pp. 46. ETD project ‘Integrity of P91 welded components’, Project 1088-gsp-prop06. Cipola L et al., ‘Formation of Z phase in a 12 %CrVNbN model steel’, in Shibli I A and Holdsworth S R (eds). ECCC Conference. Proceedings on Creep and Fracture in High Temperature. Components – Design and Life Assessment Issues, April 2009, by DEStech Publishers, Lancaster, Pennsylvania, USA. Maile K, ‘Evaluation of microstructural parameters in 9–12 % Cr–steels’, Special Issue, Int. J. Pressure Vessels and Piping, 84, Issues 1–2, January– February 2007, 62–68. Masuyama F, ‘Integrity and life assessment of P91 Components’, ETD Training Course on P/T91, London, September 2006,
[email protected]. Masuyama F, ‘Integrity and life assessment of P91 components’, in Proceedings of International Seminar on Industry and Research Experience in the Use of P/ T91 in HRSGs/Boilers, 7–8 December 2005, London, European Technology Development,
[email protected].
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8 Creep, fatigue and microstructural degradation in gas turbine superalloys P . A U E R K A R I , VTT Technical Research Centre of Finland, Finland
Abstract: Creep and fatigue are important and life-limiting damage mechanisms for many gas turbine components and therefore must be accounted for in design. Creep is promoted by high levels of stress and temperature that also drive microstructural changes in turbine materials. Fatigue is caused by fluctuating loads from, for example, rotation, vibration or thermal cycling, and can also limit component design and life at more modest temperatures. The main characteristics of creep, fatigue, microstructural damage and their combined impact are introduced below, with emphasis on aspects relevant in life assessment of high temperature components such as turbine blades made of superalloys. Key words: creep, fatigue, microstructure, degradation.
8.1
Introduction
The expected life-limiting damage mechanisms such as creep and fatigue must be accounted for in the design of many gas turbine components that will operate at high temperatures. The rate of creep, or viscous timedependent deformation, is accelerated by increasing stress and temperature, and will set limits to stress and time to maximum allowable strain (or fracture) in components like turbine blades, vanes and other stressed parts of the turbine hot path. Fatigue, or damage by fluctuating (cyclic) stress, will progress with the loading cycles and for pure fatigue is considered to be time independent. Fatigue can also limit component design and life at low or modestly elevated temperatures, e. g. in rotating or vibrating parts such as shafts, discs and compressor blades. In components that operate under conditions where both creep and fatigue damage can be expected, combined impact needs to be considered in design and life assessment. At high service temperatures the gas turbine materials and particularly the 307 © Woodhead Publishing Limited, 2011
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nickel and cobalt based superalloys of the hottest parts are not stable but experience microstructural degradation that will reduce the material performance from that of the initial state. Creep, fatigue and microstructural degradation will proceed faster with increasing temperature and stress, but the rates of these mechanisms also depend on the material and its initial state. As the resulting damage grows towards the end of the component life, the extent of observed damage and degradation can also serve as useful indicators of the in-service condition, residual life or maintenance limits.
8.2
Creep
At sufficiently high temperatures, but well below the melting range of solid materials, technically significant viscous (time-dependent, liquid-like) flow, or creep, takes place in the direction of stress. This happens in, for example, gas turbine blades, vanes, combustors and other mechanically loaded parts of the turbine hot path. Because increasing temperature and pressure levels improve the process performance of gas turbines but also make creep faster and shorten creep life, creep strength of the structural materials is one limiting factor in the design of the hot parts of gas turbines. Creep of most solids when applying a constant tensile stress or load takes place in typical successive stages that are conventionally described as shown in Fig. 8.1. Similarly to loading under conditions without significant creep, after an initial elastic (recoverable) strain the material will show irrecoverable plastic strain that is fast enough to follow the rate of loading. Thereafter, the first or primary stage of creep shows hardening or a gradually decreasing strain rate with continuing loading (and strain). The hardening mechanisms are competing with simultaneous mechanisms of weakening and recovery that gradually become more prominent while the relative effect of hardening is waning, and so the overall creep (strain) rate will pass through a minimum. The region of minimum creep rate is also often called the secondary stage of creep, for within this region the strain rate is low and may appear practically constant in the creep curve (strain– time curve), as shown in Fig. 8.1. After passing the minimum, the weakening mechanisms increasingly dominate in the third or tertiary stage of creep, where the strain rate is progressively increasing until failure. The shape of the creep curve depends on the material, but generally tends to show more pronounced primary creep when the stress is high and the temperature level modest, and is more dominated by tertiary creep when the opposite is true. In comparison to creep under tensile loading, the tertiary stage of creep is suppressed under compression. This is partly due to true suppression of weakening mechanisms like small scale cracking in compression and partly only apparent due to the increasing/decreasing cross-section of uniaxial specimens under compressive/tensile constant load testing. Although
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8.1 Typical shape and trends of the creep curve. Note that the secondary stage of an apparently constant strain rate may actually correspond to a minimum strain rate between attenuating primary and accelerating tertiary stages.
compressive loading is generally less damaging in creep, it is important to note that intermittent compression increases the overall local strain range when tensile loading is subsequently applied, and this promotes damage in both creep and fatigue (see below). The importance of creep strength is reflected in the development of structural materials, since sufficient creep strength is a prerequisite for improving the efficiency, fuel economy and specific power of gas turbines. Hence the historical development of the materials for the hot parts of gas turbines has consistently been towards increasing creep strength. Creep strength is also a major motive for selecting the chemical composition and the preferred microstructures of these materials for survival at ever higher operating temperatures. For example, the shift from polycrystalline to directionally solidified (DS) alloys in turbine blades removes the creep strength-limiting grain boundaries from the most heavily loaded directions transverse to the blade axis, and further improvement has been obtained by using single crystal (SC or SX) alloys that include no grain boundaries (Fig. 8.2). An important parallel contribution to creep strength has been obtained by alloying for increasing volume fraction of the creep resistant intermetallic gamma prime phase (γ´, up to 60–70 % vol) and for its stability at high service temperatures. High creep strength alone is not enough for satisfactory creep
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8.2 Development of the temperature capability (at 138 MPa/100 000 h creep life) of an Ni alloy for gas turbine blades as equiaxed (EA), directionally solidified (DS) and single crystal (SX) versions; adapted from reference [1].
performance. Sufficient creep ductility, or strain to failure, is also needed against local deviations in manufacturing, local or temporary overloads and local in-service damage. Unsatisfactory ductility or embrittlement may be caused by, for example, chemical impurities in the material or by unfavourable routes of fabrication in casting, forming, joining or heat treatment. Ductility can also be reduced by microstructural changes like precipitation and growth of embrittling phases during high temperature service. Low ductility can shorten creep life by promoting localisation of damage and cracking in the material. Creep strength is mostly determined from uniaxial specimens subjected to constant load, and is recorded as the time required to given levels of strain and/or to rupture at constant temperature and initial stress. The difference in stress between constant load and constant stress tests is small at low strains (up to about 0.1 to 1 %), which is often sufficient for design. In addition, constant load tests are cheaper and result in somewhat conservative (shorter) life for the same initial stress. The minimum creep rate at (moderate) stress σ and temperature T can be expressed as Q : e ¼ A exp ½8:1 sn kT
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where A is approximately constant for a limited range of temperatures, Q is the apparent activation energy of creep, n is the stress exponent of creep (usually n ≥ 3) and k is the Boltzmann constant. Assuming an approximately constant strain to failure B, the time to rupture (tr)to can be approximated as B B Q n tr ¼ ¼ exp ½8:2a s kT e A Taking logarithms of both sides and rearranging gives a simple expression for creep life: log tr ¼ PðsÞ þ CQ=T
½8:2b
tr ¼ 10PðsÞþCQ=T
½8:2c
or
where P(σ) includes the stress dependence often as a polynomial of stress or log (stress), and C is a constant. A number of alternative expressions can be derived for the same purpose, and it is customary to test multiple expressions for the best fit to the available data. For example, for IN738LC, a cast polycrystalline Ni alloy for gas turbine blades (buckets), an approximate expression predicting the mean time to creep rupture can be given as [2] tR ¼ 10
PðsÞ T C
½8:3a
where P(σ) = 42785.7 – 7142.86 log(σ) and C = 20. A corresponding mean creep life for the cast version of another polycrystalline blade alloy Udimet 500 can be expressed as [3] tR ¼ 10PðsÞ BT
½8:3b
where P(σ) = 45.01799 – 24.85729 log σ + 11.48121(log σ)2 – 2.094012 (log σ)3, and B = 0.01798595. For comparison, the mean creep life of the wrought polycrystalline blade alloy Udimet 720 can be given as [4] tR ¼ 10
PðsÞ61000 T
C
½8:3c
where P(σ) = 30.104990.04966 σ0.5log σ and C = 20. For the cast blade alloy GTD-111 the mean time to creep rupture can be expressed as [2] tR ¼ 10
PðsÞA T
C
½8:3d
where P(σ) = 0.0536 + 0.00393 (log σ) + 0.00347(log σ)2 – 0.00463 (log σ)3
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8.3 Creep rupture strength as predicted by expression [8.3a] for cast alloy IN738LC.
+ 0.000843 (log σ)4, A = 555556 and C = 20.125. Similar expressions that naturally differ numerically can be applied for creep life, when, for example, 1 % creep strain is used as a life limit instead of creep rupture. The predicted mean creep rupture strength as a function of time is shown for these blade alloys at selected temperatures in Figs 8.3 to 8.6. As noted above, creep is not limited to any particular pattern of loading or temperature history, and when both stress and temperature vary the material response can be quite complex. However, a few simplifying principles can often be used so that the usual creep and creep rupture data obtained in constant load testing remain useful for predicting creep life under more complex loadings. First, as the creep rate is strongly dependent on temperature and stress, the ranges for both can be truncated to regimes of practical engineering interest. Second, a common simplification of the linear life fraction rule (Robinsons, rule) states that the life limit is attained when X ti ¼D ½8:4 tri i where ti is the length of the time increment i spent at stress σi and temperature Ti, tri is the time to rupture (or to limiting strain) under the same loading conditions and D ≤ 1 is the limit level of life fraction that can be tolerated. The values of tr can be taken from, for example, the expressions [8.3] above,
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8.4 Creep rupture strength as predicted by expression [8.3b] for cast alloy Udimet 500.
8.5 Creep rupture strength as predicted by expression [8.3c] for wrought alloy Udimet 720.
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8.6 Creep rupture strength as predicted by expression [8.3d] for cast blade alloy GTD-111.
if available for the material. This rule is not thought to be generally valid for any complex cases, but is widely applied when loading remains proportional and tensile, e. g. when considering the relative effects of throttle position on the life of turbine blades, as periods of higher power will shorten creep life by the amount corresponding to the increased stress and temperature levels. Another example of a special loading pattern is tension with fixed total strain, so that when elastic strain is converted to permanent or plastic strain by creep, the accompanying (elastic) stress decreases in this process called relaxation. This is common in, for example, bolts of high temperature flange joints, but relaxation can be a part of the local stress history of many gas turbine components. For example, thermal stress peaks on turbine blade surfaces from startup and tripping will be reduced by creep relaxation under subsequent steady operation [5].
8.3
Fatigue
Fatigue is a common damage mechanism that is characterised by material weakening, crack initiation and growth caused by fluctuating or cyclic stress. Fatigue can occur at any temperature. Fatigue strength decreases with the number of comparable loading cycles and is usually considered
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separately for low cycle fatigue (LCF), up to about 103–104 cycles and high cycle fatigue (HCF). Although fatigue strength is characteristic for a material, fatigue life is also strongly reduced by stress raising local features of component geometry or defects, to such an extent that particularly HCF performance of a component is relatively less influenced by the material or temperature when compared to creep. Fatigue strength can be measured from uniaxial cylindrical specimens subjected to defined loading (stress or strain) cycles until failure or from precracked specimens to evaluate crack growth rates as the extent of crack growth per load cycle. For testing fatigue strength, the cycles can be under load or strain control depending on the component considered. Particularly important features of the loading cycles are minimum and maximum levels of stress or strain that typically define mean stress, stress (and strain) amplitude and the stress ratio (R = minimum/maximum stress). For locations where stresses and strains in the load cycles remain low, the HCF properties may be sufficient for design. However, for many critical components the life-determining loading cycles need to be quite extensive for competitive performance, and therefore the component life is more limited by the LCF. One convenient and simple way of describing the fatigue performance of a (nominally defect-free) material is a plot of the experimental (constant) stress or strain range, or amplitude, as a function of the number of cycles to failure. For such a plot to include both LCF and HCF regions and the effect of mean stress σ0, this can be modelled, for example, for the strain range Δε as De sf s0 ¼ ð2Nf Þb þef ð2Nf Þc 2 E
½8:5
where σf and εf are the fatigue strength and ductility ‘coefficients’ (close to the true fracture strength and strain in monotonic tension), respectively, and b and c are the corresponding fatigue strength and ductility exponents (fractional negative constants). Note that under strain-controlled cycling, the initial mean stress σ0 may relax so that (σf σ0)/E in the expression [8.5] is not constant. Examples of combined HCF and LCF performance of Ni-alloy IN738LC for blades and Co-alloy FSX-414 for vanes are shown in Fig. 8.7; note the relatively higher short term ductility and lower HCF strength of the cobalt alloy. Here the corresponding numerical expressions for the total strain range (%) at failure of IN738LC (850 8C) are De ¼ 0:544ð2Nf Þ0:045 þ8:7ð2Nf Þ0:655
½8:6a
De ¼ 0:7ð2Nf Þ0:044 þ19ð2Nf Þ0:575
½8:6b
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8.7 Constant amplitude fatigue performance of alloys IN738LC and FSX-414 at 850 8C. Data from reference [2] have been fitted to equation [8.5].
for the lower and upper bound curves, respectively, and for FSX-414 at 850 8C De ¼ 2:54ð2Nf Þ0:234 þ116ð2Nf Þ0:903
½8:7
These curves refer to results from constant strain amplitude testing of uniaxial specimens. At typical critical locations that will fail first in actual components, the loading patterns can be much more complex than those in the constant amplitude tests. To account for the difference, the loading cycles of the fatigue tests should preferably be made to mimic the actual service cycles. However, often the service cycles are interpreted with simplified methods such as rainflow counting, so that the service history will correspond better to a collection of cycles in the simpler constant amplitude testing. Although testing may also use specimens that include small defects (Fig. 8.8), the common S–N (or Wo¨hler) curves to characterise fatigue are generally applied for materials assumed to be free of defects, and the initial part of the fatigue life is taken to be spent in crack initiation. Increasing the temperature will reduce fatigue life, although the effect is not as strong in fatigue as in creep (Fig. 8.9). In many applications fatigue damage arises from the combined effect of
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8.8 HCF fatigue strength for IN738LC and IN939 at 850 8C. Data from reference [2].
8.9 Temperature dependence of the LCF performance for alloy IN738LC, showing the strain range Δε as a function of the number of cycles NA [6].
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8.10 Fatigue crack growth rate (approximate range) of IN738LC at 850 8C (R = 0.1). Adapted from references [7] and [8].
repeatedly changing the stress and service environment. The impact of the environment may also depend on the rate of reactive elements arriving to contact the critically stressed locations like crack tip regions. In addition, the flow of air and flue gas and the combustion process may result in significant vibration, mainly contributing to high cycle fatigue. The highest tensile stress of a turbine blade typically occurs when the turbine trips at a high power level, resulting in a sudden cooler air flow on hot blades and high tensile stresses on the surfaces. However, a significant cycle is also produced by a normal startup, operation and shutdown with combined thermal, rotational, bending and torsion stresses. Any preexisting casting pores, inclusions or other defects will help to initiate fatigue cracks that grow initially slowly or even at a retarding pace, so that many concurrent cracks appear before one of them will start to dominate. An alternative approach to fatigue of nominally defect-free material assumes existing defects to consider the crack growth rates in fatigue. As long as such defects are small, they grow slowly or even at a decreasing rate with increasing size. As a result, a period of no apparent growth (incubation or crack initiation) will dominate until a sufficiently long crack is finally present to show accelerating growth until failure. Complexities in loading may further include, for example, non-proportional thermal cycling or other
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special features, but to assess fatigue life, a common simplification is the linear life fraction rule (Miner’s rule) assuming life limit to be attained when X Nj ¼ Df Nrj j
½8:8
where Nj is the number of cycles with loading conditions (stress/strain cycle and temperature) j, Nrj is the number of cycles to failure under the same loading conditions and Df≤1 is the limit level of the life fraction that can be tolerated. The values of Nr can be taken from, for example, the expressions [8.5] to [8.7] above, if available for the material. In the case of a growing fatigue crack in constant amplitude cycling, the crack growth rate is often described as [7] da ¼ CDKm dN
½8:9
where C is a constant for a given material and temperature and ΔK is the range of the stress intensity factor for the crack above a threshold level ΔKth but below a critical upper level leading to very fast growth (Fig. 8.10). For correction according to the levels of the mean stress or stress ratio R, see reference [7].
8.4
Combined creep and fatigue
Although completing any loading cycle will take some time, pure fatigue is generally assumed as time-independent so that the growth rate of the resulting fatigue damage is independent of the loading frequency. At high temperatures with increasing time spent in the cycle, the damage growth rate will start to increase with decreasing frequency if the frequency is low enough. This implies influence by truly time-dependent damage mechanisms such as creep, oxidation or hot corrosion. When the effect is by creep, the combined effects of creep and fatigue can also be important in the intermediate frequency range. With further decreasing frequency, finally the overall rate will become dominated by the time-dependent mechanism, or creep in this case. The real local loading cycles of gas turbine components can be much more complex, but the same principles will apply so that the more time is spent, especially at high load level parts of the cycle, the more creep-dominated will be the combined damage. If the damage is measured as crack growth, the combined total growth rate per cycle (Fig. 8.11) can be expressed as [7] da da 1 ¼ CDKm þ dN dt f
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8.11 Effect of frequency on the crack growth rate of nickel alloy AP1 at 700 8C and constant Kmax=33.3 MPa√m. Adapted from reference [7].
where da/dt refers to growth rate due to creep and f is the loading frequency. Even in the absence of detailed data, the creep crack growth rate can be estimated from da/dt = (3/εf)C*0.85, where C* is the creep fracture parameter under idealised steady-state conditions and εf is creep ductility, or strain to failure, for the stress state of the crack tip [7]. For the alternative approach of considering nominally defect-free material, the combined effects of creep and fatigue can be expressed by combining the expressions [8.4] and [8.8] as X ti X Nj þ ¼D ½8:11 tri Nrj j i where D ≤ 1 is now the combined limit for the overall life fraction. This combined version of this simple linear life fraction rule, or the Robinson– Miner rule, has been widely applied in life assessment of components under combined creep and fatigue loading. It is an approximation, however, and this is reflected in Fig. 8.12, where the region for microcracking, with differing limits for creep and fatigue dominated regimes, corresponds to a modified version of the expression [8.11]. Nevertheless, the approach is widely used [4, 6], and a simple schematic application is shown in Fig. 8.13 for setting the inspection and repair/replacement limits. Typical expected
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8.12 Linear life fraction mapping from combined creep and fatigue [9].
8.13 Utilising the life fraction principle for maintenance of gas turbines (schematic only).
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values for tr may vary from about 10 000 h for some aircraft turbines to about 100 000 h for many land-based turbines; typical expected values for Nr may vary from hundreds to thousands of cycles depending on the component and type of service (cyclic or base load, cold/warm/hot starts or tripping cycles).
8.5
Microstructural degradation
Microstructural degradation can be defined as an adverse deviation from the intended or expected microstructure. When this already happens before service, it is a deviation from the state in which the material should be after all processing treatments in manufacturing or repair, e.g. after casting, forming, welding, brazing and heat treatments. However, much of common degradation is caused by thermally activated change in the microstructure, such as precipitation, growth or modification of certain phases and particles in the alloy and its coatings during service. Degradation will reduce the creep resistance by coarsening and other modifications of the strengthening gamma prime (γ´, Ni3(Ti,Al)) intermetallic precipitates in nickel superalloys (Figs 8.14 to 8.18) or embrittlement if carbides or other weakening and embrittling (like intermetallic topologically close-packed, TCP [10]) phases grow to continuous or otherwise extensive formations on, for example, alloy grain boundaries or under coatings (Figs 8.14 and 8.19). For maximum creep strength, the mean size of γ´ precipitates should be about 0.45 μm [11, 12]. To retard the growth of the effective mean size, the initial γ´ size distribution can be deliberately two-peaked with larger primary and smaller secondary particles (Fig. 8.15). However, in modern single crystal nickel alloys the initially cubic γ´ particles of similar size and orientation fill up most of the available space (Fig. 8.17), with thinner layers of surrounding matrix (gamma) phase [13]. On external or internal surfaces the degradation can be assisted by loss or addition of some elements due to the chemical potential differences between the surface and the interior. This happens, for example, when contents of aluminium and titanium of Ni alloys are reduced by diffusion to the surfaces and removed by oxidation, resulting in a loss of the gamma prime phase (Fig. 8.20). As the chemical compositions have evolved with generations of turbine materials with component-specific properties, microstructural degradation is very much dependent on the material and application [12–14]. For example, the changes are generally less dramatic in solution and carbide strengthened nickel and cobalt alloys than in alloys that for good strength include high volume fractions of gamma prime precipitates in the as-new state. The thermally activated change can also usefully indicate the local in-service temperature that may otherwise be poorly known.
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8.14 Microstructure of ex-service cast turbine blade (Udimet 500); dark particles are y´ precipitates (~ 30 % vol) that have somewhat coarsened in service. In addition, grain boundary carbides have grown to a locally nearly continuous chain.
8.15 Microstructure of cast blade (GTD-111) in a nearly as-new condition, including large primary and smaller secondary (<<0.5 μm in size) γ´ particles for strengthening.
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8.16 Microstructure of an ex-service blade (GTD-111), with coarse, partially fused γ´ particles, with no small particles left in between.
8.17 Microstructure of a single crystal (SX) blade, with a nearly as-new condition (70 % γ´ vol).
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8.18 Agglomeration and rafting of the γ–γ´ microstructure (SX blade as in Fig. 8.17).
For diffusive thermal coarsening of γ´ precipitates that have grown in time t from an initial mean effective radius r0 to the radius r, Kt ¼ r3 r30
½8:12
where the rate factor K for the coarsening process with an activation energy Q is given by 1 Q K ¼ exp B ½8:13 T RT where B is a material-related constant and R is the gas constant. With known values of B, Q and time in service, and measured mean r (and r0 from either a new blade or a well-cooled ex-service location), the effective temperature can be solved and used to support life assessment. Similarly, growth of other phases such as layers under coatings, oxides inside cracks or depleted layers on crack surfaces may be applicable for the same purpose. When the microstructural changes such as γ´ coarsening is moderate, the change of the base alloy can be largely reversible by suitable heat treatment (solution and precipitation annealing). However, extensive change is
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8.19 Plate-like TCP (sigma) phase precipitates below the coating (SX blade as in Fig. 8.17).
8.20 Whitish zone on crack surfaces due to oxidation-induced depletion of aluminium and titanium (and therefore loss of the γ´ phase) in the microstructure; SX blade as in Fig. 8.17. © Woodhead Publishing Limited, 2011
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8.21 Cross-section of a turbine blade with severe irreversible degradation leading to replacement; note that extensive internal damage often combines with strong oxidation and/or hot corrosion.
irreversible and will typically result in at least local repair or replacement (Fig. 8.21).
8.6
Future trends
From the point of view of high temperature performance and efficiency, the development of gas turbines is more likely to be gradual evolution than major leaps forward. This is partly due to diminishing returns from higher temperatures and pressures in the combustion turbine process and partly because of the rather gradual improvements in the protective cooling and coating solutions, and in the development of the (metallic) turbine materials. The course of development in land-based gas turbines is in part driven by cost and availability of suitable fuels, such as shale gas and syngas from various sources, and affected by the rather political process towards the low carbon world. In the past, new materials for the high temperature end of gas turbines have often been first introduced to aircraft engines, and the trend may well continue due to the demand for ever higher efficiency and affordability. Nevertheless, any new alloy must fulfil the basic requirements for mechanical (creep and fatigue) strength, supported by the microstructural stability. The near future development appears to stay with coated
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single crystal nickel superalloys in the hot end turbine blades, allowing for material temperatures up to about 1000–1150 8C, depending on loading details and the targeted design life. For future turbines, Nb alloys have been suggested up to about 1200 8C and Mo alloys up to about 1300 8C under favourable loading conditions, with the added benefit in efficiency if blade cooling and the related losses can be reduced or even avoided [15].
8.7
Conclusion
Together with the effect of cooling and coating systems, the damage mechanisms of creep, fatigue and microstructural degradation in gas turbine materials, and particularly in nickel and cobalt superalloys of the hot end turbine blades, will largely determine the life, performance and efficiency of the turbine. The damage mechanisms need to be accounted for in design and life assessment of the critical components, and also in the development of new materials or new turbines, whether for land-based power generation or for aircraft engines. Creep resistance is particularly required for the highly stressed rotating blades, with sufficient ductility and microstructural stability to resist early crack growth and thermal fatigue. Fatigue damage from fluctuating (cyclic) stress is relatively less than creep promoted by increasing temperature, but can limit the design and life of components also at low or modestly elevated temperatures. In components and operating conditions where both creep and fatigue can be expected, the combined impact can be considered by using, for example, suitable versions of the Robinson–Miner life fraction rule. Examples of the expected rates of creep and fatigue damage have been given above for selected blading materials and testing conditions, and examples of microstructural damage have been presented on selected alloys. Creep, fatigue and microstructural degradation will proceed faster with increasing temperature and stress, but the rates of these mechanisms also depend on the material and its initial state. The mechanisms tend to accelerate towards the end of life, and the observed damage and degradation can usefully support life assessment by indicating the effective condition and service temperatures of hot path components.
8.8
References
1. Stringer J R, ‘Alloys for advanced power systems’, in Natesan K, and Tillack D J (eds), Heat Resistant Materials, ASM International, Metals Park, Ohio, 1991, pp .9–23. 2. Viswanathan R, Damage Mechanisms and Life Assessment of High-Temperature Components, ASM International, Metals Park, Ohio, 1989, 497 pp. 3. NRIM Creep Data Sheet 34B, NRIM, Tokyo, 1993, 62 pp.
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4. Kloos K H, Granacher J and Windecker T, ‘High temperature behaviour of materials for gas turbine blades manufactured in new technology’ (in German), Materialwissenschaft und Werkstofftechnik, 25, 1994, 235–243. 5. Meetham G W (ed.) The development of gas turbine materials. Applied Science Publishers 1981, 306 p. 6. Gampe U, Raddatz M, de Dios R P and Freimark M, ‘Effects of non-uniform gas temperatures at turbine inlet on component lifetime and performance’ (in German), VGB Powertech, 8, 2005, 75–81. 7. Webster G A and Ainsworth R A, High Temperature Component Life Assessment, Chapman & Hall, 1994, 327 pp. 8. Holdsworth S R and Hoffelner W, ‘Fracture mechanics and crack growth in fatigue’, in Brunetaud R et al. (eds), High Temperature Alloys for Gas Turbines, Reidel Publishing, Dordrecht, 1982, pp. 345–368. 9. Auerkari P, Salonen J, Ma¨kinen S, Karvonen I, Tanttari H, Kangas P, Scholz A and Vacchieri E, ‘Life assessment of gas turbine blades after long term service’, in Proceedings of the 9th Liege Conference on Materials for Advanced Power Engineering, Liege, Belgium. 27–29 September 2010. 10. Seiser B, Drautz R and Pettifor D G, ‘TCP phase predictions in Ni-based superalloys: structure maps revisited’, Acta Materialia, 59, 2011, 749–763. 11. Reed R C, The Superalloys – Fundamentals and Applications, Cambridge University Press, 2006, 372 pp. 12. Durand-Charre M, The Microstructure of Superalloys, Gordon and Breach Science Publishers, 1997. 124 pp. 13. Scholz A, Nazmy M, Fedelich B, Tinga T and Huls R, ‘Modelling of microstructure and mechanical property changes in gas turbine alloys’, in Proceeding of the 9th Liege Conference on Materials for Advanced Power Engineering, Liege, Belgium. 27–29 September 2010. 14. Donachie M J, ‘Relationship of properties to microstructure in superalloys’, in Donachie M J (ed.), Superalloys Source Book, ASM International, Metals Park, Ohio, 1984, pp .102–111. 15. Perepezko J H, ‘The hotter the engine, the better’, Science, 326, 2009, 1068–1069.
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9 Gas turbine materials selection, life management and performance improvement T . A´ L V A R E Z T E J E D O R , Endesa Generacio´n, Spain
Abstract: The aim of this chapter is to provide a comprehensive review of the material technology used for high-temperature applications and their impact on the operational life of the gas turbine. Market advantage today relies on increasing performance and efficiency as well as reducing life-cycle costs. The first statement is responsible for pushing present materials and coatings to their limits with significant consequences in terms of the durability and maintenance costs for machines that rely on these advanced hot section designs. Gas turbine material selection will greatly impact both gas turbine performance and life-cycle costs, in such a way that the correct selection will make gas turbine technology successful in the electric power generation market. Key words: hot gas path, superalloy, thermomechanical fatigue, creep, fracture mechanics, life management.
9.1
Introduction
The gas turbine industry must focus on several key factors that will make its future power generation technology successful in the electric power generation market. These factors are summarized in the list below and in Fig. 9.1 [1]: . . . .
competitive economic performance (i.e. higher efficiency and optimized life-cycle costs); reliable operation under a cycle duty (repeated gas turbine startups and shutdowns); increased dependability of current and future plants (reliability, availability, maintainability and durability, or RAM-D); ability to meet regulatory emissions levels and achieve high thermal efficiencies; and
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9.1 The new competitive arena for power generation [1].
.
reliable fuel-switching capability and fuel flexibility.
Gas turbines will be one of the most important horizontal technologies and will play an essential role in meeting these requirements. Gas turbine technology is considered as horizontal due to its capacity to be widely applied across many different types of power plant configurations, while running with different fuels (coal gas, natural gas, hydrogen, liquid fuels, etc.). Gas turbine performance and life management arise as the way to achieve competitive advantages that will enable gas turbine technology to gain the edge over their competitors. The focus is therefore on the ‘heart’ (core) of gas turbine technology. The ‘hot gas path’ of a gas turbine is the core of the engine, which includes the combustion chamber, the transition pieces and the turbine section. The main drivers for improving hot gas path behavior are: .
.
Gas turbine performance – this is highly dependent on the turbine inlet temperature, which results in a greater need for the hot gas path components to achieve high thermal efficiencies with low nitrogen oxide (NOx) emissions. Gas turbine life-cycle costs – this is strongly affected by the costs of hot gas path components and maintenance, which gives rise to maintenance practices and inspection techniques that in turn allow the improvement of gas turbine dependability, i.e. its RAM-D.
Gas turbine material selection will greatly impact both gas turbine performance and life-cycle costs in such a way that the correct selection will
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make gas turbine technology successful in the electric power generation market. The primary philosophy is to build a reliable, efficient, cost-effective machine for the intended service. This includes a materials development program, which is expensive and time-consuming. First, new ideas and emerging developments are screened to select the one or two with the best potential for satisfying the material design goals. Extensive testing follows to ensure that the materials will perform satisfactorily in heavy duty gas turbines for tens of thousands of hours. Long-term creep testing at the expected operating temperatures of the material is conducted to characterize alloy performance. Additionally, laboratory evaluations typically include items such as tensile, rupture, low- and high-cycle fatigue, thermal mechanical fatigue, toughness, corrosion/oxidation resistance, production/ processing trials and complete physical property determination. This phase of testing can last several years for a new nozzle or blade material. After laboratory testing comes the actual machine-operating experience, the best and final test of a new material to be compared and evaluated with the current baseline material.
9.2
Superalloys
As mentioned above, the hot gas path of a gas turbine includes the combustion chamber, the transition pieces and the turbine section. The turbine section is constructed around several rows of blades and vanes. The vanes in the first stage will become the hottest as they are located closest to the combustion chamber. Then a significant performance parameter is defined, the firing temperature [2], which is thought to be the highest temperature reached in the Brayton cycle. It is usually defined as the mass–flow mean total temperature at the stage 1 vane trailing edge plane. Currently all first stage vanes are cooled to keep the temperatures within the operating limits of the materials being used. The two types of cooling currently employed are air and steam. The blades and vanes in the turbine section will determine to a large extent the ultimate efficiency of the gas turbine. These parts have to work under extreme conditions, operating in high temperatures in an oxidizing environment while being subjected to large thermal and mechanical stresses. In order to increase the durability of the blades and vanes in these extreme conditions, special metal superalloys have been developed. The high-quality technologies used in the manufacture of the turbine blades make them the most expensive parts of the gas turbine. To achieve higher thermal efficiencies, higher combustion temperatures are needed; however, higher combustion temperatures – from around 1540 8C (2800 8F) – exacerbate NOx emissions. To combat excessive NOx
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emissions, ‘dry’ low oxides of nitrogen (NOx) and ‘ultralow’ NOx combustors are introduced, as well as alternative methods for achieving ultralow NOx emissions, including rich-burn, quick-quench, lean-burn and catalytic combustors [3]. Adding further to these technological limitations, extremely high operating temperatures – greater than 1290 8C (2350 8F) – are beyond the material tolerances of the turbine blades and vanes. Therefore, the goal of achieving 60 % efficiency while staying below 10 ppm of NOx emissions is constrained by the thermal emission reduction and material limits of the gas turbine system. There are four main innovations that are critical in meeting this need for high efficiency and low emissions: . . . .
advanced cooling systems; single-crystal superalloy casting; thermal barrier and metallic coating; and lean pre-mix dry low-NOx combustors.
On the other hand, to optimize the life-cycle cost of gas turbines, special attention must be paid to the hot gas path components: typically, more than around 70 % of the total gas turbine maintenance cost corresponds to scheduled maintenance, parts and materials. This will lead to the establishment of mechanisms for risk mitigation, such as long-term service agreements (LTSAs) [4], business interruption insurance, extended guarantees and part-cost guarantees. Apart from the above considerations, it is also necessary to take into account current operational conditions in a deregulated electricity market. These conditions require more flexible operations with high efficiency and low emissions for the whole power range, high operational reliability and better maintainability. Most heavy-duty gas turbines for operation in land-based applications use proven technology, derived from aircraft and steam turbine applications. However, the unique requirements and special conditions for heavyduty gas turbines demand special materials and processes. The materials used in stationary applications can be classified in three groups: stainless steels (iron-based), nickel-based alloys and cobalt-based alloys. The alloy composition has to be a compromise between mechanical strength and the corrosion and oxidation resistance while ensuring a proper economical lifetime. Most of the alloys applied in the turbine section have a composition with high nickel or cobalt contents, resulting in good mechanical properties. Eventually, the cast superalloys for the highest temperatures are protected against oxidation and corrosion by chromium and aluminum coatings. The development of the increase in firing temperature and material properties is illustrated in Fig. 9.2. In the early years of turbine development, increases in blade alloy temperature capability accounted
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9.2
Firing temperature trend with blade material improvement [5].
for the majority of the firing temperature increase until the 1970s when aircooling was introduced, which decoupled the firing temperature from the blade metal temperature. Also, as the metal temperatures approached the 870 8C (1600 8F) range, oxidation and corrosion of blades became more life limiting than strength until the introduction of protective coatings. During the 1980s, emphasis turned toward two major areas: improved materials technology, to achieve greater blade alloy capability without sacrificing alloy corrosion resistance, and advanced, highly sophisticated air-cooling technology, to achieve the firing temperature capability required for the new generation of gas turbines. The use of steam cooling to further increase combined-cycle efficiencies in combustors was introduced in the mid to late 1990s. Since 1950, turbine blade material temperature capability has advanced by approximately 472 8C (850 8F), approximately 10 8C (20 8F) per year. The importance of this increase can be appreciated by noting that an increase of 56 8C (100 8F) in the turbine firing temperature can provide a corresponding increase of 8–13 % in output and a 2–4 % improvement in simple-cycle efficiency [5]. This technological development has been mainly possible thanks to the new generation of advanced materials called superalloys. The denomination of superalloys is used to those alloys generally used at temperatures above around 540 8C (1000 8F), i.e. nickel-base, iron–nickelbase and cobalt-base corrosion-resistant alloys. The iron–nickel-base superalloys are an extension of stainless steel technology and generally are wrought, i.e. formed to shape or mostly to shape by hot rolling, forging, etc. The cobalt-base and nickel-base superalloys, on the other hand, may be
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9.3 Crystal structure of gamma and gamma prime.
either wrought or cast depending on the application or the alloy composition involved. Superalloys consist of an austenitic face-centered-cubic (fcc) crystal structure matrix phase, gamma (γ), plus a variety of secondary phases. Important secondary phases are gamma prime (γ´) fcc ordered Ni3(Al, Ti) and various MC, M23C6, M6C and M7C3 (rare) carbides in nickel-base and iron–nickel-base superalloys (Fig. 9.3). Carbides are the principal secondary phases in cobalt-base alloys. Also, γ´, a body-centered tetragonal (bct) phase of ordered Ni3Nb, a hexagonal ordered Ni3Ti (η) phase and the δorthorhombic Ni3Nb intermetallic phase can be found in nickel-base and iron–nickel-base superalloys. The strengthening process in superalloys, and hence the mechanical properties of superalloys, can be modified considerably by manipulating the strengthening level achieved. The superalloys derive their strength from solid-solution hardeners and secondary precipitate phases that form in the γ matrix and produce precipitation (age) hardening. The principal strengthening precipitate phase in nickel-base and iron–nickel-base superalloys is γ’ (gamma prime). Additionally, carbides may provide limited strengthening directly (e.g. through dispersion hardening) or, more commonly, indirectly (e.g. by stabilizing grain boundaries against movement). The δ and η phases are useful (along with γ´) in controlling the grain structure of wrought superalloys during processing. By controlling grain structure, strength can be significantly influenced. The extent to which the second phases contribute directly to strengthening depends on the alloy and its processing. It should be noted that improper distributions of carbides and precipitate phases can be detrimental to the mechanical properties. In addition to those elements that produce solid-solution hardening and/or promote carbide and γ´ formation, other elements (e.g. boron, zirconium, hafnium) are added to enhance mechanical or chemical properties.
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Table 9.1
Role of alloying elements in superalloys [6]
Effect
Cobalt base
Solid-solution strengtheners fcc matrix stabilizer Carbide form MC M7C3 M23C6 M 6C Carbonitrides: M(CN) Promotes general precipitation of carbides Forms γ´ Ni3(Al,Ti) Retards formation of hexagonal η (Ni3Ti) Raises solvus temperature of γ´ Hardening precipitates and/or intermetallic phases Oxidation resistance Improve hot corrosion resistance Sulfidation resistance Improves creep properties Increases rupture strength Grain-boundary refiners Facilitates working Retard γ´coarsening
Nb, Cr, Mo, Ni, W, Ta Co, Cr, Fe, Mo, W, Ta, Re Ni —
a b c
Nickel base
Ti Cr Cr Mo, W C, N
W, Ta, Ti, Mo, Nb, Hf Cr Cr, Mo, W Mo, W, Nb C, N
— —
— Al, Ti
—
—
—
Co
Al, Mo, Tib, W, Ta
Al, Ti, Nb
Al, Cr
Al, Cr, Y, La, Ce
La, Y, Th Cr — B, Zr — Ni3Ti —
La, Th Cr, Co, Si B, Ta Bc B, C, Zr, Hf — Re
Not all these effects necessarily occur in a given alloy. Hardening by precipitation of Ni3Ti also occurs if sufficient Ni is present. If present in large amounts, borides are formed.
Superalloy microstructure, chemical composition and proper control thereof are complex. As many as 14 elements may be added in some superalloys. The complexity of the metallurgy is best illustrated by Table 9.1, indicating the effect of the major alloying elements. The Ni- and Co-based alloys, usually indicated as superalloys, are applied because of their high strength at high temperatures. Co-based alloys are mainly used for (stationary) vanes, whereas in general the Ni-base alloys are used for (rotating) blades. Of course the materials selection varies per manufacturer and per gas turbine type. Although many different alloys exist there are a number of alloys that are widely applied by most of the manufacturers. Table 9.2 gives the chemical composition of a number of alloys, which at this moment are considered to be the ‘state of the art’ for the industrial gas turbine (IGT). Although most manufacturers use identical
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Buckets U500 U700 (Rene 77) Alloy 738 MAR M247 GTG-111™ GTD-444 ™ PWA 1483 Rene N5 CMSX-4® PWA 1484 Nozzles FSX414 GTG-222™ GTG-111™ Rene N5
Wt %
18.50 15.00 16.00 8.25 14.00 9.80 12.80 7.00 6.50 5.00
28.00 22.50 14.00 7.00
10.2 Bal Bal Bal
Cr
Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal
Ni
Bal 19.00 9.50 7.50
18.50 17.00 8.30 10.00 9.50 7.50 9.00 7.50 9.00 10.00
Co
1.00
0.20
Fe
2.30 1.50 1.50
4.00 5.30 1.75 0.80 1.50 1.50 1.90 1.50 0.60 2.00
Mo
7.00 2.00 3.80 5.00
2.60 10.00 3.80 6.00 3.80 5.00 6.00 6.00
W
1.20 4.90
1.00
1.00
1.20 4.90
3.00 3.35 3.40 1.00 4.90 3.50 4.00
Ti
3.00 3.35 3.40 1.00 4.90 3.50 4.00
Al
0.50
0.90
Nb
1.00 2.80 6.50
1.75 2.80 2.80 4.80 4.00 6.50 6.50 9.00
Ta
Table 9.2 Nominal composition of IGT cast Co-base and Ni-base superalloys [7] Mn
V
0.010 0.008 0.010 0.004
0.05
0.25 0.10 0.10 0.05
0.004
0.10 0.08
B 0.006 0.020 0.001 0.015 0.010 0.009
0.07 0.07 0.10
C
Re 3.0;Hf 0.15;Y 0.01
Re 3.0; Hf 0.15; Y 0.01 Re 3.0; Hf 0.10 Re 3.0; Hf 0.10
Hf 0.15
Hf 0.15
Other
Gas turbine materials selection
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alloy identifications there may well be differences in alloy composition or heat treatment, resulting from improvements by each manufacturer. Since the design of turbomachinery is complex and efficiency is directly related to material performance, material selection is of prime importance. Turbine components must operate under a variety of stress, temperature and corrosion conditions. Compressor blades operate at a relatively low temperature but are highly stressed. The combustor operates at a relatively high temperature and low-stress conditions. The turbine blades operate under extreme conditions of stress, temperature and corrosion. Advances in alloys and processing, while expensive and time-consuming, provide significant incentives through increased power density and improved efficiency.
9.2.1 Metallurgical behavior The required material characteristics in gas turbine applications for high performance and long life include limited creep, high-rupture strength, resistance to high-temperature corrosion, good fatigue strength, low coefficient of thermal expansion and high-thermal conductivity to reduce thermal strains. High-temperature corrosion plays an important role in the selection of materials for gas turbine applications. The principal modes of hightemperature corrosion frequently responsible for equipment problems are oxidation, carburization, sulfidation, nitridation, halogen gas corrosion, ash/salt deposit corrosion, molten salt corrosion and liquid metal corrosion. Oxidation and hot corrosion (sulfidation) mechanisms are the most important ones in this discussion and are described in the next sections. Thus, the failure mechanism of a turbine blade is related primarily to creep and corrosion and secondarily to thermal fatigue. Satisfying these design criteria for turbine blades will ensure high performance, long life and minimal maintenance. Understanding mechanical behavior and how temperature affects the properties of the materials is an essential part for a proper material selection and design. All material properties change with temperature. Some do so in a simple linear way making compensation easy, for instance the density and the modulus. Others, however, particularly the yield strength and the rates of oxidation and corrosion, change in more sudden ways, which if not allowed for, can lead to disaster. Thermal conductivity and conductivity for matter flow (diffusion) change in more complex ways. The last of these is particularly important in our discussion, due to diffusion (the intermixing of atoms in solids and the ways it allows creep and creep fracture) has a profound effect on mechanical properties when temperatures are high. To understand and use diffusion we
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need the idea of thermal activation, the ability of atoms to jump from one site to another, using thermal energy as the springboard. Mechanical properties of interest for elevated-temperature applications include short-time elevated-temperature tensile properties, creep and stressrupture, low-cycle and high-cycle fatigue, thermal and thermomechanical fatigue and creep–fatigue interaction. Thus, extensive testing is conducted to ensure that the materials will perform satisfactorily in tens of thousands of hours. The influence of temperature on the strength of materials can be demonstrated by running standard, short-time tensile tests at a series of increasing temperatures where materials are taken to failure. Such test conditions are often called ‘static’ or monotonic conditions, and allow for the strain to develop with the load being applied gradually until the specimen fails. Typical tensile stress–strain curves for an alloy are defined at different temperatures, from which such useful properties as the ultimate tensile strength (UTS), yield stress (proof stress), elastic modulus, ductility and toughness modulus can be obtained. Stresses above the elastic limit cause permanent deformation (ductile behavior) or brittle fracture. Gas turbine components should not fail (break) when subjected to the ultimate load. This would mean that any amount of plastic deformation is allowable providing that the component does not break in a brittle fracture. Ductility and toughness (resistance to fracture) properties are required for alloys in gas turbine components. Ductility is an important mechanical property and commonly measured by elongation and reduction in area (Table 9.3). It is a measure of the degree of plastic deformation that has been sustained at fracture. A material that experiences very little or no plastic deformation upon fracture is termed brittle (brittle materials are considered to be those having a fracture strain of approximately less then about 5 %). Ductility may be expressed quantitatively as either percent elongation or percent reduction in area, i.e. lf l0 %EL ¼ 6100 ½9:1 l0 or %ROA ¼
A0 Af 6100 A0
½9:2
The area under the elastic part of the stress–strain curve is identified as the elastic energy stored per unit volume (σy2/(2E)). Beyond the elastic limit plastic work is done in deforming a material permanently by yield or
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9.3 Room-temperature mechanical properties (in tension) for various materials [8] Yield strength Tensile strength Material
MPa
ksi MPa metal alloys
ksi
Ductility, %EL (in 50 mm (2 in))
Molybdenum Titanium Steel (1020) Nickel Iron Aluminum
565 450 180 138 130 35
82 65 26 20 19 5
95 75 55 70 38 13
35 25 25 40 45 40
655 520 380 480 262 90
crushing. The increment of plastic work done for a small permanent extension or compression dL under a force F, per unit volume V = AL0, is dWpl ¼
F dL F dL ¼ ¼ s depl V A0 L0
½9:3
Thus, the plastic work per unit volume at fracture, important in energyabsorbing applications, is Z ef Wpl ¼ s depl ½9:4 0
which is just the area under the stress–strain curve (toughness module). Toughness requires a new material property, fracture toughness (resistance of materials to cracking and fracture), which is developed in the next sections. In load-limited designs, the best material selection involves a combination of fracture toughness and Young’s modulus. Having assessed the loads that will act upon a gas turbine component with a defined geometry (stress analysis), the effect of the applied load compared with the strength and the other relevant properties of the material selected will reveal whether it is favorable for the intended service. It is common to impose a safety factor into the design, in order that an adequate margin of safety be established or to allow for uncertainty in material properties, i.e. variability. The common factors used are the proof and ultimate factor, scatter factor, casting factor, stress concentration factor, etc. Once we consider the ability of a given material to resist load it becomes quite apparent that the way in which the load is applied and the conditions under which it is applied are very important. It leads us to consider the failure modes as well as the material’s ability to resist these failure modes. In particular, the availability of creep-resistant materials has proved to be extremely useful whenever components are operating at high temperatures,
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9.4 Creep-rupture schematic curve (time-dependent deformation under constant load at constant high temperature followed by final rupture, where all loads are below the short-time yield strength) [9].
i.e. in turbomachine casing, bolts and studs, turbine blades and nozzles, compressor and turbine disc applications. The phenomenon known as creep, in which progressive deformation may occur under the application of a constant load, has been known for many years. Creep is defined as the tendency of a solid material to slowly deform plastically, under the influence of (elastic) stresses. Creep is a temperatureand time-dependent phenomenon. High temperature results in a higher mobility of dislocations by the mechanism of climb and in an increase in the equilibrium concentration of vacancies. Grain boundaries become less well defined at relative low temperatures (as low at 0.4Tm, where Tm is the melting point), and there is a greater mobility of atoms at elevated temperatures. Then, we have to consider creep as a failure mode at running temperatures well below the melting point of the material. The melting point of different metals varies considerably, and their strengths at various temperatures are different. At low temperatures all materials deform elastically, then plastically, and are time-independent. However, at higher temperatures, deformation is noted under constant load conditions (within the elastic range of the material). This high-temperature, time-dependent behavior is called creep-rupture. Figure 9.4 shows a schematic of a creep curve with the various stages of creep. In many respects such materials behave in a viscoelastic manner, and when subject to a constant tensile load at elevated temperature undergo a time-dependent increase in dimension, i.e. they creep. Fig. 9.4 shows the generally accepted idealization of the three-stage creep process, where ε0 is an instantaneous elastic stage prior to stage I and de=dt /or e_ is known as the creep rate. After the initial, virtually instantaneous, elastic straining (ε0), stage I
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9.5
Time-dependent deformation.
represents a region of primary creep in which the creep resistance of a material increases as a function of its own deformation and is characterized as a decreasing creep rate. Stage II creep, known as secondary creep, is a period with a nearly constant creep rate, resulting from the balance between the competing processes of strain-hardening and recovery. Hence, second stage creep is often referred to as steady-state creep and the average value of the creep rate during this stage is called the minimum creep rate (Fig. 9.5). Stage III or tertiary creep reveals itself in the form of cavities (voids) at grain boundaries, being the behavior used to identify whether or not a creeploaded component is approaching end-of-life. Due to the fact that superalloy creep stage III is developing very fast, it is almost impossible to clearly identify the point in time of transition from stage I/II to stage III. In actual superalloy components the formation of creep voids is hardly ever observed and used as end-of-life criteria. Stage III or tertiary creep occurs mainly in constant load–creep tests at high stress and temperature when there is an effective reduction in crosssectional area usually produced by necking. The tertiary strain rate increases rapidly until fracture (rupture) occurs. There are quite often metallurgical changes associated with tertiary creep. Andrade [10] attempted to characterize the creep curve putting forward that creep is composed of two separated processes: (a) transient creep with dε/dt decreasing in time and (b) a constant dε/dt viscous creep component. Andrade’s equation [10] in terms of strain is e ¼ e0 ð1 þ btð1=3Þ Þekt
½9:5
where ε is the strain at time t. It should be noted that Andrade’s equation [10] does not include allowance for tertiary creep (if it exists). Walles and Graham [11] introduced a third term into the Andrade equation [10] to arrive at the complete equation: e ¼ at1=3 þ bt þ ct3
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Norton [12] suggested a simplified approximate form as (the Norton creep law) e_ ¼ Bsn
½9:7
where B, n are material parameters and e_ is the stationary strain rate. The nature of this creep depends on the material, stress, temperature and environment. Limited creep (less than 1 %) is desired for turbine blade application design. Cast superalloys fail with only a minimum elongation. These alloys fail in a brittle fracture even at elevated operating temperatures. The current design for creep is very much based upon empirical materials data. One method of plotting tensile creep data is shown in Fig. 9.6. This format fits in well with a well-established mathematical model for creep, namely: C ¼ Bsn
½9:8
where C = creep rate in e_ tension B = stress intercept for a long-creep rate n = slope of line on a log–log plot σ = applied stress Designing to cope with creep in a gas turbine where clearances are critical means considering creep strain as a design-limiting factor. We need to know how the strain rate or time to failure tf depends on the stress σ and temperature T to which it is exposed. That requires creep testing. The creep test is simple to comprehend since it requires the application of a steady load to a specimen held at constant temperature and the measurement of the strain of the specimen at intervals of time, i.e. the extension is measured as a function of time. Metals have creep curves with the general shape shown in the Fig. 9.4. Creep requires the use of four parameters for its description: time, temperature, stress and strain. Creep tests can be carried out for periods of 2000 to 10 000 h (or more), and be so arranged that strains of less than 0.5 % occur in this time. The initial elastic and the primary creep strains occur quickly and can be treated in much the same way as elastic deflection is allowed for in a structure. Thereafter, the strain increases steadily with time in what is called the secondary creep or the steady-state creep regime. Plotting the log of the steady-state creep rate, e_ , against the log of the stress, σ, at constant time T, as in Fig. 9.7, shows that e_ ¼ Bsn
½9:9
where n, the creep exponent, usually lies between 3 and 8 and for that reason
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9.6 Time to rupture (tr) as a function of the steady-state creep rate (ess ) for single crystals tested in tension at several temperatures [13].
this behavior is called power-law creep. At low σ there is a tail with slope n = 1 (the part of the curve labeled ‘diffusional flow’). As creep continues, damage accumulates. It takes the form of voids or internal cracks that slowly expand and link, eating away the cross-section and causing the stress to rise. This makes the creep rate accelerate, as shown in the tertiary stage of the creep curve of Fig. 9.4. Since ε is proportional to σn with n = 5, the creep rate goes up even faster than the stress: an increase in stress of 10 % gives an increase in the creep rate of 60 %. Materials can deform by dislocation plasticity or, if the temperature is high enough, by diffusional flow or power-law creep. If the stress and temperature are too low for any of these, the deformation is elastic. This shows the range of stress and temperature in which we expect to find each sort of deformation and the strain rate that any combination of them produces (the contours). Diagrams like these (Fig. 9.8) are available for many metals and are a useful summary of creep behavior, helpful in selecting a material for high-temperature applications. Where selecting materials for creep resistance we must therefore consider diffusional flow, which is important when grains are small and when the component is subject to high temperatures at low loads. The way to avoid
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9.7 The stress and temperature dependence of the creep rate [14].
9.8 A deformation mechanism map, showing the regime in which each mechanism operates [14].
diffusional flow is to choose a material with a high melting temperature and a large grain size, so that diffusion distances are long. Single crystals are best of all; they have no grain boundaries to act as sinks and sources for vacancies, so diffusional creep is suppressed completely. This is the rationale behind the wide use of single-crystal turbine blades in jet and industrial engines. That still leaves power-law creep. Materials that best resist power-law creep are those with high melting points, since diffusion and thus creep rates scale as T/Tm, and with a microstructure that maximizes obstruction to
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dislocation motion through alloying to give a solid solution and precipitate particles. Current creep-resistant materials of superalloys are remarkably successful in this. The prediction of rupture times at various combinations of stress and temperature usually involves some measure of extrapolation from shortterm creep tests to long component lifetimes. One method of extrapolation, from short to longer times, is to formulate an equation that describes the creep strain in terms of stress and temperature. This involves the use of relationships knows as time–temperature parameters. The most popular is known as the Larson–Miller parameter. This parameter can be used for long life extrapolation or for assessing cumulative creep damage. Thus, stress-rupture data are often presented in a Larson–Miller curve, which indicates the performance of an alloy in a complete and compact graphical style. While widely used to describe an alloy’s stress-rupture characteristics over a wide temperature, life and stress range, it is also useful in comparing the elevated temperature capabilities of many alloys. The Larson–Miller parameter is PLM ¼ Tð20 þ log tÞ6103
½9:10
where PLM = Larson–Miller parameter. T = temperature (8R) t = rupture time (h) The Larson–Miller parameters are plotted in Fig. 9.9 for the specified turbine blade alloys. Larson and Miller [16] first proposed their method for creep data, i.e. the life expressed at t would be that which reached a particular strain, say 0.1 %. However, this technique has been extended to cover rupture strength, in which case t would be the life to reach fracture. There is some doubt as to whether rupture strength can be considered in the same way as creep strength. It is true that rupture is the terminus at the creep curve, but the point of rupture is dictated by the ductility of the material. In general, the creep ductility is determined by the superposition of strains accumulated in void formation and growth phases separately (Ashby et al. [17]), and is affected by rupture time, which depends on applied stress as well as the steady-state creep rate of alloy. In other words, the decrease in ductility can be tied to quick void formation behavior, causing brittle fracture, while an increase in ductility may be regarded as the result of delayed void formation and the change in the fracture mechanism from intergranular to transgranular, which are very effective in changing crack growth characteristics of steel leading to ductile behavior. In many cases, all three stages of creep shown in Fig. 9.3 are not present.
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9.9 Larson–Miller parameter for various types of blades [15].
At high temperatures or stresses, very little primary creep is seen, while in the case of cast superalloys failure occurs with just a small extension. This amount of extension is ductility. In a time–creep curve there are two elongations of interest. One elongation is from the plastic strain rate and the second elongation is the total elongation or the elongation at fracture. Ductility is erratic in its behavior and is not always repeatable, even under laboratory conditions. Ductility of a metal is affected by the grain size, the specimen shape and the techniques used for manufacturing. A brittle fracture is intergranular with little or no elongation. A ductile fracture is transgranular and typical of normal ductile tensile fracture. Turbine blade alloys tend to indicate low ductility at operating temperatures. As a result, an alloy with low ductility will be sensitive for surface notches and then cracks may develop rapidly from these notches by fatigue or impact loads. As discussed, another important mechanical term in our discussion is toughness, defined as a measure of the ability of a material to absorb energy up to fracture. To determine toughness, we have to consider the specimen geometry as well as the manner of load applications. For dynamic loading (high strain rate) conditions and when a notch (or point of stress concentration) is present, notch toughness is assessed by using an impact test. Furthermore, fracture toughness is a property indicative of a material’s resistance to fracture when a crack is present. For the static (low strain rate) situation, toughness may be obtained from the result of a tensile stress– strain test. It is the area under the σ–ε curve up to the point of fracture. The units for toughness are energy per unit volume. In practice, however, and in particular for the case of rotating machinery,
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the applied loads are seldom constant (static or monotonic condition) and usually fluctuate, either about some mean stress or with complete reversal in sign. This leads to fluctuations in the stresses and strains existing within the components. If these fluctuating stresses are large enough, even though the maximum applied stress may be considerably less than the static strength of the material, failure may occur when the stress is repeated often enough. The connection between the cyclic loading and failure is known as fatigue. Fatigue is defined as the progressive, localized and permanent structural damage that occurs when a material is subjected to cyclic or fluctuating strains at nominal stresses that have maximum values less than (and often much less than) the static yield strength of the material. There are different stages of fatigue damage in an engineering component where defects may nucleate in an initially undamaged section and propagate in a stable manner until catastrophic fracture ensues. For this most general situation, the progression of fatigue damage can be broadly classified into the following stages: . . .
. .
Substructural and microstructural changes that cause nucleation of permanent damage. The creation of microscopic cracks. The growth and coalescence of microscopic flaws to form ‘dominant’ cracks, which may eventually lead to catastrophic failure. (From a practical standpoint, this stage of fatigue generally constitutes the demarkation between crack initiation and propagation.) Stable propagation of the dominant macrocrack. Structural instability or complete fracture.
The conditions for the nucleation of microdefects and the rate of advance of the dominant fatigue crack are strongly influenced by a wide range of mechanical, microstructural and environmental factors. The principal differences among different design philosophies often rest on how the crack initiation and the crack propagation stages of fatigue are quantitatively treated. It is important to note here that a major obstacle to the development of life prediction models for fatigue lies in the choice of a definition for crack initiation. The total fatigue life is defined as the sum of the number of cycles to initiate a fatigue crack and the number of cycles to propagate it subcritically to some final crack size. Classical approaches to fatigue design involve the characterization of total fatigue life to failure in terms of the cyclic stress range (the S–N curve approach) or the (plastic or total) strain range. In these methods, the number of stress or strain cycles necessary to induce fatigue failure in initially uncracked (and nominally smooth-surfaced) laboratory specimens is estimated under controlled amplitudes of cyclic stresses or strains. The
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resulting fatigue life incorporates the number of fatigue cycles to initiate a dominant crack (which can be as high as some 90 % of the total fatigue life) and to propagate this dominant flaw until catastrophic failure occurs. Various techniques are available to account for the effects of mean stress, stress concentrations, environments, multiaxial stresses and variable amplitude stress fluctuations in the prediction of total fatigue life using the classical approaches. Since the crack initiation life constitutes a major component of the total fatigue life in smooth specimens, the classical stressbased and strain-based methods represent, in many cases, design against fatigue crack initiation. Under high-cycle, low-stress fatigue situations, the material deforms primarily elastically. The failure time or the number of cycles to failure under such high-cycle fatigue has traditionally been characterized in terms of the stress range. However, the stresses associated with low-cycle fatigue are generally high enough to cause appreciable plastic deformation prior to failure. Under these circumstances, the fatigue life is characterized in terms of the strain range. The fracture mechanics approach to fatigue design, on the other hand, invokes a ‘defect-tolerant’ philosophy. The basic premise here is that all engineering components are inherently flawed. The size of a pre-existing flaw is generally determined from non-destructive flaw detection techniques (such as visual, dye-penetrant or X-ray techniques or the ultrasonic, magnetic or acoustic emission methods) [18]. The useful fatigue life is then defined as the number of fatigue cycles or time to propagate the dominant crack from this initial size to some critical dimension. The choice of the critical size for the fatigue crack may be based on the fracture toughness of the material, the limit load for the particular structural part, the allowable strain or the permissible change in the compliance of the component. The prediction of crack propagation life using the defect-tolerant approach involves empirical crack growth laws based on fracture mechanics. Various methods are available to incorporate the effects of mean stresses, stress concentrations, environments, variable amplitude loading spectra and multiaxial stresses in the estimation of useful crack growth life. In the safe-life approach to fatigue design, the typical cyclic load spectra, which are imposed on a structural component in service, are first determined. On the basis of this information, the components are analyzed or tested in the laboratory under load conditions that are typical of service spectra, and a useful fatigue life is estimated for the component. The estimated fatigue life, suitably modified with a factor of safety (or an ignorance factor), then provides a prediction of ‘safe life’ for the component. At the end of the expected safe operation life, the component is automatically retired from service, even if no failure has occurred during
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service (and the component has considerable residual fatigue life). This procedure invariably has to account for several unknowns and by selecting a large margin of safety, a safe operating life can be guaranteed, although such a conservative approach may not be desirable from the viewpoints of economy and performance. On the other hand, if fatigue cracks are nucleated in the component during service, the component may well fail catastrophically. As noted by Gurney in 1968 [19], the safe-life approach depends on achieving a specified life without the development of a fatigue crack so that the emphasis is on the prevention of crack initiation [18]. The fail-safe concept, by contrast, is based on the argument that, even if an individual member of a large structure fails, there should be sufficient structural integrity in the remaining parts to enable the structure to operate safely until the crack is detected. The fail-safe approach mandates periodic inspection along with the die requirement that the crack-detection techniques be capable of identifying flaws to enable prompt repairs or replacements. Whatever philosophy is employed in design, it is often preferable and even required in some safety-critical situations, e.g. aircraft and nuclear industries, that the critical components of a structure be inspected periodically. This step eliminates dangerous consequences arising from false estimates and errors in the design stage, especially with the safe-life approach. The three basic types of fatigue properties are [20]: . . .
stress-life (S–N) (design philosophy: safe-life, infinite-life), strain-life (ε–N) (design philosophy: safe-life, finite-life), fracture mechanic crack growth (da/dNΔK) (design philosophy: damage tolerance),
and each property plays a role in the context of a fatigue design philosophy as previously discussed. The safe-life, infinite-life philosophy is the oldest of the approaches to fatigue. Much of the technology in application of this approach is based on ferrous metals, especially steels. Steels are predominant as a structural material, but steels also display a fatigue limit or endurance limit at a high number of cycles (typically > 106) under benign environmental conditions. This limit is the highest stress level that the material can withstand for an infinite number of load cycles without failure. The infinite-life asymptotic behavior of steel fatigue life thus provides a useful and beneficial result of S– N testing. However, most other materials do not exhibit this infinite-life response (see Fig. 9.10). The stress at which a material fails by fatigue after a certain number of cycles is known as the ‘fatigue strength’. For materials such as non-ferrous metals, it is usual to define the design stress as that which occurs at some
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9.10 S–N diagram trends [21].
arbitrary number of cycles. For such materials it is common practice to set an arbitrary value for the fatigue strength (endurance limit) at, say, 107 cycles. Stress is the controlling quantity in this method (S–N data presentation). The most typical formats for the data are plots of the log number of cycles to failure (sample separation) versus either stress amplitude (Sa), maximum stress (Smax) or perhaps stress range (ΔS). In any fatigue analysis (Fig. 9.11) for a particular component it is necessary to take account of the factors that influence fatigue behavior. Some of these factors are the type and nature of loading, size of component, surface finish and directional properties, stress or strain concentrations, mean stress or strain, environmental effects, etc. Mean stress influences are also very important, and each design approach must consider them. The reversed cycles employed when deriving an S–N curve would not produce the same amount of damage as a cyclic stress superimposed upon a mean stress. Therefore carrying out a series of tests involves various combinations of ± σr and σm in such a way that a number of methods of plotting such data are found. The expressions that define the three lines shown in Fig. 9.12 are as follows: Goodman [22]: sr sm þ ¼1 se sult
½9:11
Soderberg [23]: sr sm þ ¼1 se sy
½9:12
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9.11
Allowance for factors that affect fatigue [21].
9.12
Effect of mean stress on alternating stress amplitude [21].
Gerberg [24]: 2 sr sm þ ¼1 se sult
½9:13
Almost any variation in the environmental conditions will affect the fatigue life of a component. In particular, the effects of temperature and corrosive materials are most pronounced. The combined action of repeated loading and a corrosive environment is usually known as corrosion fatigue. The combined effect of cyclic stress and
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corrosion usually (but not always) reduces the fatigue life of a component. Although we can appreciate that corrosion on its own can produce pitting of the surface and that this could in turn provide a notch from which fatigue can propagate, the combined effect of cyclic stress and corrosion is much more than this. It is found that the chemical attack greatly accelerates the rate at which fatigue cracks propagate. Some materials that have a definite fatigue (endurance) limit, when tested in air, are found to have either a lower or no such limit when tested in a corrosive atmosphere. The effects of corrosion fatigue can be reduced in a number of ways. However, in general the best approach is to emphasize the corrosionresisting properties of the material rather than the mechanical fatigue properties. Protection of the metal from contact with the corrosive medium by means of metallic coatings has been found to be successful providing that the coating does not become ruptured by the cyclic strain, which will be discussed in the next section. Strain life is the general approach employed for a continuum response in the safe-life, finite-life regime. It is primarily intended to address the lowcycle fatigue area (e.g. from approximately 102 to 106 cycles). The ε–N method can also be used to characterize the ‘long-life’ fatigue behavior of materials that do not show a fatigue limit. From a properties standpoint, the representations of strain-life data are similar to those for stress-life data. However, because plastic strain is a required condition for fatigue, strain-controlled testing offers advantages in the characterization of fatigue crack initiation (prior to subsequent crack growth and final failure). The S–N method is based on just one failure criterion, the total separation of the test coupon. In contrast, any of the following may be used as the failure criterion in strain-controlled fatigue testing: separation, modulus ratio, microcracking (initiation) or percentage of maximum load drop. This flexibility can provide better characterization of fatigue behavior [20]. The S–N and ε–N techniques are usually appropriate for situations where a component or structure can be considered a continuum (i.e. those meeting the ‘no cracks’ assumption). In the case of a crack-like discontinuity, the S– N and ε–N techniques offer little or a quantitative basis for assessment of fatigue life. Once a crack has formed in a component, even static loads producing average tensile stresses well below the material’s nominal strength may produce fracture, particularly in relatively brittle materials. The reason lies in the formation of high-stress concentrations at the leading edge of the crack. ‘Fracture mechanics’ investigations have shown that the fracture toughness of a material at a given temperature is proportional to a stress level and to the square root of a crack dimension. The fracture toughness can thus be expressed by a single parameter, the critical stress intensity
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factor K, which has units of MPa m0.5 (or MN/m1.5) and which can be determined experimentally by producing a crack from cyclic tests and then loading it statically until it fractures. Fracture mechanics methodology offers considerable promise for improved understanding of propagation of fatigue cracks and problem resolution in designing to prevent failures by fatigue. The characterization and quantification of the stress field at the crack tip in terms of stress intensity in linear elastic fracture mechanics allow us to recognize the singularity of stress at the tip and provides a controlling quantity and measurable material property. A more accurate calculation of this critical crack size can be obtained by elastic–plastic fracture mechanic calculations. This is, however, a much more complicated calculation technique, where accurate material properties (KIC) are needed. The use of stress intensity as a controlling quantity for crack extension under cyclic loading thus enhances the engineering analysis of the fatigue process. Initiation of fatigue cracks in structural and equipment components occurs in regions of stress concentrations, such as notches, as a result of stress fluctuation. The material element at the tip of a notch in a cyclically loaded component is subjected to the maximum stress range, Δσmax. Consequently, this material element is most susceptible to fatigue damage and is, in general, the origin of fatigue crack initiation. It can be shown that, for sharp notches, the maximum-stress range on this element can be related to the stress intensity factor range, ΔKI, as follows: 2 DKI Dsmax ¼ pffiffiffi pffiffiffi ¼ Dsðkt Þ x r
½9:14
where ρ is the notch-tip radius, Δσ is the range of applied nominal stress and kt is the stress concentration factor. pffiffiffi The data show (Fig. 9.13) that DKI = r and, therefore, Δσmax is the primary parameter that governs fatigue crack initiation behavior in regions of stress concentration for a given steel tested in a benign environment. The data also indicate the existence of a fatigue crack initiation threshold, pffiffiffiffiffiffi DKI = rth , below which fatigue cracks would not initiate at the roots of the tested notches. For instance, the value of this threshold is characteristic of the steel and increases with increasing yield or tensile strength of the steel. The data show that the fatigue crack initiation life of a component subjected to a given nominal-stress range increases with increasing strength. Due to the inevitability of cracks (or imperfections) in engineering structures, fatigue crack propagation is important from a designing point of view. This approach attempts to determine the safe load or safe inherent fault dimension that will preclude failure. The fatigue crack propagation behavior of metals is primarily controlled by the stress intensity factor range
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9.13 Fatigue crack initiation behavior of various steels at a stress ratio of +0.1 [20].
ΔK, which can be divided into three regions, as shown in Fig. 9.14. The behavior in region 1 exhibits a fatigue crack propagation threshold, ΔKth, which corresponds to the stress intensity factor range, below which cracks do not propagate under cyclic-stress fluctuations. The fatigue crack propagation threshold for steels is primarily a function of the stress ratio and is essentially independent of chemical or mechanical properties. In 1963, Paris and Erdogan [25] published an analysis with considerable fatigue crack growth rate data and demonstrated that a correlation exists between da/dN and the cyclic stress intensity parameter, ΔK. They argued that ΔK characterizes the magnitude of the fatigue stresses in the crack-tip region; hence, it should characterize the crack growth rate. The data for intermediate fatigue crack growth rate values can be represented by a simple mathematical relationship, commonly known as the Paris equation. This region (Fig. 9.14) represents the fatigue crack propagation behavior above ΔKth (region 2), which can be represented by the power-law relationship: da ¼ CðDKÞn dN
½9:15
where a is the crack length, N is the number of cycles, and C and n are constants. Thus, the fatigue crack growth rate behavior expressed as da/dN versus ΔK can be regarded as a fundamental material property analogous to the yield and ultimate tensile strengths and plane strain fracture toughness, KIC.
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9.14 Schematic illustration of the variation of the fatigue-crack-growth rate, da/dN, with alternating stress intensity, ΔK, in steels, showing regions of primary crack growth mechanisms [20].
The acceleration of fatigue crack growth rates that determines the transition from region 2 to region 3 appears to be caused by the superposition of a brittle or a ductile-tearing mechanism on to the mechanism of cyclic subcritical crack extension, which leaves fatigue striations on the fracture surface. These mechanisms occur when the strain at the tip of the crack reaches a critical value. Thus, the fatigue-rate transition from region 2 to region 3 depends on the maximum stress intensity factor, on the stress ratio and on the fracture properties of the material. Low-cycle fatigue (LCF) conditions are frequently created where the repeated stresses are of thermal origin, which is denominated thermomechanical fatigue (TMF). TMF is a structural failure mode in many hightemperature components. Thermal fatigue loading is induced by temperature gradients during transient heating or cooling from one high temperature of operation to another. Thermal fatigue loading can also occur when heating and cooling are present simultaneously and thermal gradients are maintained during steady-state operation. Internally aircooled high-temperature turbine blades are very representative examples. Thermal gradients produce differential expansion as the hottest material wants to expand more than the cooler, but is constrained from doing so by the cooler and stronger material. The constraint is perceived by the hottest material as a compressive thermal strain that is no different in its effect on the material than would be a mechanically induced strain of equal magnitude. Similarly, the coldest material is forced by the hottest to expand
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9.15 Comparison of the average thermal fatigue lives of conventionally cast, DS cast and SX-cast nickel-base superalloys [13].
more than normal. The thermally induced strain in the colder material is tensile. The analysis of thermal fatigue (Fig. 9.15) is essentially a problem in heat transfer and properties such as modulus of elasticity, coefficient of thermal expansion and thermal conductivity (see Fig. 9.14). The most important metallurgical factors are ductility and toughness. Highly ductile materials tend to be more resistant to thermal fatigue. They also seem more resistant to crack initiation and propagation. As a result of observations of environmental effects, these play a major role in high-temperature fatigue crack growth of superalloys. The presence of sulfur has a significant impact, which provokes profound changes in the material strength. It is also meaningful to remark on the effect of oxygen and the combined effect of oxygen and carbon on the base material. Exposure in air at high temperatures (greater than about 900 8C, or 1650 8F) could lead to profound embrittlement at intermediate temperatures (700 to 800 8C, or 1290 to 1470 8F). With nickel-base superalloys, it has been found that surface cracks related to environmental attack may develop at strains as low as 0.5 %. Since these cracks result in severe loss in fatigue life, this is an appropriate failure criterion rather than rupture life. Gas turbine blades may therefore be designed on the basis of time to 0.5 % creep with a suitable safety factor on stress. We can conclude that in order to develop an improved design methodology for machines and equipment operating at high temperatures,
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several key concepts and their synergism must be considered. Particularly, it must include [26]: . . . . . .
Plastic instability at elevated temperatures, which leads to tertiary creep. Deformation mechanisms and strain components associated with creep processes. Stress and temperature dependence. Fracture at elevated temperatures. Environmental effects. Cycle stress and strain range.
As discussed, the nature of the design process requires serious consideration of the relationships between predicted machine conditions, such as stress, strain and temperature, and the capability of the component materials to withstand those conditions. Engineers will utilize the most appropriate analytical methods and the most precise mechanical and thermal boundary conditions in the design efforts. They will then modify the analytical results by factors of safety, correlations or experience to arrive at the specific for stress and temperature for assessing component life. This value is understood to be a reasonably close and conservative approximation. It is of particular significance that this value is specific, and it becomes the standard against which the design and materials are measured to judge acceptability [2]. On the other hand, engineers have to consider the variability of materials properties. If many tests are run at a specific temperature, a scattering of the property about some mean value is noted. It should also be noted that there is a finite probability (generally greater then 5%) that values for the measured property can fall outside the scatterband of actual data. This characteristic of material properties requires the engineer to determine just what value of the property will be used to judge the acceptability of the design [27]. The nature of superalloys is that they resist the creep-rupture process better than other materials, have very good higher temperature short-time strength (yield, ultimate), very good fatigue properties (including fatigue crack propagation resistance) and combine these mechanical properties with good to exceptional oxidation resistance. Consequently, superalloys are the obvious choice when structures are to operate at higher temperatures. Generally, the temperature range of superalloy operation is broken up into the intermediate range of about 540 8C (1000 8F) to 760 8C (1400 8F) and the high-temperature range that occurs above about 816 8C (1500 8F).
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9.2.2 Creep–fatigue interaction The modes of cracking have frequently exhibited creep-like fractures intermixed with cycle-dependent fatigue-type cracking – hence the descriptive name, creep–fatigue interaction. Creep–fatigue interaction is a special phenomenon that can have a detrimental effect on the performance of metal parts or components operating at elevated temperatures. When temperatures are high enough, time-dependent creep strains as well as cyclic (i.e. fatigue) strains can be present and the interpretation of the effect that one has on the other becomes extremely important. For example, it has been found that creep strains can seriously reduce fatigue life and/or that fatigue strains can seriously reduce creep life [13]. Creep–fatigue interaction testing and modeling have been intense activities due to seemingly premature failures of components in structural equipment operating at elevated temperatures, including gas turbine engines. The interaction between thermally activated time-dependent processes such as creep and mechanical fatigue mechanisms severely complicates life prediction at elevated temperatures. Factors such as frequency, wave shape and creep/relaxation, which are of small consequence at room temperature, take on a significant importance at high temperatures. Hold times at a given stress or strain (e.g. a gas turbine component at constant load) often figure strongly in high-temperature load histories. Under constant stress conditions creep or crack extension may occur, which naturally results in a change in deformation. Under constant strain conditions relaxation may occur, which results in a reduction of the applied stress. Manson early associated the time-dependent fatigue lifetime with intergranular cracking and reasoned that this damage mechanism was intimately associated with time-dependent inelastic strain (i.e. creep or relaxation) whereas time-independent plasticity was accompanied by transgranular cracking. Manson also conducted cyclic creep tests between fixed strain limits and found that the lifetime did not correlate with monotonic time to rupture in a creep test (i.e. t in the creep-rupture test). This led to tests of four simple uniaxial cycle types involving creep and plasticity in the increasing and decreasing halves of the strain cycle. Manson partitioned these strains into four inelastic strain ranges (Fig. 9.16) that may be used as basic building blocks for any conceivable hysteresis loop: Δεpp = tensile plasticity reversed by compressive plasticity Δεcp = tensile creep reversed by compressive plasticity Δεpc = tensile plasticity reversed by compressive creep Δεcc = tensile creep reversed by compressive creep
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9.16
Partitioning of the strain range into different component strains.
Manson generalized this concept into a procedure for evaluating any strain– time–temperature cycle, which he named strain-range partitioning (SRP) [28], which will be discussed in section 9.6.2 on the strain-range partitioning model. Creep–fatigue interaction testing is conducted at an isothermal temperature, sufficiently high that thermally activated, diffusion-controlled creep deformation mechanisms can operate under stress as a function of both time and temperature. The addition of creep to a cycle of normal fatigue loading will invariably reduce the cyclic life, although the clock time to failure may remain constant or actually increase. Conversely, the superposition of fatigue cycling and conventional monotonic creep will also alter the rate of creeping and the time to rupture. The strain-range partitioning technique referred to earlier has been used in the analyses of several high-temperature LCF situations including combustion liners.
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9.2.3 Oxidation and hot corrosion The evolution of conventional superalloys has been dictated by balancing requirements of strength, hot corrosion resistance, oxidation resistance, freedom from deleterious sigma phase and forgeability. As explained, the first requirement is mainly fulfilled thanks to the Ni3 (Al, Ti) phase. However, increasing the volume fraction of this phase required the reduction of chromium and the addition of cobalt. The loss of Cr resulted in both a loss of solid solution strength and oxidation resistance, which were compensated for by the addition of molybdenum and aluminum respectively. The use of Ni-base superalloys as turbine blades in an actual end-use atmosphere produces deterioration of material properties. This deterioration can result from erosion, corrosion or oxidation. Erosion results from hard particles impinging on the turbine blade and removing material from the blade surface. The particles may enter through the turbine inlet or can be loosened scale deposits from within the combustor. Metal oxidation occurs when oxygen atoms react with metal atoms to form oxide scales. The higher the temperature, the more rapidly this process takes place, creating the potential for failure of the component if too much of the substrate material is consumed in the formation of these oxides. At higher temperatures > 899 8C (> 1650 8F), a relatively rapid oxidation attack of some materials can occur unless there is a barrier to oxygen diffusion on the metal surface. Additionally, oxygen may also penetrate along grain boundaries at high temperatures, causing a rapid decrease of material strength. Aluminum oxide (Al2O3) provides such a barrier. Aluminum oxide will form on the surface of a superalloy at high temperatures if the superalloy’s aluminum content is sufficiently high. Thus, the alloy forms its own protective barrier in the early stages of oxidation by the creation of a dense, adherent aluminum oxide scale. Hot corrosion is a rapid form of attack that is generally associated with alkali metal contaminants, such as sodium and potassium, reacting with sulfur in the fuel to form molten sulfates. Hot corrosion is an accelerated oxidation of alloys caused by the deposition of Na2SO4. Hot corrosion results from the ingestion of salts in the engine and sulfur from the combustion of fuel. The presence of only a few parts per million (ppm) of such contaminants in the fuel, or equivalent in the air, is sufficient to cause this corrosion. Sodium can be introduced in a number of ways, such as salt water in liquid fuel, through the turbine air inlet at sites near salt water or other contaminated areas, or as contaminants in water/steam injections. As well as alkali metals such as sodium and potassium, other chemical elements can influence or cause corrosion on hot gas components. Notable in this connection are vanadium, primarily found in crude and residual oils, and
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9.17 Schematic illustration of the variation in corrosion rate with temperature [29].
lead. Corrosion causes deterioration of blade materials and reduces component life. There are now two distinct forms of hot corrosion recognized by the industry (Fig. 9.17), although the end result is the same. These two types are high-temperature (type I) occurring at 816–927 8C (1500–1700 8F) and lowtemperature (type II) hot corrosion occurring at 593–760 8C (1100–1400 8F), both requiring a higher chromium to aluminum ratio and substitution of molybdenum by other refractory elements such as W, Ta and Nb. High-temperature hot corrosion has been known since the 1950s. It is an extremely rapid form of oxidation that takes place in the presence of sodium sulfate (Na2SO4). Sodium sulfate is generated in the combustion process as a result of the reaction between sodium, sulfur and oxygen. Sulfur is present as a natural contaminant in the fuel. Low-temperature hot corrosion was recognized as a separate mechanism of corrosion attack in the mid 1970s. This attack can be very aggressive if the conditions are right. It takes place at significant partial pressure of SO2. It is caused by low melting eutectic compounds resulting from the combination of sodium sulfate and some of the alloy constituents such as nickel and cobalt.
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The lines of defense against both types of corrosion are similar. First, reduce the contaminants. Second, use materials that are as corrosionresistant as possible. Third, apply coatings to improve the corrosion resistance of the component alloy. A high-nickel alloy is used for increased strength at elevated temperatures and a chromium content in excess of 20 % is desired for corrosion resistance. An optimum composition to satisfy the interaction of stress, temperature and corrosion has not yet been developed. Thus, many highstrength superalloys in use today cannot form sufficient protective scales because the compositional requirements for achieving other properties, such as high strength and metallurgical stability, do not allow for the optimization of oxidation/corrosion resistance in the superalloy itself. Therefore, most of today’s superalloys must receive their oxidation protection from specially engineered coatings.
9.2.4 Properties of superalloys Superalloys were initially developed for use in aircraft piston engine turbosuperchargers, and their development over the last 60 years has been paced by the demands of advancing gas turbine engine technology. The design for high-temperature applications requires specific material characteristics based on the main failures modes affecting its intended function. For gas turbine applications, the performance characteristics are limited by the operating conditions that can be tolerated by the material used. Superalloys possess a remarkable ability to maintain their properties at high temperature and this will be briefly discussed. As defined, the basis of superalloys are iron, cobalt and nickel, i.e. transition metals located in a similar area of the periodic table of the elements in the 8th group of the periodic system of the elements. Table 9.4 shows the physical properties of the superalloy base elements. It can be seen that pure iron has a density of 7.87 g/cm3 (0.284 lb/in3), while pure nickel and cobalt have densities of about 8.9 g/cm3 (0.322 lb/ in3). The superalloys are created usually by adding significant levels of the alloy elements chromium, aluminum and titanium, plus appropriate refractory metal elements such as tungsten and molybdenum to the base metal. Densities of superalloys are a function of the amounts of these elements in the final compositions. Aluminum, titanium and chromium reduce superalloy density whereas the refractory elements such as tungsten, rhenium and tantalum increase it. Table 9.5 gives the density, melting range and physical properties of some nickel-base and cobalt-base superalloys [31]. The main properties of superalloys are that they exhibit some combination of high strength at temperature; resistance to environmental attack (including nitridation, carbonization, oxidation and sulfidation); excellent
© Woodhead Publishing Limited, 2011
hcp fcc bcc
© Woodhead Publishing Limited, 2011
8.00
b
1493 1452 1535
0.32 0.32 0.28
8.9 8.9 7.87
7.0 7.4 6.7
— 1305– 1420
1260– 1290 1290– 1320
— 2380– 2590
2300– 2350 2350– 2410
8F
At 538 8C (1000 8F)
At 1093 8C (2000 8F) At 93 8C (200 8F)
At 538 8C (1000 8F)
Thermal conductivity
12.4 13.3 11.7
At 1093 8C (2000 8F)
464 610 493
Mean coefficient of thermal expansion (106/K)a
0.215 0.165 0.175
cal/cm2/s/8C/cm
Btu/ft2/h/8F/in
8C6106
8F6106
— 420
440
420
— 0.10
0.105
0.10
— 530
565
565
— 0.126
0.135
0.135
— 645
710
710
— 0.154
0.17
0.17
— 11.0
10.7
10.9
— 76
74
76
— 21.8
16.7
17.0
— 151
116
118
— 32.1
25.3
26.4
— 222
176
183
— 12.2
10.1
10.6
— 14.4
15.8
13.5
— —
18.9
17.1
J/kg K Btu/lb 8F J/kg K Btu/lb 8F J/kg K Btu/lb 8F W/m K Btu in/h ft2-8F W/m K Btu in/h ft2-8F W/m K Btu in/h ft2-8F At 938C At 5388C At 10938C (2008F) (10008F) (20008F)
At 21 8C (70 8F)
Specific heat
From room temperature to indicated temperature. Liquidus temperature.
Cobalt base FSX-414 8.3 Haynes 1002 8.75
7.91
8C
Density Melting range (g/cm3)
IN-713 LC
a
2723 2647 2798
g/cm3
Thermal conductivitya
Expansion coefficienta
Physical properties of cast nickel-base and cobalt-base alloys [15]
Nickel base IN-713 C
Alloy
Table 9.5
a At room temperature.
Co Ni Fe
lb/in3
8F
8C
Density
Melting point
Some physical properties of superalloy base elements [30]
Crystal structure
Table 9.4
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creep resistance, stress-rupture strength toughness and metallurgical stability; useful thermal expansion characteristics and resistance to thermal fatigue and corrosion. The influence of temperature on the strength has been discussed and can be demonstrated by running standard, short time tests at a series of increasing temperatures. This leads to the conclusion that the melting temperature of the material is a critical parameter for hightemperature behavior. Thus the first required characteristic is the ability to withstand loading at an operating temperature close to its melting point. If the operating temperature is denoted Top and the meeting point Tm, a criterion defined based upon the homologous temperature defined as Top/ Tm is important in material selection, which should be greater than about 0.6. The melting temperatures of the base superalloy elements are nickel at 1452 8C (2647 8F), cobalt at 1493 8C (2723 8F) and iron at 1535 8C (2798 8F). When metals are alloyed, there is no longer a single melting point for a composition. Instead, alloys melt over a range of temperatures. The lowest melting temperature (incipient melting temperature) and melting ranges of superalloys are functions of the composition and prior processing. Just as the base metal is higher melting, so generally are incipient melting temperatures higher for cobalt-base superalloys than for nickel-base or iron–nickel-base superalloys. Nickel-base superalloys may show incipient melting at temperatures as low as 1204 8C (2200 8F). However, advanced nickel-base single-crystal superalloys having limited amounts of melting point depressants tend to have incipient melting temperatures equal to or in excess of those of cobalt-base superalloys [31]. Other physical properties such as electrical conductivity, thermal conductivity and thermal expansion of superalloys tend to be low (relative to other metal systems). These properties are influenced by the nature of the base metals (transition elements) and the presence of refractory-metal additions. A second characteristic is a substantial resistance to mechanical degradation over extended periods of time. As discussed, time-dependent deformation and fracture of structural materials at elevated temperatures are among the most challenging engineering problems. In order to develop an improved design methodology for machines and equipment operating at high temperatures, key aspects to be considered are plastic instability at elevated temperatures and deformation mechanisms and strain components associated with creep processes. The superalloys have low ductility compared to iron-based steels; the ductilities of cobalt-base superalloys are generally less than those of iron– nickel-base and nickel-base superalloys. Short-time tensile ductilities as determined by elongation at failure generally range from as low as 10 pct to
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as high as around 70 pct, but γ´ hardened alloys are in the lower end, usually between about 10 and 40 pct [31]. Creep-rupture ductilities are generally lower than tensile ductilities. At the 760 8C (1400 8F) tensile ductility minimum area, creep-rupture ductilities of castings have gone below 1.5 pct; however, most current high-strength polycrystalline (PC) equiaxed cast alloys have rupture ductilities in excess of 2.0 pct. Single-crystal directionally solidified (SCDS) superalloy ductilities will vary with orientation of the single crystal relative to the testing direction [31]. Superalloys typically have dynamic moduli of elasticity in the vicinity of 207 GPa (306106 psi), although moduli of specific PC equiaxed alloys can vary from 172 to 241 GPa (25 to 356106 psi) at room temperature depending on the alloy system. Processing that leads to directional grain or crystal orientation can result in moduli of about 124 to 310 GPa (about 18 to 456106 psi) depending on the relation of grain or crystal orientation to the testing direction [31]. Short-time tensile yield properties of γ´-hardened alloys range from around 550 MPa (80 ksi) to 1380 MPa (200 ksi) at room temperature. Actual values depend on the composition and processing (cast versus wrought). Ultimate strengths range from around 690 MPa (100 ksi) to 1520 MPa (230 ksi) at room temperature, with γ´-hardened alloys in the high end of the range [31]. Superalloys tend to show an increase of yield strength from room temperature up to about 760 8C (1400 8F) and drop off thereafter. This is in contrast to ordinary alloys that tend to continuously decrease in short-time strength as temperatures increase. Ultimate tensile strengths generally do not show this trend. Concurrently, tensile ductility tends to decrease, with a minimum at around 649 8C (1200 8F). The highest tensile properties are found in the finer grain size wrought or powder metallurgy superalloys used in applications at the upper end of the intermediate temperature regime, perhaps to about 760 8C (1400 8F). The highest creep-rupture properties invariably are found in the coarser grain cast superalloys used in the high-temperature regime. Rupture strengths are a function of the time at which they are to be recorded. The 1000 h rupture stress capability is obviously lower than the 100 h capability. Creep-rupture strengths for 100 h failure at 982 8C (1800 8F) may range from 45 MPa (6.5 ksi) for an older γ´ hardened wrought alloy such as U500 to 205 MPa (30 ksi) for the PC equiaxed cast superalloy Mar-M 246. Columnar grain and single-crystal alloys can be much stronger [31]. A third characteristic is the ability to withstand cyclic mechanical and thermal-induced loading. This is known as fatigue failure due to the sudden and catastrophic separation of a machine part into two or more pieces as a result of the application of fluctuating loads or deformations over a period
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of time. Failure takes place by the initiation and propagation of a crack until it becomes unstable and propagates suddenly to failure. The loads and deformations that typically cause failure by fatigue are far below the static or monotonic failure levels. When loads or deformations are of such magnitude that more than about 10 000 cycles are required to produce failure, the phenomenon is usually termed high-cycle fatigue. When loads or deformations are of such magnitude that less than about 10 000 cycles are required to produce failure, the phenomenon is usually termed low-cycle fatigue. When load or strain cycling is produced by a fluctuating temperature field in the machine part, the process is usually termed thermal fatigue. Thermal fatigue is the single most frequent cause of machine repair or failure, and understanding it requires substantial analytical, experimental and metallurgical effort. Cyclic properties are not commonly tabulated for superalloys. Properties of interest would be the 103–105 and 106–108 cycle fatigue strength capabilities. This could mean stress for a fixed-cyclic life-toa-particular-sized crack or stress for a fixed-cyclic life-to-fracture for LCF (low-cycle fatigue) regimes or only stress for a fixed-cyclic life-to-fracture for HCF (high-cycle fatigue) regimes. Also, crack propagation rates versus toughness parameter (da/dN versus ΔK) are desired. The life values, when available, lend themselves to tabulation, but the da/dN values are best represented by graphs. LCF strengths are usually related to an alloy’s yield strength while HCF strengths are usually related to an alloy’s ultimate strength for wrought alloys used at intermediate temperatures. For cast alloys used in the hottest sections of a gas turbine, there appears to be a relation of TMF strength to the creep strength of an alloy for a given alloy form, as there is for columnar grain directionally solidified nickel-base superalloys [13]. The interaction of creep and fatigue is probably synergistic and leads to combined creep and fatigue failure, in which all of the conditions for both creep failure and fatigue exist simultaneously, each process influencing the other to produce failure. This is an essential criterion for remaining life evaluation. A fourth characteristic is tolerance of severe operating environments, in which the role of the ceramic and metallic coating applied will play a critical role. On the other hand, coatings have their own failures modes. Metal coatings have been designed to withstand three types of environmental attack: high-temperature oxidation, high-temperature hot corrosion (type I) and low-temperature hot corrosion (type II). However, the coatings degrade at two fronts during service: the coating/gas-path interface and the coating/ substrate interface. They can be summarized as follows: hot corrosion, mechanical distress, solid-state diffusion, spallation and thermomechanical fatigue cracking.
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9.3
Protective coatings
Both metallic and ceramic coatings have become essential to withstand current operational conditions of gas turbines. The metallic coatings of blades and vanes are required to protect the component from corrosion, oxidation and mechanical property degradation by corrosive attack. As superalloys have become more complex, it has been increasingly difficult to obtain both the higher strength levels that are required and a satisfactory level of corrosion and oxidation resistance without the use of coatings. Thus, the trend toward higher firing temperatures increases the need for coatings. This trend has resulted in environmental degradation of materials, deteriorating the mechanical properties and shortening the service life of components. The idea to apply a layer with protective properties on the surface of Nibased superalloys was first practiced in the 1960s. Two types of protective coatings have been most widely used: diffusion aluminide coatings based on the formation of a β-NiAl phase at the component surface and application of MCrAlY (M = Ni, Co or NiCo) overlay coatings based on a mixture of β-NiAl and γ´-Ni3Al or γ phases. The function of these coatings is to provide a surface reservoir of elements such as Al and Cr that will form stable, adherent oxide layers that will protect the substrate alloy from environmental attack. As the temperature capability of Ni-based superalloys approaches their intrinsic limit, further improvements in their temperature capability have become increasingly difficult. With turbine firing temperatures exceeding 1120 8C (2050 8F) and thereby challenging the ability of air cooling to maintain surface temperatures below about 900 8C (1650 8F), there has been very considerable growth in the use of ceramic coatings, i.e. thermal barrier coatings. The function of the thermal barrier coating (TBC) is to retard the conduction of heat from the combustion gas to the internal cooling air, and thus either reduce the amount of cooling air needed or allow an increase in the turbine inlet temperature for a given metal temperature. Therefore, during the past two decades, the emphasis in gas turbine materials developments has shifted to TBCs, which are ceramic coatings with a very low thermal conductivity that reduce the alloy surface temperature by insulating it from the hot gas. Current state-of-the-art thermal barrier coatings comprise two layers: a diffusion aluminide or MCrAlY bond coat and a low thermal conductivity partially stabilized zirconia (YSZ: 7 to 8 wt% Y2O3–ZrO2) top coat. Thermal barrier coatings were first successfully tested in a research turbine engine in the mid 1970s. By the early 1980s they entered revenue service on the vane platforms of aircraft engines, and today they are flying in revenue service on vane and
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blade surfaces, and combustor section components such as transition pieces and liners. Thermal barrier coatings are expected to play an increasingly significant role in advanced gas turbine engines both in aero and industrial applications in the future.
9.3.1 Diffusion coatings Diffusion coatings (Table 9.6) are based on the formation of intermetallic compounds such as β-NiAl and β-CoAl via a diffusion process. Their usefulness comes from the protective nature of an Al2O3 scale that forms on the coated part at the service operating temperature. Developed in the 1950s, diffusion coatings still are widely used. The majority of these coatings are manufactured by means of slurry fusion, pack cementation and related gas-phase (out-of-contact) processes. Aluminide coatings degrade in service through cyclic oxidation, hot corrosion, erosion, interdiffusion and thermomechanical fatigue cracking. Aluminum in the coating combines with oxygen at the substrate surface, forming a protective Al2O3 scale. When the scale cracks and spall from thermal cycling, additional aluminum from the coating diffuses to the surface to reform the scale. Aluminum from the coating also diffuses into the substrate, and as aluminum is depleted in the coating, β-NiAl converts to γ´-Ni3Al and then to γ-Ni solid solution. When the aluminum content in the coating drops to about 4 to 5 wt%, the Al2O3 scale can no longer form and rapid oxidation occurs. The rate of aluminum diffusion is influenced by the substrate composition. If the coating undergoes thermal fatigue cracking, refractory elements in the diffusion zone may be exposed to the oxidizing environment and oxidize rapidly. This effect must be considered when selecting a coating for use in cyclic conditions. Incipient melting in the diffusion zone also can result in rapid oxidation penetration. Such melting can occur at temperatures as low as 1120 8C (2050 8F), well below the melting point of NiAl (1590 8C, or 2900 8F). It is well established that platinum in diffusion aluminides extends the lives of such coatings in oxidizing environments. Chromide diffusion coatings rapidly form a continuous, adherent Cr2O3 scale for protection. Because interdiffusion is small, in lower-temperature environments, chromide coatings are relatively thin (0.04 to 0.05 mm, or 1.5 to 2.0 mils). The thinner coating is more favorable with respect to mechanical properties, because chromium compositions are prone to cracking due to lower ductility. Chromizing or chromizing followed by aluminiding of high-temperature alloys is used to improve resistance to hot corrosion and high-temperature oxidation. Chromium is added either through codeposition with aluminum or more effectively in a two-step pack
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Table 9.6
Commercial diffusion coatings [32]
Coating designation
Producer
Composition
PWA 73 PWA 273
Pratt & Whitney Pratt & Whitney
Aluminide Aluminide
Comments
High-activity pack Gas-phase Cr, low-activity pack PWA 275 Pratt & Whitney Aluminide Low-activity above-the pack PWA 70 – 73 Pratt & Whitney Chromized + aluminized High-activity pack MDC 200 Howmet Aluminide High activity MDC 210 Howmet Aluminide Low activity RT 21 Chromalloy Aluminide High activity RT 22 Chromalloy Pt aluminide High-activity pack RT 44 Chromalloy Pt – Rh aluminide High-activity pack (Co base) JML 1 Johnson Matthey Pt aluminide (fused High-activity pack salt deposition + aluminized) SS 82A Turbine Component Pt aluminide High-activity Corp. above-the-pack LDC 2E Howmet Pt aluminide High-activity pack MDC 3V Howmet Chromized High-activity pack MDC 150 Howmet Pt aluminide High-activity pack MDC 150L Howmet Pt aluminide High-activity CVD MDC 151L Howmet Pt aluminide + Hf, Si, Zr Low-activity CVD SermaTel W Sermatech Slurry aluminide Low-activity CVD SermaLoy J Sermatech Slurry aluminide with Si High-activity pack Elcoat 360 Elbar Slurry Ti with Si High-activity pack Jo Coat Pratt & Whitney Slurry aluminide with Si High-activity pack CODEP General Electric Aluminide High-activity pack ALPAK Rolls Royce Aluminide High-activity pack SermAlcote Sermatech Pt aluminide High-activity pack Snecma CIA SNECMA Chromized + aluminized Plating + slurry + diffuse
process in which pack or gas phase chromizing is followed by pack aluminiding. Silicon is beneficial in improving both resistances to oxidation as well as type II hot corrosion. It can be incorporated as a silicide coating by a pack process with metallic Si as the source. It can be introduced during or prior to the aluminiding process.
9.3.2 Overlay coatings Overlay coatings (Table 9.7) differ from diffusion coatings in that interdiffusion of the coating with the substrate is not required to generate the desired coating structure/composition. Overlay coatings do not rely on
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Gas turbine materials selection Table 9.7
371
Commercial overlay coatings [32]
Coating designation
Producer
PWA 286 PWA 270 BC 21 (Marine) BC 22 BC 23 (Marine) three-layer
Pratt & Whitney NiCoCrAlY + Hf, Si Pratt & Whitney NiCoCrAlY GE Co22.5Cr10.5Al0.3Y GE Co26Cr10.5Al2.5Hf GE Co26Cr12Al1Hf5Pt
BC23 (Marine) BC51 BC52 GT29 (Plasmaguard) GT29+ GT23 SICOAT 2231 SICOAT 2442 SICOAT 2453 SICOAT 2412 SDP 2
GE GE GE GE
Co26Cr12Al1Hf5Pt Ni4CoCr6Al0.3Y CoCrAlY CoCrAlY + aluminized
LPPS EB-PVD PVD LPPS PVD + aluminize + Pt + diffusion heat treatment LPPS LPPS LPPS LPPS
GE GE Siemens AG Siemens AG Siemens AG Siemens AG N K Engines, Moscow Sermatech Allied Signal
NiCoCrAlY CoNiCrAlY NiCoCrAlY Ni10Co23Cr12Al13Re0.6Y CoNiCrAlYRe NiCrAlY Overaluminized
LPPS + aluminized LPPS LPPS LPPS LPPS/HOVF LPPS EB-PVD
MCrAlY NiCrAlY
LPPS + pack Overlay
SermAlcote SCC-103
Composition
Comments
reaction with the substrate for their formation; instead, a prealloyed material applied over the substrate determines the coating composition and microstructure, and adhesion of the coating to the substrate is effected by some elemental interdiffusion. Coatings of this type in current use are generally called MCrAlY overlay coatings, where M represents Ni, Co, Fe or some combination of these metals. These coatings essentially comprise a monoaluminide (MAl) component contained in a more ductile matrix of solid solution (γ), in the case of CoCrAlY coatings, or a mixture of γ and γ´-Ni3Al phase, in the case of NiCrAlY coatings; NiCoCrAlY and FeCrAlY modifications are also available. The oxidation pattern of MCrAIY overlay coatings is similar to that of diffusion aluminide coatings. Elements such as chromium and yttrium improve the resistance of the Al2O3 scale to spallation. The most widely used overlay coating for oxidation resistance is the NiCoCrAIY type. The addition of cobalt also improves coating ductility. In the oxidation process, grains of aluminum-rich β phase convert to islands of γ´ eventually leaving only the less oxidation-resistant γ matrix phase. Substrate composition can influence oxidation resistance. MCrAlY overlay coatings have a higher melting point than diffusion
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coatings, so melting does not occur in the interdiffusion zone at temperatures lower than the melting point of the bulk coating. Coatings have survived exposure temperatures as high as 1290 8C (2350 8F) without melting. The addition of elements such as silicon, tantalum and hafnium can improve oxidation resistance at the expense of some ductility. Compositional flexibility offers the opportunity to tailor coatings for optimum performance.
9.3.3 Thermal barrier coatings The hot corrosion (and oxidation) resistance of coated blades can be further improved by applying a layer of thermal insulation. This thermal barrier coating (TBC) must be sufficiently thick, have a low thermal conductivity and high thermal-shock resistance, and have a high concentration of internal voids to reduce further thermal conductivity to a value well below that of the bulk material. The temperature difference between the outer surface of a TBC and the outer surface of the underlying corrosion-resistant film can be as high as 150 8C (270 8F). In addition to reducing the temperature at the surface of the superalloy, these coatings also reduce thermal-shock loads on the blades; it means rapid changes in ambient temperature are moderated and attenuated before they reach the substrate. A TBC system consists of an insulating ceramic outer layer (top coat) and a metallic inner layer (bond coat) between the ceramic and the substrate. Both the top coat and bond coat can be applied by plasma spraying (air and low pressure) and EB-PVD (electron-beam physical vapor deposition). A schematic of a typical system is shown in Fig. 9.18. Current state-of-the-art TBCs are zirconium oxide (ZrO2), zirconia with 6 to 8% (by weight) of yttrium oxide (Y2O3), or yttria, to partially stabilize the tetragonal phase for good strength fracture toughness, and resistance to thermal cycling. The coatings are relatively inert, have a high melting point and have low thermal conductivity. Thermal barrier coatings are currently used in the combustion section (combustor cans/liners and transitions) of some gas turbines, including the newer models from main OEMs, which also began using TBCs on the first stage stationary vane and continued this application through the 1425 8C (26008F) model, which have TBCs applied to the first and second stage blades and vanes. Recent development in electron-beam physical vapor deposition of TBCs has resulted in potential improvements over the air plasma spray process for the ceramic layer. This is expected to result in the wider use on the airfoils of vanes and blades. Key factors affecting TBC performance are bond coat surface finish, bond coat oxidation, bond coat surface imperfections, thermal conductivity of YSZ, sintering of YSZ and phase transformation of YSZ. Other problems
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9.18 Schematic of a multilayer thermal barrier coating system [13].
must be overcome to make a good TBC. It must stick to the blade, and that is not easy. To achieve it, the blade surface is first plated with a thin bond coat (a complex Ni–Cr–Al–Y alloy) between the blade and coating. Most ceramics have a lower expansion coefficient than the superalloy of the blade, so when the blade heats up it expands more than the coating – we know what that means: thermal stresses and cracking. The problem is solved by arranging for the coating to be already cracked on a fine scale, with all the cracks running perpendicular to its surface, making an array of interlocking columns. When the blade expands the columns separate very slightly, but not enough for hot gas to penetrate to any significant extent; its protective thermal qualities remain.
9.3.4 Failure modes Once it became apparent that it would be extremely difficult to develop structural materials that possessed all the desired high-temperature mechanical properties as well as environmental resistance, the functions were uncoupled [33]. Structural alloys were developed by Stringer and Viswanathan [34] to optimize mechanical properties and coatings were developed to serve as physical barriers between aggressive environments and the substrate. The structural materials degradation modes that are moderated by the traditional metallic protective coatings include oxidation and hot corrosion. In addition, TBCs insulate the underlying structure against the full effect of gas-path heat. They thus retard creep degradation and reduce the severity of the thermal gradients and transients in the structure that drives low-cycle fatigue processes [33]. Historically, metal coatings (i.e. alumide, chromide and MCrAlY) have been designed to withstand three types of environmental attack: hightemperature oxidation, high-temperature (type I) hot corrosion and low-
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temperature (type II) hot corrosion. Figure 9.17 shows the range of temperatures over which these attacks occur. However, it should be kept in mind that the oxidation of the aluminide and chromide coatings at the coating/gas-path interface results in the formation of a protective oxide scale. Therefore, within this context, high-temperature oxidation is not purely a degradation mechanism. In addition to these failure modes, thermomechanical fatigue of coatings (and substrates) can occur as a result of cyclic and thermal loading of the component. Thermomechanical fatigue crack loading in coatings, particularly around film-cooling holes, has often been observed in advanced engines. A TBC can reduce the magnitude of the thermomechanical fatigue strain range and can minimize or eliminate thermomechanical fatigue cracking. Several other damage modes cause coating loss and accelerate the overall failure mechanism, such as mechanical distress to the coating, solid-state diffusion of elements between the coating and substrate and spallation caused by differential thermal expansion between the coating and the substrate. The coating degrades at two fronts during service: the coating/gas-path interface and the coating/substrate interface. At temperatures well below the incipient melting point of coating of the conventional superalloys, deterioration of the coating surface at the coating/gas-path interface tends to be a consequence of oxidation or hot corrosion. As the temperature rises, diffusion across the coating/substrate interface plays a greater role in degradation; compositional changes at this interface can also compromise the structural properties of the substrate. High-temperature oxidation HTO refers to a solid–gas chemical reaction that produces the oxide(s) of constituents within the solid. The rate of oxidation increases exponentially with temperature; certain oxides (notably those of aluminum and chromium) are slow growing and protective of the underlying substrate. At high temperatures, coatings that protect against oxidation form a compact, adherent oxide scale (usually Al2O3), which provides a barrier between the high-temperature gases and the underlying metal. Chromia scales (Cr2O3) have been used but offer less protection than alumina above 840–870 8C (1550–1600 8F) because chromia scale tends to sublimate to CrO3 above these temperatures. Without the protective scale, the coating, and ultimately the substrate, come under rapid attack. The general case is that the oxide spalls repeatedly until it can no longer form. Then internal oxidation, caused by diffusion of oxygen into the coating, degrades the protective oxide scale. Segregation at the metal/scale interface causes disruption of the bond and spallation of the oxide. In
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addition, oxygen-active elements such as hafnium and yttria are often added to coatings to improve the adhesion of the oxide. Hot corrosion Corrosion of materials including Ni-base and Co-base superalloys induced by molten salts in an oxidizing gas elevated temperature is termed ‘hot corrosion’. It is the result of accelerated oxidation at temperatures typically between 700 8C (1300 8F) and 925 8C (1700 8F) when metals and alloys become covered with contaminant salt films. Hot corrosion involves attack by molten salts, typically sodium and potassium sulfates, that enter the turbine hot section as contaminants from the air and fuel and can result in rapid loss of material (Stringer and Viswanathan [35]). As indicated in Fig. 9.3, there are two forms of hot corrosion: high-temperature (type I) and low-temperature (type II). The key to continued protection against hot corrosion is maintaining the protective scale and keeping the hot corrosion process in its initial, or incubation, period. Rapid loss of the coating and the underlying metal substrate will occur once the scale is lost or penetrated. The ideal protective scale would minimize solubility of the fused salt, a property determined by the combustion chemistry and operational environment. This criterion for petroleum-base fuels can generally be satisfied by either chromia or alumina, although alumina protects at higher operating temperatures than chromia. A TBC degradation mechanism has been observed in service in which molten surface deposits can accelerate spalling of electron-beam physical vapor deposition (EB-PVD) and plasma-sprayed TBCs. Salt species in vapors from such sources as dirty fuel or marine environments can condense on to the coating surface. If the coating temperature is above the melting point of the salt, it can wick through the interconnected cracks and porosity of the coating (Miller [36]). The wicked-in salt freezes on cooling, eliminating the strain tolerance mechanism of the coating and accelerating spallation (Palko et al., [37], Miller [36] and Strangman [38]). Mechanical distress Erosion and impact damage caused by the ingestion of articles in the air stream leads to mechanical distress. Two related failure modes – mechanically induced erosion and impact damages – merit special consideration. Erosion results from particles in the air stream that abrade the coating and accelerate its loss. Impact damage occurs when a solid object in the air stream strikes the coating. These failure modes are sometimes overlooked because they can often be avoided by proper operating practice. However, these modes will become increasingly
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important as substrate materials are operated under conditions in which they offer little inherent protection against oxidation or other major failure modes. These failure modes are also a key concern in the retention of TBCs. These ceramic coatings are inherently more vulnerable to the impact damage and erosion than metallic coatings. Creep of standalone metallic coatings at higher temperatures can cause rumpling of the surface of coated components, resulting in degradation of performance and durability. Application of a TBC may alleviate the problem by lowering the temperature of the metallic coating (Bose and DeMasi-Marcin [39]). However, plastic instability of bondcoats for TBCs is still a potential problem at higher temperatures, and work on creep-resistant bondcoats is continuing (Brindley [40]). Solid-state diffusion Solid-state diffusion leads to reduction of the aluminum content of the coating because of interdiffusion with the base metal. The loss of aluminum in the coating, caused by diffusion of the coating with the base metal, ranks as another important degradation mode. Diffusion of aluminum into the base metal and base metal elements into the coating reduce the concentration of aluminum in the coating that is available for forming alumina. A protective oxide can no longer re-form after spallation once the aluminum concentrations fall below a certain level. Spallation Spallation is the loss of protective oxide at the coating/oxide interface. Loss of oxide scale by spallation is the major concern for all coatings and poses a particular problem for coating based on aluminum or chromium. Once again, a protective scale can no longer be maintained when the aluminum (or chromium) concentration falls below a critical level, usually cited as approximately 4 to 5 wt % aluminum. Loss of the protective oxide at the coating/oxide interface is most damaging since an entirely new protective scale must be formed. Recent research indicates that spallation at this interface can be exacerbated by the presence of sulfur in the material. With the reduction of sulfur impurities, stepwise improvements have been observed in the adherence of protective oxides to superalloys (Smialek [41] and Smialek and Tubbs [42]), resulting in vastly improved oxidation resistance. TBC thickness is yet another concern. While it is obvious that a finite amount of TBC must be deposited in order to reap the benefit of a thermal barrier, increasing the TBC thickness adds to component weight and accelerates spallation (Bose and DeMasi-Marcin [39]).
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9.19 Failure mechanism in a thermal barrier coating (TBC) [13].
Spallation of TBCs is an important failure mode. These ceramic coatings are not self-replenishing. Hence, spallation rapidly degrades the insulating properties of the coating and can accelerate attack of the underlying metallic bondcoat or substrate. Thermomechanical fatigue cracking Thermomechanical fatigue (TMF) is the long-term (i.e. over many cycles) formation and propagation of cracks because of external mechanical stresses and residual stresses from lack of thermal expansion compatibility between the substrate and the coating. Failure of TBCs in service is generally attributed to stress that develops during cooling after high-temperature exposure and to transient thermal stress that develops during rapid thermal cycling. Failure occurs primarily due to thermal expansion mismatch between the ceramic and metallic layers and environmental attack or the bondcoat. The stress state in the ceramic layer is biaxial compressive in the plane of the coating. Strains induced by these stresses increase during repeated thermal cycling, which results in crack initiation and eventual spalling of the coating, as shown in Fig. 9.19. The use of Y2O3 to stabilize the ceramic coat, together with an MCrAlY-type bondcoat, significantly improves the thermal fatigue resistance of TBCs.
9.4
Material applications
Table 9.8 shows the chronological evolution of large land-based GTs. During the last 40 years the power ratings of GTs have increased from 50 to nearly 340 MW and the thermal efficiencies have increased from 29 % to 44 % in simple cycle and from 43 to 60 % in combined cycle configurations. These changes have been largely possible because of increases in the turbine inlet temperature from 900 to 1426 8C (1650 to 2600 8F), the compressor pressure ratio from 10.5 to 23 and also in air mass flow.
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46
43
49
70–105 34
R1 and 2 vane, R1 and 2 blade
1120 (2050) 14 530 (986)
1979
53
147–240 36
R1, 2, 3 vane, R1, 2, 3 blade
1260 (2300) 14.5 582 (1080)
1990a
58
167–271 39
R1, 2, 3 vane, R1, 2, 3 blade
1417 (2583) 19–23 593 (1100)
2000b
60
R1, 2, 3 vane, R1, 2, 3 blade, steam-cooled concept 244–340 40
1425 (2600) 19–32 625 (1157)
2010c
Corresponds approximately to GE 7F, FA/9F, FA, S-W 501F and MHI 501F/701F conditions; Siemens V84.3/94.3 is in the same class, in terms of output and heat rate. b Corresponds approximately to S-W 501G, MHI 701G1 conditions; GE 7FA+/9FA+, Siemens V84.3A/94. A and Alstom (ABB) GT2/26 are in the same class in terms of output and heat rate. c Corresponds approximately to MHI 701G2 and S-W SGT5-8000H. GE 7H/9H are in the same class in terms of output and heat rate. Note: SC and CC denote simple cycle and combined cycle. RIT is rotor inlet temperature. Source: Based on adaptation of reference [43]
a
60–80 31
50–60 29
R1 and 2 vane, R1 blade
R1 vane
SC power range (MW) Approx. SC efficiency (%) Approx. CC efficiency (%)
1010 (1850) 11 482 (900)
900 (1650) 10.5 427 (800)
Approx. RIT (8C (8F)) Compressor ratio Exhaust temperature (8C (8F)) Cooled turbine rows
1972
1967
Approx. 1 per year
Table 9.8 Evolution of large land-based gas turbine design features
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In models developed during the last decade gas turbine manufacturers have concentrated their efforts on improving gas turbine efficiency, emissions, operating flexibility and introducing novel designs in a race to dominate the power market: .
Siemens Power Generation has developed the new generation of H-class gas turbines that set unparalleled standards for high efficiency, low life-cycle costs and operating flexibility. Siemens’ latest gas turbine for the 50 Hz market, the SGT5-8000H gas turbine, is the world’s largest exclusively air-cooled gas turbine with the following main features: – – – – – – –
Gas turbine output net: 340 MW Pressure ratio: 19.2:1 Gas turbine efficiency: 40 % Exhaust mass flow: 820 kg/s Exhaust temperature: 625 8C (1157 8F) Hydraulic clearance optimization (HCO) Four-stage turbine with stage 1 single-crystal blades and vanes, stage 1 to 3 TBC coated blades and vanes, air-cooled – High-temperature can-annular combustion system, air-cooled.
This is the first new frame developed after the merger of Siemens and Westinghouse and combines the best features of the existing product lines and advanced innovative technology. Siemens is installing the world’s highest output gas turbine at E.ON Energie’s Irsching site in Bavaria, Germany. Irsching 4 is a turnkey combined cycle power plant (SCC5-8000H) around the turbine with 570 MW power output in combined cycle operation with an efficiency of over 60 %. .
GE continues to develop and market the H system, which has steamcooled stationary and rotating parts. The GE H System™ gas turbine uses closed-loop steam cooling of the turbine. This unique cooling system allows the turbine to fire at a higher temperature for increased performance, yet without increased combustion temperatures or their resulting increased emissions levels. It is this closed-loop steam cooling that enabled the combined cycle GE H System™ to achieve 60 % fuel efficiency while maintaining adherence to the strictest, low NOx standards. Closed-loop cooling eliminates the need for film cooling on the gaspath side of the airfoil and increases the temperature gradients through the airfoil walls. This method of cooling results in higher thermal stresses on the airfoil materials and has led General Electric to use single-crystal superalloys for the first stage, in conjunction with thin
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ceramic thermal barrier coatings. The main features of the MS9001H gas turbine are: – Firing temperature: 1430 8C (2600 8F) – Single-crystal materials. The use of these advanced materials, on the first stage nozzles and buckets, and thermal barrier coatings, on the first and second stage nozzles and buckets, ensures that these components will stand up to high firing temperatures while meeting maintenance intervals. – Air flow: 685 kg/s (1510 lb/s) – Pressure ratio: 23:1 – Four-stage turbine with stage 1 single-crystal blades N5 coated, steam-cooled – Can-annular lean pre-mix DLN-2.5 dry low NOx (DLN) combustor system – NOx 25 (ppmvd at 15 % O2) The first H System™ located at Baglan Bay, Wales, has been in commercial operation since September 2003 and has achieved significant operating experience. The 109H Baglan Bay Power Station aimed to validate H System™ technology with 520 MW power output in combined cycle operation and an efficiency of over 60 %. This was followed by the following new projects: – 36109H TEPCO Futtsu 4 CCGT Power Station with 1520 MW power output in combined cycle operation – 26107H Calpine Inland Empire CCGT Power Station with 800 MW power output in combined cycle operation .
Mitsubishi takes G technology one step further with a secondgeneration G-Series turbine that offers near-H performance. The M701G2 gas turbine is the latest 50 Hz version and incorporates a number of H-Series advances, including steam-cooling of both the combustor transition pieces and turbine vane rings, as well as an advanced dry low-NOx combustion system and an improved compressor that delivers an impressive 21:1 pressure ratio with only 14 stages. Mitsubishi expects that efficiencies close to 60 % can be achieved without extending steam cooling to rotating components (vanes). Gas turbine design features include: – – – – –
Multiple circular arc airfoils in compressor blades Controlled diffusion airfoils in compressor blades Full-coverage film cooling of turbine blades Thermal barrier coating on turbine blades Directionally solidified cast blades
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– Tangential on-board injection of cooled and filtered compressor discharge air – State-of-the-art seal technology (blade ring cooling scheme is designed to provide turbine tip clearance control for better efficiency) – Advanced dual fuel pre-mix combustion with steam-cooled combustor The M701G2 combines the field-proven pedigree of a G with near-H output and efficiency – 334MW/39.5 % in simple cycle applications and 489MW/58.7 % in combined cycle configurations. The first commercial application of the 334 MW M701G2 gas turbine has been the Tokyo Electric Power Co.’s (TEPCO) Kawasaki Combined Cycle Thermal Power Station, one of the most advanced 50 Hz singleshaft combined cycle plants in operation, with demonstrated 59.1 % combined cycle efficiency. On 12 March 2009, Mitsubishi Heavy Industries, Ltd. (MHI) completed development of the ‘J-series’ gas turbine. The unit is designed to operate at a very high temperature near 1600 8C at the turbine inlet. The company also launched activities towards the turbine’s commercial production. Mitsubishi, which has over 60 G-class units sold with steam cooling, has chosen to apply the same steam cooling scheme to the J-class, its latest gas turbine model. This steam cooling scheme is applied to stationary components only, the combustion liners and the first two stages of blade rings. Turbine cooling is based on air, just as in the G. The Mitsubishi J features a 60 Hz simple cycle output of 320 MW and a combined cycle output of 460 MW. The J targets efficiencies of over 60 % in combined cycle. .
Alstom continues improving its GT26 gas turbine for the 50 HZ market and suggests that steam cooling is not the only way to improve efficiency. Alstom and Rolls-Royce signed a long-term technology agreement which will enable Alstom to use Rolls-Royce aeroengine technology in the development of its heavy-duty gas turbine product range. The expertise and knowledge in these areas gained by RollsRoyce in developing its world-leading aero engines are now applied to Alstom’s heavy-duty gas turbines to improve efficiency, power output and durability. Alstom’s GT26/GT24 gas turbine is the only one of its class that makes use of the sequential combustion concept, which consists of a high-pressure combustor followed by a high-pressure turbine, a lowpressure combustor and a low-pressure turbine. Advantages of GT26 Sequential Combustion are seen in the field of low emissions from 40 % to 100 % load.
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Alstom’s new compressor design with a 32:1 compression ratio on the GT26 allows a higher mass flow to be obtained in the gas turbine, which involves a redesign of the compressor with different blading and number of stages. Bleeds are taken from the final compressor stages to provide cooling air for the turbine blades. If the delivery pressure is too high then the bleed air will be too hot for effective blade cooling. Alstom solves this problem by guiding the two air bleeds once through coolers that are connected on the boiler steam-side between the HP economizer and the superheater outlet. Heat is recovered in the HP steam system and not the reheat-IP circuit. The main GT26 gas turbine features are: – – – – – – – – –
Gross electrical output: 288.3 MW Gross electrical efficiency: 38.1 % Pressure ratio: 32:1 Exhaust mass flow: 650 kg/s Exhaust temperature: 616 8C NOx emissions << 25 vppm (at 15 % O2 dry) EV/SEV burners Welded rotor One HP turbine stage with single crystal (CMSX4 with SV coating and film and convection cooled with air from the compressor via internal cooling circuits) and four LP turbine stages, stage 1 with single crystal (CMSX4 with SV20 coating) (convection cooled) and stage 2 with directionally solidified (CM247LC)
The Alstom Combined Cycle portfolio includes the following plant configurations aimed to reach around 60 % efficiency without using steam cooling: KA26-1 Single Shaft 1 on 1 Plant net output: 424.0 MW Plant net efficiency: 58.3 % Plant net heat rate: 6172 kJ/kW h KA26-2 Multi Shaft 2 on 1 Plant net output: 850.3 MW Plant net efficiency: 58.5 % Plant net heat rate: 6156 kJ/kW h KA26-2 ICSTM Multi Shaft (ICS = integrated cycle solution) Plant net output: 857.7 MW Plant net efficiency: 59 % Plant net heat rate: 6103 kJ/kW h The Alstom development path is based on upgrading mature technologies
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with proven elements and calibrated tools. Evolutionary steps ensuring a conservative approach towards 60 % CCPP efficiency with existing platforms. The service life requirements for these industrial gas turbines are significantly longer compared to aircraft engines. This places more emphasis on time-dependent phenomena such as creep and creep–fatigue interactions. Creep–fatigue interactions become more pronounced when materials are thermally cycled while operating at higher temperatures for long times. The effect of creep–fatigue interactions can significantly reduce the fatigue strength of the material. This phenomenon places an additional demand for fatigue data with hold times to simulate the effect of this interaction. The current generations of land-based gas turbines introduce advancements in cooling technology, aerodynamic design and mechanical innovation. Materials technology has necessarily kept pace, with some evolutionary improvements derived from recent versions of the prior generation and some significant introductions of advanced concepts. The key components of the hot gas path are the combustor, transition pieces, turbine vanes, blades and discs. The materials technology relating to these components is discussed in the following sections.
9.4.1 Compressor section Advances in compressor development during the last few years consist of increases in the compressor ratio and air flow, which lead to higher temperatures and stresses in the compressor airfoil. Thus, higher strength materials must be used. Compressor blading (stationary vanes and rotating blades) is variously made by forging, extrusion or machining. All production blades, until recently, have been made from Type 403, 403 Cb (both 12 Cr), or 410 stainless steels. This class of alloy provides the tensile strength at ambient temperature and the strength retention and creep resistance needed for the entire temperature range of compressor vane service (up to about 398 8C (750 8F)). Other qualities of this class of alloy that are attractive for the compressor blading applications include toughness and resistance to impact caused by foreign objects, good mechanical damping and resistance to atmospheric corrosion in a variety of environments. Rotating blades are either precision-forged or machined, with very thin leading and trailing edges for aerodynamic efficiency. The stationary vanes, like blades, may be produced as individual pieces by forging, extrusion, machining or rolling. Corrosion of the compressor vanes is generally not a problem unless the turbine is operated in salty or acidic environments or intermittently in humid climates.
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Table 9.9
Compressor blade alloys [44]
Compressor blades Cr AISI 403 AISI 403 + Cb GTD-450
Ni
Co Fe
12 — — 12 — — 15.5 6.3 —
W Mo Ti
Bal — — Bal — — Bal — 0.8
Al Cb
V
C
B
Ta
— — — — 0.11 — — — — 0.2 — 0.15 — — — — — — 0.03 — —
Compressor efficiency is significantly deteriorated by vane surface roughness. However, a protective coating can be applied to the blades and vanes to maintain performance. Nickel–cadmium plating has a history of successful resistance to compressor vane corrosion in marine environments. Aluminum-rich sacrificial coatings have also proved effective, particularly in acid environments. During the 1980s (Table 9.9), a new compressor blade material, GTD450, a precipitation-hardened martensitic stainless steel was introduced into production for advanced and uprated machines. This material provides increased tensile strength without sacrificing stress corrosion resistance. Substantial increases in the high-cycle fatigue and corrosion fatigue strength are also achieved with this material, compared to Type 403. The use of coating on Type 403 or 410 stainless steel has significant cost benefits over the use of a corrosion-resistant nickel alloy or titanium. However, these materials may be necessary in future generations of gas turbines with high compression ratios. In particular, nickel-based alloys may be required for later stages where compressed air temperature would exceed 398 or 426 8C (750 or 800 8F).
9.4.2 Combustion section Designing to reliably achieve a turbine inlet temperature of around 1426 8C (2600 8F) for long periods requires advanced materials and cooling technologies. Since combustion system components often have the shortest life in the turbine, early attention should be given to selecting and developing the most promising materials technologies. Potentially applicable technologies to be investigated include advanced thermal barrier coatings, ceramic composites, oxide-dispersion strengthened alloys and fabrication techniques allowing improved air cooling. There are three basic designs for combustion systems: the can, annular and external silo types. Current models use mainly can type or annular combustors and the materials issues involved are the same for both. The (can type) combustion system is a multiple-chamber assembly composed of three basic parts: the fuel injection system, the cylindrical combustion liner and the transition piece, which directs the hot gases to the first stage guide vanes of the turbine with an acceptable temperature pattern.
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Table 9.10 Superalloy sheet materials used for industrial gas turbine combustors and transition pieces [45] Alloy Nickel base Hastelloy X Inconel 617 Inconel 625 Nimonic C-263 Cobalt base Haynes Alloy 188 Iron base AISI Type 310
Ni
Fe
Bal Bal Bal Bal
18.5 1.5 22.0 9.0 1.5 12.5 22.0 9.0 2.5 — 21.5 9.0 — 20.5 20.0 5.9
22.0
Co
1.5 Bal
22.0 Bal
—
Cr
Mo W 0.6 — — —
Al
Ti
C
Other
— 1.2 0.2 0.5
— 0.3 0.2 2.2
0.10 0.10 0.05 0.06
— 0.2 Cu 3.6 Cb —
22.0 —
14.0 —
—
0.1
0.07 La
25.0 —
—
—
0.1
—
—
Fuel injection systems are complex systems containing oil and gas burning nozzles. A mixture of compressor discharge air and fuel is burned near the neck. The combustor is designed to contain the flame, mix in diluent air to control the temperature, emissions and smoke, and provide for air-cooling of the metal walls. Burners can be of dry or wet low-NOx type. A key problem in burners is fuel nozzle wear. In addition, material overheating problems due to flashbacks are often evident. Combustion burner components are repairable and replaceable. For this type of service, materials chosen must possess high-temperature strength, including tensile and creep strength, as well as resistance to high-cycle fatigue, low-cycle fatigue, oxidation and carburization. Further, the materials should maintain metallurgical stability in service to avoid embrittlement and be fabricable and weldable in sheet forms, both for initial manufacture and for ease of repairing service-induced defects. Typical alloys used for industrial gas turbine combustor and transition pieces are listed in Table 9.10. The austenitic grades of stainless steel (e.g. Type 310) were initially chosen for combustors and are still being used successfully. Turbines with higher firing temperatures have generally employed nickel or cobalt alloy sheet, Hastelloy XTM and Haynes 188 being popular choices. For improved creep strength retention, alloys such as Inconel 617 and 625 and Nimonic C-263 are being used for some of the newer designs. Combustor wall temperatures can be very high. The materials undergo very abrupt temperature changes on start and stop, so that low-cycle fatigue (LCF) is an important failure mode. Steady-state thermal stresses can also be significant due to the nature of the combustion process and the need for wall cooling. A design that is flexible but still able to withstand thermal fatigue is a must. The combustion process itself generates high-frequency
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Table 9.11 OEM
Combustion turbine materials [43]
Combustor
Transition
Turbine discs
Alstom
Annular-nickel alloy (e.g. IN 617 (duct with silo Alloy steel (A469/ Hastelloy) with configuration) 565) (12CrNiMoV) ceramic refractory tile (welded) GE Cans-Hastelloy and Nimonic 263 with Nickel alloy (IN 706 Haynes 188 with TBC TBC and IN 718) IN 617 (duct with silo Alloy steel (A565) (22 Siemens Annular-Hastelloy XTM with TBC configuration) CrMoV & 12 CrNiMo) S-W Cans-Hastelloy with TBC IN 617 with TBC and Alloy steel Haynes 230 with (NiCrMoV) TBC TomilloyTM (similar Alloy steel MHI Cans-TomilloyTM with TBC to IN 617) (NiCrMoV)
vibrations, which can result in high-cycle fatigue (HCF) failure. The relatively thin walls of the combustor can mean that oxidation is also an important failure mode. Finally, the pressure outside the combustor is somewhat higher than that inside. Since the walls are thin, this pressure difference is sufficient for creep rupture and buckling to be a concern where temperatures are highest. Easy maintenance of the combustor is a requirement, while clearly some introduction of advanced technology that would allow the use of higher inlet temperature machines with extended inspection, repair or replacement intervals would be desirable. A summary of the materials used in combustors in the current model GT is provided in Table 9.11. The combustion gas from the combustor exit is directed toward the first ring of stationary airfoils (nozzle guide vanes). It is important that the flow of hot gas into this stage is distributed as uniformly as possible. The ducts that perform this function are the transition pieces. In the case of an engine with can combustors, the transition pieces have a circular cross-section at one end, which joins the can. The other end has an approximately rectangular cross-section, which on its narrow sides adjoins the neighboring transition pieces, forming a more-or-less continuous ring facing the inlet nozzle guide vanes. The gas entering the transition piece is at the burner outlet temperature and leaves the transition at about the same temperature. Although the transition pieces are less complicated than the liners, they have been more challenging from a materials/process standpoint because of the combination of stresses and temperatures encountered in service and because less effective cooling is possible. Initial 1950s transition pieces were made from AISI 309 stainless steel. In the early 1960s, nickel-base alloys Hastelloy-XTM and RA-333 were used in the more life-limiting parts. The
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Nimonic 263 material is being phased into the higher firing temperature gas turbine models and will be used in future uprated machines. Since the early 1980s, TBCs have been applied to the transition pieces of the higher firing temperature gas turbine models and to uprated machines. Field experience over thousands of hours of service has demonstrated good durability for this coating on transition pieces. The primary basis for the material changes that have occurred has been increased high-temperature creep-rupture strength. These material changes had to be done while maintaining satisfactory oxidation/corrosion resistance. Nimonic 263, the most recently introduced alloy, is some 140 8C (2508F) stronger than the original AISI 309 stainless steel. HastelloyXTM, which was used in the 1960s through the early 1980s, is intermediate in strength between the two. As firing temperatures increased in the newer gas turbine models, HS-188 has recently been employed in the latter section of some combustion liners for improved creep-rupture strength. In addition to the base material changes, the use of a thermal barrier coating (TBC) on combustion liners of advanced and uprated machines has been incorporated.
9.4.3 Turbine section The increase in firing temperature has been the result both of design improvements and of the development of hot-gas path materials with increased temperature capability. Design improvements include the use of lower stress hollow blades, control of the gas-path radial temperature profile and the use of vanes and blade air cooling (which started in about 1960). Materials developments include the use of vacuum-processed nickel-base superalloy, improved cobalt-base alloys and coatings resistant to oxidation and hot corrosion. The function of the stationary vanes is to take the hot gas from the combustor (or previous stage) and turn it (generally accelerating the gas as well) so that it reaches the next rotating stage at the optimum angle. The first stage vanes experience the highest gas temperature in the turbine itself, and the highest gas velocity. Since the vanes are stationary, the only stresses are due to the aerodynamic loading and thermal gradients. They are therefore routinely subjected to impingement of the highest temperature gases and experience the highest metal temperatures of any component in the turbine. Even though the superalloys used for vanes are capable of creep resistance at temperatures above 926 8C (1700 8F) for short periods of time, the desire for a component lifetime of 50 000 to 100 000 hours for industrial turbine vanes means that a high degree of cooling is necessary. Although there are no centrifugal stresses on the vanes, the combination of gas bending loads and the thermal gradients caused by the vane cooling
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Table 9.12 Nominal compositions of selected alloys used for industrial gas turbine vanes
Cobalt base X-40 (Stellite 31) X-45 FSX-414 Mar M 509 ECY-768
Co
Ni
Fe
Bal
10.0 1.5 25.0 7.5 —
Bal Bal Bal Bal
10.5 10.5 10.0 10.0
2.0 2.0 1.0 1.0
Cr
25.0 29.5 21.5 23.5
W
7.0 7.0 7.0 7.0
Ta Ti
— — 3.5 3.5
Iron base Multimet N-155 20.0 20.0 Bal 21.0 2.5 — AISI Type 310 Nickel base GTD-222
—
20.0 Bal 25.0 —
19
Bal
—
—
22.5 2.0 —
C
B
Zr
Other
—
0.5
—
—
—
— — 0.2 0.2
0.25 0.25 0.60 0.60
0.01 0.012 0.01 0.01
— — 0.5 0.05
— — — 0.15 Al
—
0.15
—
—
3.0 Mo; 1.0 Cb
—
0.1
—
—
1.2 0.008 1.00
—
2.3 Mo; 0.8 Al; 0.10 V
Source: Based on adaptation of reference [45].
result in rather high localized steady-state operating stresses in stationary vanes. Thermal stresses from uneven heating and cooling of leading and trailing edges during startup and shutdown are also a factor and can cause cracking at these locations. A list of the properties required by the first stage vanes (‘nozzle guide vanes’) are: (1) excellent oxidation and corrosion resistance, (2) high resistance to thermal fatigue, (3) relatively good weldability for ease of manufacture and repair, (4) good castability and (5) creep resistance to withstand steady-state thermal stresses and the weight of the vane assemblies themselves, which are cantilevered from the outer diameter. Stationary vanes used for industrial gas turbines as currently built are single or multiple airfoil investment castings made from a cobalt-base or nickel-base alloy. Cobalt-base alloys generally possess superior strength at very high temperatures, compared to nickel-base alloys. In addition to the mechanical property requirements already discussed, the material must be easily cast into large, complex (internal cooling passages) configurations. A further requirement is weldability for ease of fabrication (cooling inserts are welded in place) and for repair of service-induced damage. Table 9.12 shows selected alloys used for an industrial gas turbine vane. GE’s practice has been to use cobalt-base alloys for the stator vanes. X-40 and X-45 were used in the early years, but were eventually replaced by FSX414, which has less carbon to improve weldability and more chromium to enhance oxidation resistance. Long-life tests in a simulated gas turbine
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combustion chamber have demonstrated a two- to three-fold increase in oxidation resistance compared to X-40 and X-45. This improvement permitted an increase in the firing temperatures of approximately 56 8C (100 8F) for an equivalent nozzle oxidation life. Currently, first stage nozzles through the ‘F’ series and the majority of later stage nozzles are made of this alloy. Recently, a new nickel-base alloy, GTD-222, has been developed to provide improved creep strength in the second and third stage nozzles and is currently being used in the later stages of the MS7F/9F machines. It offers an improvement of more than 668C (150 8F) in creep strength compared to FSX-414 and is weld-repairable. An important additional benefit derived from this alloy is enhanced low-temperature hot corrosion resistance. The vane alloys used by Siemens-Westinghouse are ECY768 (a modified Mar-M 509 alloy), WES 100 and X-45. The choice of alloy and degree of cooling depend on mechanical design and performance considerations. In the W501G, all stages of vanes are cast nickel-base alloy, IN 939. Alstom’s current practice is to use the cast nickel-base alloys IN 738 and IN 939. Casting porosity becomes more of a problem as the size of the castings increases, and the porosity can act as a site for fatigue crack initiation. This has led to the general introduction of hot isostatic pressing (HIPing) into the production process, to close up the pores. This results in a very significant improvement in fatigue properties. The IN 738 LC and IN 939 investment cast vanes are cooled to achieve average metal temperatures not exceeding 850 8C (1562 8F). Siemens uses the same alloys for the latest models of its large industrial gas turbines in the ‘V64 series’. The repair of service run vanes is an important consideration. This technique is routinely used to maximize the usefulness of these components. All of the alloys used are repairable by welding, although the difficulty increases with the strength of the alloy and the scope becomes more limited. The use of thermal barrier coatings to facilitate tolerance of higher temperatures and/or reduced cooling flow can have a pay-off approaching 93 8C (200 8F). On the other hand, the turbine blades probably represent the most difficult materials application in the gas turbine. Normal operating conditions include high temperatures and high centrifugal loadings, which are especially significant when one considers the last-row blades. In addition, there are various sorts of vibrations loadings, thermal stresses from cooling and thermal fatigue associated with starts and stops. The first stage blade is very significant as it must withstand the most severe combination of temperature, stress and environment; it is generally the limiting component in the machine. Since 1950, turbine blade material temperature capability has advanced approximately 472 8C (8508F), i.e. approximately 10 8C (20 8F) per year. Advances in alloys and processing,
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Ni
15.5 19.0 18.0 18.0 19.0 16.0 22.5 14.0
12.5 15.0 8.0 10.0 9.0 14.0
— 18.5 10.0 15.0 10.0 9.5
Cr
— 12.0 15.0 15.0 18.0 8.5 19.0 9.5
Co
4.2 5.0 6.0 3.0 — 4.0
— 6.0 3.0 3.0 4.0 1.75 — 1.5
Mo
— — — — 12.5 4.0
— 1.0 1.5 1.25 — 2.6 2.0 3.8
W
— — 4.3 — — —
— — — — — 1.75 1.4 2.8
Ta
6.1 4.4 6.0 5.5 5.0 3.0
0.7 2.0 2.5 2.5 3.0 3.4 1.9 3.0
Al
0.8 3.5 1.0 4.7 2.0 5.0
2.5 3.0 5.0 5.0 3.0 3.4 3.7 4.9
Ti
0.12 0.07 0.1 0.18 0.15 0.17
0.04 0.05 0.07 0.035 0.08 0.10 0.15 0.10
C
0.012 0.025 0.015 0.014 0.015 0.015
— 0.005 0.02 0.035 0.005 0.01 0.01 0.01
B
2.0 — — — 1.8 —
1.0 — — — — 0.9 1.0 —
Cb
Nominal composition of selected nickel-base superalloys used for turbines blades [45]
Industrial turbine alloys Inconel X-750 Bal U-520 Bal U-710 Bal U-720 Bal U-500 Bal IN-738 Bal IN-939 Bal GTD-111 Bal Aircraft turbine alloys IN 713C Bal U-700 Bal B-1900 Bal IN 100 Bal Mar-M 200 Bal R 80 Bal
Alloy
Table 9.13
0.1 — 0.08 0.06 0.05 0.03
— — — 0.035 — 0.05 0.1 0.03
Zr
— — — 1.0 V — —
7.0 Fe — — — — — — —
Other
Cast C/W Cast Cast Cast Cast
Wrought Wrought Wrought Wrought C/W Cast Cast Cast
Form
390 Power plant life management and performance improvement
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MGA2400 (CC Ni alloy) (four rows)
MGA2400 (CC Ni alloy) (four rows)
501/701F
501/701G
IN939 (four rows)
501G
Coatings
NiCrAlY+Si (some rows) TBC (R1V)/NiCrAlY+ Si (R2–4 B&V)Uncoated (R5V). chromized (R3) GTD111(R1)/IN738(R2)/U-500(R3) RT22→GT29-In+(R1B) DS GTD111 (R1)/GTD111 (R2and3) GT33-In(R1)/GT29-In+(R2)/ chromized (R3) SC Rene N5 (R1)/DS GTD111 (R2)DS TBC (R1&2 B&V)All others GTD111(R3and4) (DS version of Rene N4) GT 33 IN738LC (R 1, 2 and 3)/IN792 (R4) CoNiCrAlY+Si (some rows) SC PWA1483 (R1 and 2) TBC (EB-PVD) (R1B)/MCrAlY +Re IN738 (R1)/U-520 (R2, 3, and 4) TBC (partial on R1V)/MCrAlY +Re IN738 LC (four rows) TBC (R1B&V)/MCrAlY (R2 and 3B)/Sermalloy J (R4B)* DS MM002 (R1 and 2)/EA CM247 TBC (EB-PVD) (R1 B&V)TBC (R3 and 4)Or DS CM 247 (R1 and 2) (R2 B&V) MCrAlY (R3 B&V) MGA1400 CC (similar to IN 792) (four rows) TBC (R1 B&V)/MCrAlY (R2 and 3B) MGA1400 DS (R1 and 2)/MGA1400 CC (R3 TBC (R1 and 2 B&V)/MCrAlY and 4) (R3 B&V)
IN 738L DS CM247LC(R1-3)/Mar–M247LC (R4&5)
Blades
Rene N5: 7.5 Co, 7Cr, 1.5Mo, 5W, 3Re, 6.5Ta, 6.2 Al, 0.05C, 0.2B, 0.01Y SC PWA 1483: 12%Cr alloy similar to INCO792 GT29: CoCrAlY: GT29+ CoCrAlY + diffusion aluminide top coat GT33: MCRALY with improved oxidation resistance a TBC with MCrAlY bond coat for row 1 blade and vane (plasma sprayed). b TBC with MCrAlY bond coat for rows 1 and 2 blades and vanes.
MHI
ECY-768 (R1 and 3)/X-45 (R2 and 4)
501F
SC Rene N5 (R1)/DS GTD222 (R2)Rene108 (R3)/GTD222 (R4) Siemens V84/94.2 IN 939 (all four rows) V84//94.3A (SC)PWA1483 (R1 and 2)/IN939 (R3 and 4) S-W 501D5/D5A ECY-768 (R1 and 3)/X-45 (R2 and 4)
7H
7/9EA 7/9FA
GE
FSX-414 (all stages) FSX-414 (R1)/GTD222 (R2 and 3)
11N2 IN939 GT24/GT26 DS CM247LC(R1-3)/Mar-M247LC (R2 and 3)/IN738 (R4 and 5)
Alstom
Vanes
Model
Combustion turbine alloy and coatings for hot section blading [43]
OEM
Table 9.14
Gas turbine materials selection 391
392
Power plant life management and performance improvement
while expensive and time-consuming, provide significant incentives through increased power density and improved efficiency. Thus, the rotating blades are more highly stressed than the vanes, but experience a slightly lower gas temperature. Required properties include those listed for the stationary vanes (oxidation and corrosion resistance, thermal and low-cycle fatigue resistance, excellent microstructural long-term stability, repairability, ease of manufacture-castability and creep resistance) and also high-cycle fatigue resistance to withstand various sorts of vibratory loading. Furthermore, in today’s highly cooled airfoils, there is a need to tolerate large, highly localized stresses adjacent to cooling passages. Finally, the metallurgical properties of long-term microstructural stability and coatability must be considered [46]. Table 9.13 shows selected nickel-base superalloys used for turbine blades in aircraft and industrial gas turbine applications. In combustion turbines (Table 9.14), the creep-rupture lifetime of both the wrought and cast blades is generally limited, in part, by cracking of grain boundaries normal to the direction of centrifugal loading. To minimize or completely preclude intergranular creep cracking, the concept of directional solidification (DS) was introduced by VerSnyder. The DS is a blade with an oriented grain structure that runs parallel to the major axis of the part (aligned with the loading direction) and contains no transverse grain boundaries. The elimination of these transverse grain boundaries confers additional creep (to 1 or 2 % strain) and rupture strength on the alloy, and the orientation of the grain structure provides a favorable modulus of elasticity in the longitudinal direction to enhance fatigue life. As well as significant improvements with respect to resistance to thermal fatigue, crack growth and oxidation have also been noted compared to conventionally cast alloys. Table 9.15 lists the composition of some candidate alloys under investigation for both DS and single-crystal applications. DS alloys initially resembled the composition of more advanced equiaxed materials. It was soon discovered, however, that the use of the DS technique plus the addition of hafnium made it possible to use high-strength eutectic compositions, producing significant increases in rupture strength. Lowering the Ti/Al and addition of Ta also helped optimize the alloys in terms of oxidation resistance. By directionally solidifying the alloy GTD-111TM, an increase of approximately 23 8C (40 8F) in creep strength and an increase of approximately 10 times in fatigue life can be realized. More significant improvements in creep and fatigue properties are achieved by adopting alloys developed in a single-crystal (SC) configuration. The creep strength as well as the thermal fatigue resistance of the SC alloy shows a nine-fold improvement with respect to the polycrystalline material.
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C
14.0 8.25 9.0 9.0 22
Cr
9.5 10 10 10 10
Co 3.0 5.5 5.0 5.5 2.3
Al 4.9 1.0 2.0 1.5 3.5
Ti 1.5 0.78 — — —
Mo 3.8 10 12.5 10 2.0
W
1.0 — 0.8
Nb
Composition (wt%)
2.8 2.8 — 2.5 1.1
Ta
0.05 0.05 0.1
0.03
Zr 0.01 0.015 0.02 0.015 0.1
B
Composition (wt%) of cast DS superalloy for turbine blade application [43]
Directionally solidified DSGTD 111 0.10 DSMAR-M247 MAR-M200 Hf 0.15 MAR-M002 DS 0.15 IN 6203 0.15
Alloy
Table 9.15
1.5 Hf 2.0 Hf 1.5 Hf 0.75 Hf
Other
1970 1975 1981
1987
Approx. year of introduction
1040 1045 1020
Temperature capability (8C)
Gas turbine materials selection
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9.20
Benefits of a single crystal [45].
The oxidation/hot corrosion resistance is increased by a factor of three. As the Cr content increases, the hot corrosion resistance increases, but at the expense of the creep strength. The best combination of the two properties at a given Cr level is achieved in single crystals. The current alloys of choice by the GT manufacturers seem to favor lower Cr content SC alloys to achieve high creep strength. Figure 9.20 shows the benefits in service life of the SC application. The use of one alloy such as Rene N5 can produce an increase of more than 35 8C (~60 8F) in creep strength and 26 to 36 increase in fatigue life compared to DS GTD-111TM. Over a period of about 20 years, the chemistry of the single-crystal superalloys has been refined in order to improve their properties. Thus, the composition of the single-crystal superalloys have evolved significantly in the period since 1980 when the first ones emerged. Table 9.15 gives the compositions of important first, second, third and fourth generation singlecrystal superalloys. As previously discussed, both the protective coating and thermal barrier are required to assure the mechanical and metallurgical integrity of the base metal. Protective coatings are used to provide oxidation resistance, protection against environmental embrittlement and occasional hot corrosion resulting from possible upsets in the gas clean-up fuel system or inlet filtration, while thermal barrier coatings are used on stationary and rotating airfoils to increase service life, thus reducing thermal stresses and increasing turbine efficiency through reduced cooling requirements.
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Gas turbine materials selection Table 9.16
395
The composition of single-crystal superalloys [47]
First generation single-crystal superalloy Alloy Cr Co Mo W Al Nasair 100 9 — 1 10.5 5.75 CMSX-2 8 4.6 0.6 8 5.6 CMSX-6 9.8 5 3 — 4.8 PWA 1480 10 5 — 4 5 SRR99 8 5 — 10 5.5 RR2000 10 15 3 — 5.5 Rene N4 9 8 2 6 3.7 AM1 7.8 6.5 2 5.7 5.2 AM3 8 5.5 2.25 5 6 TMS-6 9.2 — — 8.7 5.3 TMS-12 6.6 — — 12.8 5.2 Second generation single-crystal superalloy Alloy Cr Co Mo Re W CMSX-4 6.5 9 0.6 3 6 PWA1484 5 10 2 3 6 Rene N5 7 8 2 3 5 MC2 8 5 2 — 8 TMS-82+ 4.9 7.8 1.9 2.4 8.7 Third generation single-crystal superalloy Alloy Cr Co Mo Re W CMSX-10 2 3 0.4 6 5 Rene N6 4.2 12.5 1.4 5.4 6
Ti 1.2 1 4.7 1.5 2.2 4 4.2 1.1 2 — —
Ta 3.3 6 2 12 3 — 4 7.9 3.5 10.4 7.7
Nb — — — — — — 0.5 — — — —
V — — — — — 1 — — — — —
Hf — — 0.1 — — — — — — — —
Ni Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal
Al 5.6 5.6 6.2 5 5.3
Ti 1 — — 1.5 0.5
Ta 6.5 8.7 7 6 6
Nb — — — — —
Hf 0.1 0.1 0.2 — 0.1
Ni Bal Bal Bal Bal Bal
Al Ti 5.7 0.2 5.75 —
Ta 8 7.2
6
Hf Others Ni 0.03 0.1 Nb Bal 0.15 0.05C Bal 0.004B 0.01Y 0.1 — Bal
TMS-75 3 12 2 5 6 Fourth generation single-crystal superalloy Alloy Cr Co Mo Re Ru MC-NG 4 <0.2 1 4 4 MX4/PW1497 2 16.5 2.0 5.95 3
6
—
W 5 6.0
TMS-138 TMS-162
6.1 5.8
Al Ta Hf Others Ni 6.0 5 0.10 0.5 Ti Bal 5.55 8.25 0.15 0.03C Bal 0.004B 5.8 5.6 0.05 — Bal 5.8 5.6 0.09 — Bal
2.8 2.9
5.8 2.9 5.8 3.9
5.1 4.9
1.9 6
9.4.4 Rotor The turbine rotating blades are commonly mounted on the discs by means of ‘fir-tree’ roots. Rotational forces are thus transmitted through the interconnected discs that comprise the turbine rotor. The discs are highly loaded by the centrifugal force of the rotating blades and would run at temperatures approaching those for blades, unless cooled. The potential for increased disk temperature, of course, is directly related to the turbine inlet temperatures. As discussed in a prior section, this presents a significant design challenge. The continuous increase in turbine inlet temperature to 2600 8F will require a disc alloy capable of operating at higher stresses and temperatures.
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Nickel alloys will most likely be required. Properties of such alloys must be determined and non-destructive tests for large discs must also be established. In general, there are two types of rotor designs, the multidisc type and the welded rotor. The first can be classified depending on how torque is transmitted along the shaft, i.e. shrunk-on discs with torque transmission by shrink-fit, a central bolt with torque transmission by Hirth serration and a single bolt (Siemens design) and multiple bolts/friction with torque transmission by friction between discs. The solid welded rotor consists in a one-piece design with forged discs welded together. This design has been used in all Alstom gas and steam turbines since 1929. Since the rotor is welded, it has no through-bolts and no risk of misalignment. The solid welded rotor design eliminates scheduled maintenance work such as restacking and disc replacement during the service life of the rotor. Compressor rotors have been made by shrinking discs on to forged shafts, bolting discs together or by welding discs together at the outer diameter to form a monolithic construction. The disc alloys currently in use are shown in Table 9.17. Because of the relatively low temperature of the compressor (up to approximately 398 8C (750 8F)), it is possible to use low alloy steel for the discs. A major consideration for the material selection and processing is protection against brittle fracture. Because of the large section sizes a quenched and tempered vacuum degassed NiCrMoV steel is commonly used on stages where the disc service temperature is below the onset of temper embrittlement (approximately 315 8C (600 8F)). Beyond this temperature, material selection and processing follow practices used for (steam) turbine discs (i.e. CrMoV, 12 Cr, hot spinning). It is general practice to use vacuum-degassed steel for the manufacture of NiCrMoV steel to avoid hydrogen flaking. Magnetic particle and ultrasonic inspection are used to check for critical defects. The operating temperatures for turbine discs are considerably lower than for turbine blading, but the stresses experienced are much greater. Under these circumstances, the properties required of a turbine disc alloy can be summarized as follows: (i) high yield stress and tensile strength to prevent yield and fracture, (ii) ductility and fracture toughness, to impart tolerance to defects, (iii) resistance to the initiation of fatigue cracking and (iv) fatigue crack propagation rates that are as low as possible. Creep resistance is also important, but traditionally it has been given less emphasis due to the lower temperatures and because a stress relaxation capability around notches and features of stress concentration is desirable. Table 9.18 provides the composition of some turbine disc alloys that are used in gas turbine engines. Depending on the gas turbine duty the potential and predominant rotor life damage mechanisms are creep life exhaustion,
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Gas turbine materials selection Table 9.17
397
Use of iron-base alloys in industrial gas turbines
Component
Alloy
Nominal composition
Casing
Carbon steel plate (e.g. A-515 Fe; 0.3C; 0.75 Mn; 0.25 Si Gr.70) Inlet, compressor Carbon steel casting (e.g. A- Fe; 0.3 C; 0.8 Mn; 0.5 Si 356 Gr. 1) and combustor Cast iron (e.g. A-278) Fe ; 3.8 total C max. Turbine
Low-alloy steel plate (e.g. A- Fe; 2.25 Cr; 1.0 Mo; 0.15 C 387 Gr. 22) Low-alloy steel castings (e.g. Fe; 2.25 Cr; 1.0 Mo; 0.15 C A-217 Gr. WC9) Modular iron (e.g. A-395) Fe; 3.0 total C min.; 2.5 Si
Compressor blades AISI Type 403 AISI Type 403 + Cb GTD-450
Fe; 12 Cr; 0.12 C Fe; 12 Cr; 0.15 C; 0.2 Cb Fe; 15.5 Cr; 0.03 C; 6.3 Ni
Compressor vanes
AISI Type 403
Fe; 12 Cr; 0.12 C
Discs
Low-alloy steel (e.g. AISI4140) Low-alloy steel (e.g. AISI4340) Low-alloy steel (e.g. A-471 Gr. 10) Super 12 Chrome steel (e.g. AISI Type 422)
Fe; 1.0 Cr; 0.2 Mo; 0.4 C
Super 12 chrome steel (M152) Multimet (N-155)
Fe; 2.5 Ni; 12 Cr; 1.7 Mo; 0.3 V; 0.12 C Fe; 20 Ni; 20 Co; 21 Cr; 3 Mo; 2.5 W; 1 Cb; 0.15 C; 1.5 Mn Fe; 20 Ni; 25 Cr; 0.1C
Compressor
Turbine
Turbine vanes
Fe; 2.0 Ni ; 0.75 Cr; 0.25 Mo; 0.4 C Fe; 1.2 Cr; 1.15 Mo; 0.25 V; 0.3C Fe; 0.5 Ni; 12 Cr; 1.1 Mo; 0.3 W 1.1 W; 0.75 Mn; 0.5 Si; 0.25 C Super 12 chrome steel (e.g. Fe; 0.5 Ni; 6 Co; 1.1 Cr; 0.75 FV535) Mo; 0,25 V; 0.4 C; 0.9 Mn; 0.5 Si; 0.09 Low-alloy steel (similar to A- Fe; 3.5 Ni; 1.75 Cr; 0.5 Mo; 0.1 471) V; 0.35 C; Low-alloy steel (similar to A- Fe; 1.2 Cr; 1.15 Mo; 0.25 V; 471) 0.3C Iron-base superalloy (e.g. Fe; 26 Ni; 13.5 Cr; 2.75 Mo; Discalloy) 1.75 Ti; 0.1 Al; 0.04 C; 0.9 Mn Iron-base superalloy (e.g. A- Fe; 26 Ni; 15 Cr; 1.25 Mo; 0.3 286) V; 2.15 Ti; 0.2 Al; 0.5 C; 1.4 Mn
AISI Type 310 Source: Based on an adaptation of reference [45].
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Cr
11.5 15.0 16.0 19.0 13.1 12.4 11.5 16.0 14.0 13.1 15.0 18.0 19.0 15.0 18.0 17.9 16.0 19.5
Alloy 10 Astroloy Inconel 706 Inconel 718 ME3 MERL-76 N18 Rene 88DT Rene 95 Rene 104 RR1000 Udimet 500 Udimet 520 Udimet 700 Udimet 710 Udimet 720 Udimet 720LI Waspaloy
15 17.0 — — 18.2 18.6 15.7 13.0 8.0 18.2 18.5 18.5 12.0 17.0 15.0 14.7 15.0 13.5
Co
2.3 5.3 — 3.0 3.8 3.3 6.5 4.0 3.5 3.8 5.0 4.0 6.0 5.0 3.0 3.0 3.0 4.3
Mo 5.9 — — — 1.9 — 0.6 4.0 3.5 1.9 — — 1.0 — 1.5 1.25 1.25 —
W 1.7 — 2.9 5.1 1.4 1.4 — 0.7 3.5 1.4 1.1 — — — — — — —
Nb 3.8 4.0 0.2 0.5 3.5 0.2 4.35 2.1 3.5 3.5 3.0 2.9 2.0 4.0 2.5 2.5 2.5 1.3
Al 3.9 3.5 1.8 0.9 3.5 4.3 4.35 3.7 2.5 3.5 3.6 2.9 3.0 3.5 5.0 5.0 5.0 3.0
Ti 0.75 — — — 2.7 — — — — 2.7 2.0 — — — — — — —
Ta — — 40.0 18.5 — — — — — — — — — — — — — —
Fe
The chemical composition of some common turbine disc alloys, in wt% [48]
Alloy
Table 9.18
— — — — — 0.35 0.45 — — — 0.5 — — — — — — —
Hf 0.030 0.06 0.03 0.04 0.030 0.050 0.015 0.03 0.15 0.030 0.027 0.08 0.05 0.06 0.07 0.035 0.025 0.08
C
0.020 0.030 — — 0.030 0.03 0.015 0.015 0.010 0.030 0.015 0.006 0.005 0.030 0.020 0.033 0.018 0.006
B
0.05 — — — 0.050 0.06 0.03 0.03 0.05 0.050 0.06 0.05 — — — 0.03 0.05 —
Zr
Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal Bal
Ni
398 Power plant life management and performance improvement
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399
low-cycle fatigue due to the start–stop cycle, thermal softening and reduction in strength, thermal fatigue damage from temperature cycles, high-cycle fatigue crack initiation and propagation, material embrittlement in service and loss of fracture toughness, crack growth during service from pre-existing defects or material anomalies above a threshold size. The most sensitive location for the above-mentioned damage mechanism that have to be inspected are the central bore surface and near bore region, periphery grooves, fillets, blade attachments, locking slots and cooling slots.
9.4.5 Casing The casings of the land-based gas turbine are a horizontally split and bolted design similar to that used in steam turbine practice. They are cast in grey iron or ferritic modular iron or steel or are weld-fabricated from carbon or Cr steel. An advantage of cast iron is its low cost and excellent producibility. Grey iron is restricted to a service temperature below about 232 8C (450 8F) and thus is used for forward compressor casings. Compressor discharge casings and turbine casings are made from ferretic modular cast iron for service up to about 343 8C (650 8F) or from welded carbon or CrMo steel. Combustor casings and wrappers are also fabricated from carbon or CrMo steel. Major design considerations for these casings are thermal fatigue cracks, particularly at nozzle support hooks, and tolerance for defects to prevent low-cycle fatigue growth and brittle fracture.
9.4.6 Seals The efficiency of a gas turbine can be significantly improved by optimizing various seals and clearances throughout the engine. The high-pressure, hightemperature turbine blade outer air seal is especially important. Technology developed for aircraft (e.g. abradable seals) is being modified to obtain the longer life required for land-based engines.
9.5
Advanced materials and coatings
All classes of materials are being pushed to higher temperature capabilities to fulfill the ever-growing demands for increasing performance and efficiencies, while they must also satisfy stringent gas turbine dependability criteria, as explained in the introduction (RAMD). This leads to more complex material processing and higher cost. This long process is the result of very intensive activity in material research and development right from the beginning (Fig. 9.21). The earliest industrial gas turbines used forged turbine blades with cast
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Power plant life management and performance improvement
9.21 Evolution of the high-temperature capability of the superalloys [48].
nozzles. This changed in the late 1960s when castings began to be used in blade applications. Many nozzle and blade castings used in industrial gas turbines are made by using the conventional equiaxed investment casting process. Vacuum casting is used in most cases, except for some of the cobalt alloys, to prevent the highly reactive elements in the superalloys from reacting with the oxygen and nitrogen in the air. With proper control of metal and mold thermal conditions, the molten metal solidifies from the surface toward the center of the mold, creating an equiaxed structure. In the late 1980s directional solidification (DS) was first introduced into industrial gas turbines. Although it has been applied in aircraft engines for over 25 years, considerable process development was necessary to scale-up the process to the sizes of blades used in industrial gas turbines. By exercising careful control over temperature gradients, a planar solidification front can be developed in these large blades. More recently, industrial gas turbine OEMs have introduced singlecrystal (SC) castings that offer additional creep and fatigue benefits through the elimination of grain boundaries. Major improvements in turbine inlet temperatures can be achieved by replacing Ni-base superalloy hot section components with silicon-base ceramic matrix composite (CMC) and silicon nitride (Si3N4) ceramics. Ceramic matrix composite, consisting of a ceramic matrix with dispersed ceramic fibers or particles, or possibly both, have attracted increasing attention because of the favorable changes in mechanical behavior they commonly produce.
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These materials have superior high-temperature mechanical properties, such as strength and creep resistance, compared to Ni-base superalloys. They are also light and possess excellent high-temperature oxidation resistance in clean, dry air, due to the formation of slow-growing, protective silica scale. One major disadvantage of these materials is the lack of environmental durability in combustion environments. Water vapor, a combustion reaction product, reacts with the protective silica scale, forming gaseous reaction products, such as Si(OH)4. In high-pressure, high-gasvelocity combustion environments, this reaction results in rapid recession of these materials, which also suffer from severe hot corrosion in environments contaminated by molten salt. A new class of coatings, the environmental barrier coating (EBC), has been developed in the 1990s to protect Si-base ceramics and ceramic composites from the degradation by water vapor. The current state-of-theart environmental barrier coating comprises three layers: a silicon bond coat, a mullite-based intermediate coat and a barium–strontium–aluminosilicate top coat. Ceramic composites offer important changes in mechanical behavior that could make more desirable materials for use in large land-based gas turbines than present materials. These composites also have critical issues associated with them, in particular mechanical fatigue for almost all such composites and oxidation or environmental embrittlement in ceramic fiber composites. Other very promising alternatives are the oxide dispersion strengthened alloys (ODS) developed by The International Nickel Company. In this process, mixtures of elemental metal powders, master alloy powders and very fine refractory oxide particles are prepared and charged into a highenergy ball mill or attritor, where mechanical alloying occurs. The individual particles are cold-welded and fractured by repeated high-energy mechanical impact until the alloy consists of a uniform dispersion of a highly refractory oxide (e.g. Y2O3) in a metallic superalloy matrix. The most commonly investigated ODS alloy, MA 6000, contains 15 % Cr, 2.5 vol % Y2O3 with various amounts of W, Al, Ti, Fe, Ta, Mo, C, B and Zr. The lower density of MA 6000 compared to DS Mar-M 247 provides lighter weight blades; the higher modulus of MA 6000 may, however, be unfavorable since peak load stresses can be generated by smaller transient strains. It is reported that the ODS alloy MA 6000 has greater rupture strength than several conventional and single-crystal superalloys at high temperatures/high stresses. The fact that ODS superalloys are better than SC superalloys at high temperatures/low stresses but somewhat worse at low temperatures/high stresses clearly means that some turbine components will benefit from being manufactured from an ODS alloy while some others will not. Another
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important property that must be considered when selecting superalloys for turbine blade applications is fatigue. ODS alloys have the potential for use in the combustion section as well as in uncooled turbine hardware in applications that would require air cooling of conventional superalloys. It may also be possible to use ODS alloys with cooled hardware to operate at higher temperatures. However, they present problems for fabrication by conventional fusion welding techniques because of the agglomeration of the oxides in the overheated zone and the accompanying strength loss. Diffusion bonding, brazing and mechanical joining are, however, possible options. Ductile ordered alloys in the iron, nickel, titanium and cobalt aluminide series constitute a new class of material that, based on developmental results now being obtained, appear attractive as potential materials for gas turbine engine construction. Ordered alloys operate in areas where there are two or more atomic species and these two species occupy specific sites in the crystal lattice. Such alloys tend to occur at well-defined atomic ratios (i.e. AB, A3B, AB3, etc.). Although laboratory results are very promising and development work in titanium aluminides is well advanced, no ordered alloy has reached the point where sufficient manufacturing and design data exist to permit engineering utilization of such alloys in other than experimental hardware. However, the mechanical properties being observed, particularly at elevated temperature, combined with a high aluminum content, which promises good oxidation resistance at elevated temperatures, direct attention to the aluminides as potentially very useful alloys for future gas turbine design. Generally speaking, the modern alloys are characterized by significantly lower concentration of Cr and higher concentrations of Al and Re. Concentrations of Ti and Mo are now at very modest levels. The period since 2000 has seen the emergence of the fourth generation single-crystal superalloys, such as, MC-NG, EPM-102 and TMS-162, which are characterized by additions of ruthenium. On the other hand, thermal barrier coatings will be more aggressively designed to protect gas turbine engine hot-section components in order to meet future engine higher fuel efficiency and lower emission goals. Advanced thermal barrier coatings have been developed using a multicomponent defect clustering approach and have been shown to have improved thermal stability and lower conductivity. The coating systems have been demonstrated for high-temperature combustor applications. For thermal barrier coatings designed for turbine airfoil applications, further improved erosion and impact resistance are crucial for engine performance and durability. Erosion-resistant thermal barrier coatings are being developed, with a current emphasis on the toughness improvements using a combined rare earth and transition metaloxide doping approach.
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Ceramic thermal barrier coatings permit significantly increased gas temperature, thus reducing cooling requirements and improving engine fuel efficiency and reliability. It has been shown that decreasing the bondcoat temperature by 25 8C can double the TBC life. A different design approach can be followed, i.e. by increasing TBC thickness, by reducing TBC conductivity and by increasing cooling air. The most effective way is by reducing TBC conductivity. This leads to a recognition that in crystalline solids in general heat is transferred by three mechanism: (i) electrons, which are expected to provide the major contribution at very low temperature, (ii) lattice vibration, i.e. anharmonic phonon–phonon scattering, which dominates at intermediate temperature and (iii) radiation, i.e. photons, which will contribute strongly at very high temperatures. The above arguments indicate that thermal conductivity reduction can be won by dopants introduced to YSZ-based ceramics to encourage phonon dispersion. Rare earths added in their oxide form, e.g. erbia, ytterbia, neodymia and gadalonia, have been shown to be effective at reducing the thermal conductivity because they act as phonon scatters. Siemens and MHI are in production with advanced TBC compositions capable of tolerating higher temperatures, having improved phase stability, sintering resistance and lower conductivity. These new coatings have been designated as ‘low k TBC’ and ‘ultra low k TBC’ and could achieve around 30 % and 50 % lower conductivity respectively. General Electric is developing a new alternative for EB-PVD (electron beam–physical vapour deposition), with a thermal barrier coating made of suspension plasma spray (SPS), which provides excellent bondcoat strength. Suspension plasma spray is a novel process for producing nanostructured coatings with metastable phases using extra small particles as compared to conventional thermal spraying. Suspension spraying involves atomization, solvent evaporation and melts consolidation, which can cause substantial complexity in the system. SPS is potentially a disruptive technology. Another advanced coating system is the solution precursor plasma spray process (SPPS), a relative by new and flexible thermal spray process that can produce a wide variety of novel materials, including some with superior properties. The SPPS provides a longer coating life than APS or PVD coatings.
9.6
Life management and diagnostic
Gas turbine engine hot-section components are designed to operate in hightemperature environments, with high thermal gradients and mechanical loads. Repeated engine startup and shutdown operations subject components to cyclic strains that are generated both thermally and mechanically. These thermomechanical fatigue cycles cause microstructural material damage to
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the components and lead ultimately to fatigue crack initiation, crack growth, and failure. Advanced high-temperature fatigue life prediction methodologies are needed to decide when to replace engine components, in order to increase reliability and service life, and reduce maintenance costs. However, predicting the stress and strain response of components, and hence their lives under this TMF loading is a great challenge. Sophisticated life prediction models have been developed to take creep–fatigue interactions and other damage effects into account, and so improve life predictions under TMF conditions. Much work has been performed, with qualified success, in modeling the TMF life of materials on a crack initiation basis. The prediction of crack initiation (for instance a surface crack about 1 mm long) requires knowledge of the effects of local strain range, mean stress, temperature and operational profile on the formation, the link-up and the propagation of slip band or grain boundary microcracks or a crack from a defect. Generally the life prediction for turbine blades is based solely on crack initiation while for vanes and combustor liners, which have a much greater damage tolerance, the crack growth life may be used as well as the initiation life. In the development of realistic models for crack initiation in gas turbine materials, the following considerations should be kept in mind: . . . . . .
Experimental conditions should be established by the parameters of the engine operating environment. Material property scatter should be accounted for. Ideally, models should account for various interactive effects (e.g. environment and operational profile correlations). The use of models in structural life prediction analyses should not require parameters that can neither be reliably calculated by the available analytical tools nor be directly measured. Models should be kept as simple as possible and should rely on a minimum amount of specimen testing. Interpolative and extrapolative capabilities of the models should be assessed carefully by selective testing of simple and/or complex geometry specimens.
The basic approach used in current life prediction methods is to measure the lifetime of simple specimens under known environmental conditions and then to develop analytical models that accurately describe their behavior. The suitability of the analytical models is then tested by comparing the predicted lives with the observed (measured) lives of simulated hardware (or model) specimens, which are designed to have appropriate complex geometries and stress–strain states. Upon successful completion at this level, the analytical models are then used to predict the lifetimes of actual engine components with verification by rig testing.
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TMF life prediction may include both TMF crack initiation models and crack propagation models, but as previously noted, only the former are addressed in the following sections. The more sophisticated models require many variables and associated parameters in the life equations, to represent the principal damage mechanisms. With these considerations, a review of one of each lifetime prediction methods for thermo-mechanical fatigue crack initiation from the most popular models, such as, damage summation, frequency modified model, Ostergren’s damage function, frequency separation and ductility exhaustion, to the most advanced, such as strain-range partitioning, total-strain version of SRP and strain energy partitioning, will be discussed. The models used include damage-based criteria, stress-based criteria, strain-based criteria and energy-based criteria.
9.6.1 Damage summation model In the design of equipment operating at temperature levels where creep was a significant design parameter, methods were devised to sum the ‘damage’ from the creep loading with that from the cyclic loading. A summation of time and cycle fractions is the oldest of the hightemperature creep–fatigue life prediction methods. In this method for predicting life, it is assumed that at high temperatures, there are two independent types of damage that develop. The first is conventional fatigue damage Df, analogous to the damage that occurs in the fatigue process at low temperature where creep is absent. The second is a high-temperature form of damage Dc, which would eventually cause failure independent of cyclic loading, analogous to that obtained in elevated temperature creep rupture tests. The total damage or linear damage summation (DS) model (also called the linear life fraction or linear cumulative damage) is the simplest expression for creep–fatigue prediction: Dfatigue þ Dcreep ¼ Dtotal
½9:16
and failure occurs when Dt is equal to a critical value, usually unity. By means of the 1945 Miner rule [49] for fatigue damage and the 1952 Robinson rule [50] for creep damage, with repeated application of a single simple cycle, equation (9.16) becomes 1 th Nf þ ½9:17 ¼ Dtotal Np Tc The way in which the fatigue and creep damages are determined is as follows. The fatigue damage issuing from one applied cycle is defined as the reciprocal of the number of cycles to failure that would be obtained (in the
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absence of creep) if the applied cycle were allowed to be continually repeated. This is represented by Nf, the number of cycles to failure at a given strain-range, and Np, the pure fatigue life at that strain-range. The creep damage fraction in each cycle is th/Tc, in which th is the hold time at a given stress in a cycle and Tc is the time to rupture under static creep at that stress. This would result if this stress were maintained until failure. An ambiguity arises in defining creep damage under compressive stress, since monotonic creep rupture does not occur under compressive stress. A way to take this into account in a conservative manner is by assuming compressive stresses to be as damaging as tensile stresses. Under thermomechanical strain cycles, the creep damage consists both of a loading creep strain (as load/temperature is held constant or increases with time above a threshold creep value) and a dwell creep strain (when strain is held constant over the hold time resulting in stress relaxation). Ellison and Zamily [51] modified life prediction by this fraction rule to deal with the changing strain, summing over th/Tc (approximated as a series of steps) [52] as follows X ti X ti th total ¼ loading þ dwell ½9:18 Tc Tci Tci Analogously to summing the creep damage fraction, fatigue damage is obtained by summing the cyclic damage under differing loadings. The linear damage summation in equation [9.17] should then be written as X N f X ti þ ¼ Dtotal Npi Tci
½9:19
This approach has an advantage over other approaches in that it requires only high-frequency fatigue data and conventional monotonic creep-rupture data. Also, it is very easy to take different types of cycles into account in the summations required and is compatible with current stress analysis techniques. The disadvantages of this approach are first the inability to account realistically for compressive creep effects. Second, from the fact that since the rupture life is very sensitive to stress, it is necessary to know the stresses very accurately. Third, the result could show considerable scatter resulting from a creep–fatigue synergism. Fourth and last, the damage summation model ignores the failure process; in other words, the sequence effect on the creep–fatigue life is omitted. In this sense, it assumes that the subsequent creep life after prior fatigue is the same as the subsequent fatigue life after prior creep. However, in many cases, creep damage and fatigue damage are not independent [53].
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9.6.2 Strain-range partitioning model For high-temperature low-cycle fatigue, the failure process is affected by the time-dependent cyclic stress–strain behavior over the entire temperature range. To account for the time-dependent portions of the cycle, Manson et al. [28] developed the strain-range partitioning (SRP) method. The strainrange partitioning (SRP) approach involves partitioning of the total inelastic strain range into four possible components depending on the direction of straining (tension or compression) and the type of inelastic strain accumulated (creep or time-independent plasticity) [52]. Under cyclic reversed loading, there are four possible combination cycles of inelastic strain that the SRP model needs to consider. The cycle types are tensile plasticity reversed by compressive plasticity (PP), tensile creep reversed by compressive creep (CC), tensile creep reversed by compressive plasticity (CP) and tensile plasticity reversed by compressive creep (PC). In any established combined cycle, a maximum of three cycle types are physically possible, PP, CC and either PC or CP. Figure 9.16 shows the four generic types of hysteresis loops for the four types of strain ranges. The actual hysteresis loop from a creep–fatigue test (i.e. an LCF test with hold time) is broken down into the component strains: Δεcc, Δεpc and Δεcp. The terms Δεpp and Δεcc represent the pure reversed plastic and reversed creep strain ranges respectively, and the other two terms represent combined creep and plastic strain ranges. For each type of strain range, the Coffin– Manson relationship can be applied. For instance, Npp = A(Δεpp)α, and so on. The fractional strain for each type of strain with respect to the total inelastic strain can be expressed as Fpp ¼
Depp Decc Depc Decp ; Fcc ¼ ; Fpc ¼ ; Fcp ¼ Dein Dein Dein Dein
½9:20
where εin is total inelastic strain. By adding up the fractional damage for each type of strain, the total damage is estimated by the expression based on the following simple linear damage rule of Manson et al. [54]: 1 Fpp Fcc Fcp Fpc ¼ þ þ or ½9:21 Nf Npp Ncc Ncp Npc where Npp, Ncc, Npc and Ncp are the fatigue lives produced by PP, CC, PC and CP inelastic strain cycles respectively and represent the number of cycles to failure for each type of strain. One of the advantages of the SRP rule is that it is relatively temperatureindependent (Halford et al. [55]). The life relationships are governed by the four inelastic strain ranges and are not greatly affected by the temperature at
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which the strains are imposed. If particular strain ranges are imposed at one temperature, the life will be similar to that when the same strain ranges of the same type are applied at another temperature. This does not mean that life is independent of temperature because a given imposed load will produce different strains at a different temperature and a given strain will be partitioned differently depending on the temperature. To obtain the four unique partitioned strain-range versus cyclic life relationships, four separate PP, PC, CC and CP creep–fatigue tests have to be done. Results of these tests are fitted to the Coffin–Manson equation [56] and can expressed as C Depp ¼ App Npp pp ½9:22 C Depc ¼ Apc Npc pc
½9:23
C Decp ¼ Acp Ncp cp
½9:24
Decc ¼ Acc ðNcc ÞCcc
½9:25
The coefficients A and exponents C are experimentally determined material constants. To determine the TMF life of engine components using the SRP model, the procedures are summarized as follows. First, based on the cyclic stress versus strain response, obtain the partitioned inelastic strain ranges, Δεpp, Δεpc, Δεcp and Δεcc, and then the corresponding strain fractions, Fpp, Fpc, Fcp and Fcc. Second, from the total inelastic strain range, solve for Npp, Npc, Ncp and Ncc using equations [9.22] to [9.25]. Finally, determine the TMF lives by the interaction damage rule equation [9.21]. The lives determined in the above steps are for a theoretical zero mean stress condition. The SRP model has been extensively applied in the nuclear and aerospace industries. The frequency, hold time and temperature effects are built into the SRP model since these variables are implicit in the stress–strain hysteresis loop. However, it is important to recognize the limitations of the SRP method to make the best use of it. The model is not applicable to nonductile materials since the inelastic strain is too small to be determined correctly. The SRP model does not take into account the environmental attack caused by oxidation and may overpredict the life. It is still difficult to partition the inelastic strain experimentally. Improved prediction methods for thermomechanical fatigue life will assist in reducing life-cycle costs and increasing the availability of the hot-section components in gas turbines. Reliable life assessment of engine structural components using the state-of-the-art life prediction models, forms the basis
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of good design and of design modifications that may lead to component life extension. All the models mentioned are able to predict life with acceptable accuracy within a certain envelope of material and test conditions, but no model can consistently produce accurate life predictions for all materials under different service conditions. Therefore a prediction model should be used only when it has been validated for the material category and expected environmental conditions. If data are not available, predicted lifetimes should be validated against laboratory test data or by in-service monitoring. Although the SRP (strain-range partitioning) and TS-SRP (total strain version of strain-range partitioning) models have attracted great research effort [57], all the current design codes and rules are still based on the earlier linear damage summation model. Further research on life prediction models is needed, and long-term validation and analysis will be required to make the new models more generally acceptable.
9.7
Future trends
Market advantage today relies on increasing performance and efficiency as well as reducing life-cycle costs. The first statement is responsible for pushing present materials and coatings to their limits with significant consequences in terms of the durability and maintenance costs for machines that rely on these advanced hot section designs. The second statement emphasizes the need to focus on four key areas of development: advanced repair technology of components, non-destructive evaluation (NDE), online monitoring techniques and remaining life assessment methods supported by experimental testing and operational data. As more and more DS and SC components are used, component costs will rise significantly. To minimize costs, it is essential to rejuvenate and repair parts rather than replace them. To optimize rejuvenation and repair frequencies and minimize unit shutdown it is necessary to develop inspection techniques to determine the amount of life consumed and on-line monitoring to detect the initiation and propagation of cracks. Condition assessment of components requires a combination of both online and off-line techniques. i.e. on-line condition monitoring will allow us the continuous or periodic measurement and interpretation of operational data to indicate the condition of a component to determine the need for maintenance. It should be based not only on standard instrumentation but also on diagnostic instrumentation (i.e. new instrumentation providing additional information about component condition and potential failure modes and, in this way, assist in improving availability through shortening maintenance time and by extending the time interval between routine inspections).
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On-line monitoring techniques based on optical (to measure properties such as chemical composition, phase information, acoustic velocity, moduli and thermal diffusivity) and electromagnetic (to measure properties such as conductivity, permeability or dielectric properties) sensors are proving to be very promising. On the other hand, off-line inspection techniques (i.e. advanced non-destructive evaluation methods) are used for non-destructive degradation assessment of hot gas-path component coatings and base materials during their service life. These techniques are a key alternative to destructive sectioning of these expensive components in order metallographically to assess damage and the percentage of life consumed. The valuable information from the above-mentioned techniques constitutes the starting point for developing and tuning methods for predicting the life of critical high-temperature components of a gas turbine, which are available and used but need to be extended for the development of the next generation of gas turbine designs. The critical components of concern are blades, vanes, discs (wheels), combustors and transition pieces. The new generation of turbines is expected to operate reliably between 100 000 and 200 000 hours with shutdowns and startups perhaps as often as daily cycling and at a turbine inlet temperature of up to 2600 8F. The principal damage mechanisms are creep and thermal fatigue, complicated by environmental effects where coatings play an important role (especially oxidation and some unavoidable corrosion effects), brittle fracture and long-time material structural instability as well as embrittlement reactions (e.g. temper embrittlement, aging). It is possible to conclude that three principal sources of damage generally determine the service life of a hot-section component: coating oxidation or depletion, thermal mechanical fatigue and creep. All three have to be considered for modeling purposes. Nowadays, damage accumulation is very much machine-specific and location-specific for a given part. From it arises the necessity of verifying the life assessment models through experimental and operational measurements made on operating engines for selected models and duty cycle applications. In comparison with current maintenance practices, the remaining life of components is assessed by using analytical methods and destructive tests. These approaches have inherent limitations. Analytical methods must use worst-case assumptions about stresses, material properties and operation so that the results tend to be conservative. In the case of destructive tests, testing material from the most severely damaged location may not always be available and there is usually insufficient time to test long-term properties and insufficient material to test all relevant properties. Finally, advanced repair and rejuvenation are processes undertaken to extend the useful life of a component. Repairs can be thought of as ‘external’ processes that return the component to its original size and shape
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or replace the protective coating. Rejuvenation can be defined as the regeneration of a microstructure leading to the restoration of mechanical properties equivalent to those of the original component prior to initial service. These reconditioning technologies together with the off-line and on-line techniques and life assessment will give way to a new vision of gas turbine asset and performance management. This leads us to the concept of condition-based maintenance (CBM), which is defined as the preventive maintenance that should be initiated as a result of knowledge of the condition on an item from routine and continuous monitoring. CBM based upon non-destructive flaw evaluation and fracture mechanics arises as a way for promoting lifetime extension and maintenance cost reduction. Then crack growth approaches are considered in addition to crack initiation methods. CBM is based upon specific measurements of individual components and the tracking of its service history. As a result, we are able to obtain an accurate and reliable prediction of the current and projected (‘health’) condition.
9.8
Sources of further information and advice
The following is a list of further information related to the main topics dealt with in this chapter. – Metals: ○ Metals and Materials Society of the American Institute of Mining, Metallurgical and Petroleum Engineers, New York, USA (www.tms.org) ○ Nickel Development Institute, Toronto, Ontario, Canada (www.nidi.org) ○ ASM International is the society dedicated to serving the materials science and engineering profession, ASM Alloy Center/ (http://asmcommunity.asminternational.org/portal/site/www/ About/) (http://products.asminternational.org/alloycenter/index.jsp) ○ Cobalt Development Institute, Guildford, Surrey, UK (www.cobaltdevinstitute.com) ○ International Chromium Development Institute, Paris, France (www.chromium-assoc.com) ○ Specialty Steel Industry of North America, Washington, DC, USA (www.ssina.com) ○ Metal Suppliers Online (MSO) is a company dedicated to making the buying and selling of industrial metals easier via their industry-
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leading process (http://www.suppliersonline.com/Default.asp) – Investment casting: ○ Precision Cast Parts, Minerva, Ohio; Cleveland, Ohio; Mentor, Ohio; Douglas, California. (http://www.precast.com/index.html) ○ Doncasters Precision Castings, Droitwich, Worcs, UK; Bochum, Germany; Groton, Connecticut, USA. (www.doncasters-deritend.com) (www.doncasters-bochum.com) ○ Hitchiner Manufacturing Co., Gas Turbine Division, Milford New Hampshire, USA (www.hitchiner.com) ○ Alcoa Howmet, Cleveland, Ohio. Alcoa Power and Propulsion. United States locations: Alcoa Howmet. Dover, New Jersey, Hampton, Virginia; La Porte, Indiana; Whitehall, Michigan; Wichita Falls, Texas. Canada location: Alcoa Howmet, Laval, Quebec H7L 3H7; 53602 EvronAlcoa Howmet, 14160 Dives-surmer, 92230 Gennevilliers. Japan location: Howmet Japan Limited, Ishikawa 923-1101. United Kingdom Location: Exeter, Devon EX2 7LG. (http://www.alcoa.com/howmet/en/home.asp) (http://www.alcoa.com/app/en/home.asp) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=72) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=73) (http://www.alcoa.com/howmet/en/info_page/desc_hlc.asp) (http://www.alcoa.com/howmet/en/info_page/desc_hwc.asp) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=77) (http://www.alcoa.com/locations/canada_laval_howmet/en/home. asp) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=252) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=316) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=214) (http://www.alcoa.com/locations/japan_terai/en/home.asp) (http://www.alcoa.com/locations/alcoa_location/en/home.asp? code=194)
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Pacific Cast Technologies (PCT) supplies commercial, military and aerospace OEMs with airframe, launch vehicle, and turbine engine structural components, specializing in investment cast titanium parts exclusively (http://ladishco.com/) (http://www.paccast.com/page.asp?PageID=185) – Forgings: ○ Wyman Gordon Co., N. Grafton, Massachusetts; Livingston, Scotland; Houston, Texas (www.wyman-gordon.com) ○ Schlosser Forge, Cucamonga California (www.aerospace-engine-parts.com) ○ Ladish Co., Cudahy, Wisconsin (www.ladish.com) ○ Doncasters PLC, Monk Bridge, UK; Blaenavon, UK; Leeds, UK (www.doncastersmonkbridge.com) (www.doncastersblaenavon.com) (www.doncasters.com) ○ Thyssen Umformtechnik, Remscheid, Germany (www.tut-gmbh.com). – Coating and/or refurbishment/repair: ○ Chromalloy Gas Turbine Corporation, Carson City, Nevada; Gardena, California; Orangeburg, New York; Harrisburg, Pennsylvania; Middletown, New York; Columbus, Indiana; Manchester, Connecticut; Phoenix, Arizona, USA (www.chromalloy-cnv.com) (www.chromalloy-cla.com) (www.chromalloyhit.com) ○ Sermatech International, Limerick, Pennsylvania; Muncie, Indiana; Houston, Texas; Manchester, Connecticut, USA (www.sermatech.com) ○ Sulzer Turbo Services Venlo B.V., Lomm, The Nether–lands. Sulzer Turbo Services Houston Inc., LaPorte, Texas, USA (www.sulzerts.com) ○ WG Word Group. Word Group Gas Turbine Services. Rolls Wood Group, Wood Group Light Industrial Turbine. Wood Group Heavy Industrial Turbine Solutions (http://portal.woodgroup.com/portal/page? _pageid=11475,1677770&_dad=portal30&_schema=PORTAL30)
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9.9
References
1. 2. 3. 4.
5. 6.
7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.
18. 19. 20. 21. 22.
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58. Fatigue and Fracture Mechanics of High Risks Parts. Application of LEFM & FMDM Theory. Bahram Farahmand with Georage Bockrath and James Glassco. Chapman & Hall.
9.10
Appendix 1: nomenclature
9.10.1 Symbols Dc: conventional creep damage Df: conventional fatigue damage Dt: total damage E: Young’s modulus %EL: percentage elongation F: fractional damage K: stress intensity factor ΔK: stress intensity factor range Nc: predicted lifetime for creep dominated failure Nf: number of cycles to failure Np: predicted lifetime for fatigue dominated failure Tc: time to rupture under static creep %ROA: percentage reduction in area tc: compression hold time tcy: cycle time th: hold time tt: tensile hold time t0: time for continuous cycle portion UTS: ultimate tensile strength ε: strain e_ : creep rate εe: elastic strain range εcp: combined creep and plastic strain εcc: reversed creep strain εpc: combined plastic and creep strain εpp: reversed plastic strain εt: total strain range εp: plastic strain ρ: notch tip radius σ: stress σe: endurance limit σm: mean stress σmax: maximum principal stress σr: alternating stress σy: static yield strength
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9.10.2 Subscripts: c: pure creep or compression-going cc: tensile creep reversed by compressive creep cp: tensile creep reversed by compressive plasticity el: elastic f: failure in: inelastic i: index for specific fatigue conditions ij: pp, cc, pc, cp p: pure fatigue pc: tensile plasticity reversed by compressive creep pp: tensile plasticity reversed by compressive plasticity
9.11
Appendix 2: key definitions
Condition-based maintenance (CBM) ‘The preventive maintenance initiated as a result of knowledge of the condition of an item from routine or continuous monitoring.’ More informally, this is the periodic inspection of components by manual or automatic systems in order that their condition may be assessed and to identify their degradation rates. A decision is then taken regarding replacement, which is based upon an analysis of the monitored data of Seddon [58]. Performance monitoring Performance monitoring is a technology that uses measured and calculated thermodynamic properties to assess the condition of a machine or system. The advantages of performance monitoring systems as a part of preventive/predictive maintenance programs are the economic impact on the turnover of the power plant. Life-cycle cost (LCC) Life-cycle cost is the total cost of ownership to the customer for acquisition, operation and maintenance and disposal of a product for the whole life cycle. Dependability This is collective term used to describe the availability performance and its influencing factors: reliability performance, maintainability performance and maintenance support performance.
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Reliability (R) (How often does it fail?) can be expressed in terms of time as the ‘mean time between failure’ (MTBF). Availability (A) (Can it be used when you want it?) is a measure of an ability to satisfy demand. Maintainability (M) (How easy is it to fix?) is also expressed in terms of time as the ‘mean time to repair’ (MTTR). Durability (D) denotes long continuous run times without forced or planned outages.
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10 Gas turbine maintenance, refurbishment and repair A . D . W I L L I A M S , Wood Group GTS, UK
Abstract: This chapter provides an overview into the practices and procedures used in the repair and rejuvenation of hot section components of the gas turbine. It initially discusses the maintenance procedures currently being adopted, detailing the changes in the market and the drive for improved availability and reliability. It then links these market drivers and illustrates how advanced technology repairs can lead to part life extension and reduced overall costs while maintaining the key required metrics. Key words: field service overhaul, maintenance factors, refurbishment, inspection, welding, vacuum brazing, coating, heat treatment, life extension.
10.1
Introduction
The gas turbine, like any complex mechanical machine, requires a comprehensive maintenance programme to ensure plant optimisation, availability and reliability. Optimisation constantly involves improvements, upgrades and modifications to significantly improve operational efficiency with due consideration to current national and international legislation. It is also important that gas turbine operators understand the total life cycle costs in establishing and implementing any maintenance regime. This regime will be different from plant to plant and engine type to engine type and considers numerous factors including people, manufacturers’ recommendations, commercial contracting models and environmental constraints. The following chapter considers some of the procedures required to maintain and operate a gas turbine effectively. It starts by considering the inspection regime and practices operated in the station itself, focusing on the inspection periods, techniques and criteria applied and also considers additional factors
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that have a significant impact on the maintenance cycles. The remaining sections then address the methods and procedures established to enhance the life and performance of the hot section components specifically utilising repair and refurbishment practices. Metallurgical degradation as a result of operational conditions is considered and methods discussed that return critical components back to serviceable condition. Extending component life beyond manufacturers’ recommendations is also reviewed, illustrating the robust processes and practices required to ensure safety and reliability, underpinned with sound economic benefits.
10.2
Field service overhaul and maintenance
Predictive planned maintenance schedules are the backbone in machine reliability. Based on the design, the major engine manufacturers have produced guidelines to owners as to when maintenance should be undertaken and what scope of work should be undertaken at that period. Traditionally the manufacturers and their licencees have followed a protocol based on running hours at base load condition, which is then adjusted on a number of key variables. These periods are not the same from manufacturer to manufacturer, but notwithstanding this variation the maintenance schedules are normally categorised into three areas: 1. 2. 3.
Combustion inspection/A inspection Hot gas path inspection/B inspection Major inspection/C inspection
Clearly the amount of work varies significantly between the three categories, with a combustion inspection lasting as little as three to five days of downtime to a full major, which can run into several weeks dependent on the amount of work required, not only to the turbine but also to the ancilliary plant. It is well understood throughout the industry that, as a generalisation, the majority of breakdowns caused in a power plant are a result of the ancilliary plant surrounding the turbine and not the turbine itself. As a result, when opportunities arise to maintain and inspect other areas of the power plant, including items such as the generator and heat recovery steam generator (HRSG), then these should be built into the maintenance programmes. Combustion inspections typically involve the removal of the fuel nozzles, combustion liners and transition pieces. These are either replaced with brand new components or parts that have previously operated and been refurbished to approved and proven procedures by either the engine manufacturer or third party service providers. These inspections normally occur after about 8000 hours of operation, which equates to about one year
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of continuous firing. Once removed, visual inspection of the early stages of the turbine is also undertaken to ensure trouble-free operation. Hot gas path inspections are significantly more labour intensive than the combustion inspection mentioned above and include the removal of the turbine casing followed by the removal and replacement of the key hot gas path components. This inspection also incorporates a combustion inspection ensuring that availability is optimised. Again the replacement parts can be wholly supplied by the manufacturer or a combination of new manufacturer parts, new third party parts, third party repaired parts or manufacturer repaired parts. This combination can create significant challenges for the field service teams in rebuilding the engine and ensuring that all critical clearances are maintained. Clearance dimensions have a critical impact on engine performance, reliability and efficiency. The clearances are held to within tenths of millimetres and although these engines are over 3 m in diameter the tolerances used are extremely tight. This is made even more difficult when the engine being overhauled is many years old, has been operated at elevated temperatures in excess of 1100 oC for long periods of time and can therefore be subjected to casing and shell distortion. However, as operational performance has a marked affect on profitability it is critical that these tolerances are held. The major inspection, in general, occurs every six years or 48 000 operating hours. This constitutes a full stripdown of the complete engine including hot and cold sections and removal, inspection and, if necessary, disassembly of the rotor. Figure 10.1 illustrates a heavy industrial rotor in the process of being disassembled. These inspections can last again for several weeks but should improve the efficiency and the engine performance significantly and return it to similar levels as if the engine were new. There will always be a variation in performance as again engines subjected to harsh operating regimes for a number of years will always have a degree of degradation that cannot be recovered. Over the life of a major service interval the effects on compressor performance can be significant and it is during this inspection that the compressor can be thoroughly cleaned, inspected and if necessary have new parts installed. A rebuild of all these critical components is vital, starting with full rotor alignment and balancing. These engines rotate at up to and in some cases over 5000 rpm and minor amounts of imbalance can lead to significant vibration, resulting, in the worst cases, to engine trips and even component failure. Documentary control and inspection reporting is paramount at every stage of the process to ensure that a satisfactory build programme is completed. The traditional combustion inspection, hot gas path inspection and major overhaul detailed above are constantly being reviewed as a better understanding of engine performance is obtained. In many instances the periodicity of these inspections is being extended to 32 000 hours for a hot
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10.1 Heavy industrial gas turbine rotor being disassembled.
gas path or even 41 000 hours on some of the large Siemens designed units. This has a significant cost benefit over the life of the turbine and is clearly something that is being driven by the operators as a cost reduction process. Furthermore, the drive away from time based inspections to condition based assessments has also gathered momentum over recent years and remote monitoring and diagnostic systems are now more commonplace in ensuring safety, reliability and availability.
10.2.1 Factors affecting maintenance intervals Maintenance intervals are set on the basis of continuous operation under ISO conditions of 25 oC ambient temperature and 15 % oxygen, operating on clean natural gas. Any factor that varies from this base line has to be taken into consideration when calculating when to do the maintenance work and when to repair or replace components. The key variables that have the main effect on this calculation are: 1. 2. 3. 4.
Fuel type Startup and shutdown cycles Trips from full load Water/steam injection
Each manufacturer makes recommendations for calculating the equivalent
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operating hours (EOH) based on a combination of the factors listed above. There are significant industrial discussions as to which is right or wrong and the substance behind the calculation, but regardless of the provenance these calculations are used as the basis for deciding the maintenance strategy. Fuel types Fuel types vary tremendously from natural gas, right through the octane range to heavy crude oil. Over recent years and into the future other fuel types including biomass and hydrogen rich fuels from coal gasification need to be considered, which will not only have an effect on the maintenance cycle but also on the overall life of the engine and its constituent parts. Different fuels, depending on their chemistry and calorific value, have a significant impact on parts life and therefore maintenance schedules. Liquid fuels with significant impurities of sulphur and vanadium give rise to hot corrosion products forming on the hot gas path parts. Without significant on-line washing techniques being applied these salts will rapidly degrade the rotating blades and static vane assemblies, leading to component failure and forced outage. Hydrogen rich fuels have the opposite effect in terms of flame temperature compared with the liquid fuels, but the result of forced outage can be the same. In these cases hydrogen burns with a very hot flame, leading to increased gas temperatures in the turbine. If unchecked by the use of coating technology and improved cooling designs, the materials would be rapidly oxidised, leading to premature degradation and failure. Starts and stops Changes to commercial contracts have resulted in many engines being operated away from the traditional base load duty cycle to a highly cyclic startup based regime, where the engines are brought on line to meet an energy demand signal and then taken off line when demand is low. This can happen daily or in some cases on a more sporadic basis. In both cases reliability and availability are the key drivers from an operator’s perspective but the challenges from a maintenance perspective are even more significant. Operating any components in a thermally cyclic operation results in significant thermal and mechanical stresses being induced. Parts are heated and cooled at relatively rapid frequencies, giving rise to both low cycle and high cycle fatigue conditions. Maintenance factors are therefore introduced to increase the frequency of inspections to ensure that the integrity of the turbine is maintained.
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Trips from load This is more of an unplanned occurrence and results in fast changes in mechanical and thermal load being applied to the unit. This can happen as a result of numerous operational factors and is something that needs to be avoided where possible. This has a significant impact on component life and integrity and any trip needs to be incorporated into the overall part life calculation. Water/steam injection This feature has the direct effect of altering the heat transfer coefficient of the gas stream and is used for either power augmentation or nitrous oxide (NOx) reduction. In both instances water or steam is injected into the engine either via the fuel nozzles themselves for water or via the combustion endcover and compressor exit using steam. The way the gas turbine is operated and controlled may have a significant impact on parts life and understanding whether the turbine is operating with a dry control curve or a wet control curve is an important aspect in life management. A small increase in steam (i.e. about 2–3 %) will result in an increase in the heat transfer coefficient of about 3–4 %, which increases the actual blade metal temperature by up to 10 oC, reducing the overall part life by 25–30 %. These are significant and must be taken into consideration in the overall gas turbine maintenance strategy.
10.2.2 Condition based inspections Although the maintenance practices are extremely well documented and methodically proven, the drive to reduce the inspection intervals is significant. Operating the units from one hot gas path cycle to the next without intervention has been driven by the overall cost reduction, resulting from the time saving created by the removal of two combustion inspections and the need for improved availability, again through not having the unit off line. Reliability therefore is one of the key drivers and having confidence in the engine performance is critical. Building up this confidence that the engine is performing correctly and will continue to perform without major intervention can occur in two ways: 1. 2.
Build up a maintenance plan over a period of time using frequent noninvasive borescope examinations to inspect the engine condition. Attach a condition monitoring system to the unit to predict signals being emitted accurately, allowing early action to be taken.
In reality a combination of both of the above will be imposed in the first
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10.2 Screen data photograph detailing information that can be identified remotely from the control system.
instance until confidence is achieved. Moving to the second option is clearly preferable as monitoring and early onset diagnostic work can be achieved remotely and early solutions can be discussed. Having a greater understanding of the data through statistical trend analysis can lead to fundamental savings in alleviating significant downtime at a later date. Figure 10.2 illustrates a typical data screen viewed remotely to ascertain trends in engine performance.
10.3
Parts refurbishment: incoming inspection
Regardless of the maintenance regime employed by the operator there will come a point where the critical hot gas path components will require refurbishment. Techniques established in the aircraft gas turbine overhaul industry have been adopted and in some cases adapted for use in land based turbines. Most of the technologies have evolved over the last 30 years as turbine inlet temperatures have increased dramatically to enable an increase in power output and efficiency. In order to withstand these excessive temperatures, component materials, cooling systems and protective coatings have all had to evolve at a similar rate to cope with the hostile environments to which the parts are exposed.
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The combination of high temperatures, fossil fuels, vibration, mechanical and thermal stress all occurring over an extended period of time lead to a multitude of degradation factors. Areas such as creep, fatigue, oxidation and hot corrosion, material degradation and foreign object damage all need to be considered during the parts inspection and it is a critical part of any repair process to diagnose the condition correctly and apply technically viable processes to cure the symptom. Understanding the part history and even the engine history is important to ensure that optimum repair solutions are performed. Components are normally removed from the engine during an overhaul and sent to a specialist repair facility for refurbishment. On receipt, documentation is prepared to capture the part history. Serial numbers and part numbers are identified and a unique batch number is traditionally applied to allow full traceability throughout the repair process. Initially, and for batches of parts where the history indicates that a further operating cycle can be achieved, a small metallurgical section is removed from an area of the component that can be readily repairable. This section is prepared in the laboratory and evaluated to determine the condition of the underlying material received from the customer and the type of coating that has previously been applied. Both of these pieces of information are important in determining the correct heat treatments to be applied during the repair process and also the correct method of coating removal process to be adopted. Depending on the results of this initial evaluation, a bespoke inspection process will be created. Firstly, this may entail the removal of any protective coating to ensure any and all defects are exposed. Techniques utilised for this process are typically as follows: 1. 2. 3.
Submersion in acidic baths Water jet machining Mechanical blending
Ensuring that the coating is completely removed can be achieved through a low temperature heat treatment in air to oxidise the surface layer of the material. This produces a colour change to the surface and allows easy identification of coated and non-coated areas. Etching the component with an acidic surface etch to identify the material grain structure also has the same effect but is more time consuming and is less widely used. On completion of the coating removal, the components must be fully heat treated to ensure that the material is returned to a uniform condition to allow repairability. This is typically known as a solution heat treatment and is carried out in a clean vacuum furnace at a temperature and time that is determined by the material being treated. In addition to the primary reason for performing a heat treatment, additional benefits in the surface preparation can be achieved. These are:
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Power plant life management and performance improvement It adequately cleans the surface of the components, removing any oxide product that may be apparent. By removing the engine debris, defects are opened that enable a better and more robust inspection process to occur.
Full inspection can now take place and inspection processes will be dependent on the complexity and the criticality of the product. Only 20 years ago all static, non-rotating hot section components were exposed to a red dye penetrant inspection to determine surface defects. This was supported by a full visual inspection and dimensional checks of all critical features using manual gauging. The critical rotating components, which experience a more significant mechanical load, were always subjected to a more intense inspection as it is the rotating parts that are more liable to failure during operation. These parts underwent a more sensitive inspection using fluorescent penetrant as the technique to ensure minute defects were captured. The evolution in component design to overcome the harsher environments seen in today’s advanced engines has meant that additional processes have had to be incorporated to ensure that all the defects are identified and reported. Experience gained in the repair of flight turbines has been adopted within the industrial gas turbine industry and now the technologies shown below have become commonplace. The list below details the majority of the inspection techniques that can be applied but the actual ones utilised for a particular component are determined by the component itself, its geometry and the commercial contract. Processes used are as follows: 1. 2. 3. 4. 5. 6. 7. 8.
Flourescent penetrant inspection Radiographic inspection Borescopic inspection Thermography Eddy current Ultrasonics Magnetic particle inspection Dimensional inspection (a) Coordinate measuring technology (b) White light laser scanning (c) Manual clocks and gauges
Each of the technologies have their own advantages and disadvantages and a good working knowledge of each must be understood before recommending them for a particular task. In many instances a combination of technologies must be incorporated into the repair process to ensure a full understanding is achieved. On completion of the inspection process all the results are compiled into a
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comprehensive report. This report should include the level of defects noted, the type of defects encountered and the overall repairability of the component. A detailed engineering review will then establish the repair process to be followed and identify any hold points to be incorporated.
10.4
Parts repair
Repair and rejuvenation focuses on returning the part to a standard that allows another engine cycle to be achieved. It does not and cannot produce a part that is new or that has the exact mechanical and metallurgical properties as a new component, but it will strive to rejuvenate parts to optimise their life and maximise their performance even when received in a badly degraded condition. The defects established during the incoming phase together with the component material type and structure will determine the nature of the repair technologies to be utilised.
10.4.1 Welding The process of welding involves the fusion of two or more pieces of material under the influence of a heat source. Many individual processes fall into this category and some, such as metal inert gas (MIG) and oxyacetylene, although fairly agricultural in nature, still play a valuable role in extending the life of the less complex gas turbine components utilised in the repair of gas turbine components. These techniques have not been considered in this section, which focuses on more advanced or widely used processes. TIG welding Dimensional restoration and major crack repair is traditionally performed using tungsten inert gas welding. This comes in many forms from micro TIG and pulse TIG done under standard atmospheric conditions on a weld bench to TIG welding in a chamber under a protective environment or even at elevated temperatures. In principle, however, these are all variations on a common theme. TIG welding uses a shrouded tungsten electrode to produce an arc with the parent material. This melts the surface locally and creates a molten pool into which is fed a rod of weld filler material. The heat source is relatively intensive and the resultant joint, although strong, has a significant mixing layer between the weld and the parent material commonly referred to as the heat affected zone (HAZ). This layer’s mechanical properties are significantly different to those of either the filler material and the parent material and is typically extremely brittle. Welding is principally a melting and solidification process and during solidification shrinkage can occur if the welding parameters are not adequately controlled, giving rise to
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10.3
Schematic representation of a tungsten inert gas welding system.
microfissuring within this HAZ layer. Reducing this HAZ is therefore extremely useful when welding more complex alloys with significantly improved mechanical strengths, and different techniques have been developed to overcome these issues. Figure 10.3 illustrates a tungsten inert gas welding process. Laser powder fusion welding Laser powder fusion welding (LPW), or alternatively laser cladding, utilises a laser heat source and powder metallurgy technology to actively repair and rejuvenate the component. This technology has grown in its capability as materials have grown in complexity. Having the ability to repair tips of blades, knife edge seals or areas of critical parts that are susceptible to high thermal stresses has been a major enhancement to extending life of hot section components and even repairing high strength, complex chemistry materials, i.e. single crystal and directionally solidified alloys. As stated, the technology utilises a laser as the heat source. This is traditionally a CO2 system, but YAG lasers are now also finding an application as the technology is developed. The laser beam is focused through a series of mirrors on to the workpiece where powder is injected coaxially under a
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10.4 Schematic representation of a laser powder fusion welding system.
shrouded gas into the molten pool. The nozzle is normally mounted into either a six-axis CNC (computer numerical control) machine centre, although again robotic systems are now being utilised to provide better access to complex shapes. Figure 10.4 illustrates a schematic representation of the laser system. Both systems allow accurate control both in terms of heat input and accessibility. LPF creates a weld that is traditionally known as ‘near net shape’. This refers to the amount of weld material being deposited. The lack of heat input, excellent control of the weld deposit and accurate weld speed allow a weld to be produced that has a minimal heat affected zone (HAZ) and very closely approximates to the final shape of the component required, reducing the amount of post-weld blending or machining necessary, and as such leads to an extremely efficient welding process. Figure 10.5 illustrates a typical laser weld being produced. Electron beam welding This is a technology which is underutilised in the gas turbine hot section repair business but has found a niche in the repair of fuel nozzles and for the replacement of tip caps on turbine blades on the advanced class of engines. Instead of using a laser to create the heat source, electron beam welding
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10.5
Laser weld on the tip section of a turbine blade.
(EBW) creates the source through the generation of a stream of electrons, which are accelerated towards the workpiece. The process is carried out in a vacuum chamber but it is the size of the chamber and the access to the joint that limits the use of the technology. Traditionally this technique is used purely as a metal joining technology but some systems have been modified to allow wire to be fed into the weld pool. As a joining technology it relies on pure fusion of the parent materials and therefore the joint design is fundamentally important. Any gaps at the interface can lead to weld defects and centre-line cracking. This technique allows materials of dissimilar chemistries to be joined, but again good quality control is paramount. This technology has found its niche market outside the GT industry in electronics, medical and automotive. Figure 10.6 illustrates a schematic representation of the electron beam welding process and Fig. 10.7 illustrates the structure of the weld produced showing a typical nail head formation.
10.4.2 High temperature vacuum brazing Not to be confused with torch brazing using silver solders, high temperature vacuum brazing uses powder metallurgy and excessively high temperatures to heal defects, build up material sections and join metal. Different companies have developed their own versions of the technology and it is commonly referred to by each company’s own process name (transient phase restoration, surface reaction brazing, turbofix, activated diffusion healing, to name but a few). All these processes basically follow a fairly similar flow: 1. 2.
Cleaning Lay up
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10.6 Schematic representation of an electron beam welding process.
10.7 Photomicrograph of an electron beam weld detailing the typical morphology of the weld.
3. 4. 5.
Heat treatment Blending Inspection
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10.8
A typical hydrogen fluoride cleaning system.
Cleaning This is one of the most critical steps in the whole process. If the alloy is not metallurgically cleaned then the braze powder will not wet the surface and flow correctly. The technology used here utilises either hydrogen, hydrogen fluoride, other fluoride compounds or a combination to reduce the tenacious oxides that are formed on the component surface and on the surface of the cracks during engine operations. In one such system the components are loaded into a special furnace, purged through with an inert gas, heated to temperatures of around 1100 oC and subjected to a cycle of hydrogen and hydrogen fluoride. The gaseous nature of the process ensures that even oxides deep within the tip of a defect are meticulously cleaned. Figure 10.8 illustrates the hydrogen flouride ion cleaning process equipment. Lay up This element of the process applies the braze alloys to the surface of the component. Braze alloys can have a variety of chemical compositions but are normally made up of a combination of the parent metal to be repaired and an alloy including elements such as boron and/or silicon, which lower the overall melting point of the total braze mixture enabling the brazing process to occur at temperatures below the melting point of the parent material itself. These various alloy combinations are normally provided either in a paste form, sheet form or as pre-ordered shapes (pre-sintered preforms) ready to apply to the surface of the material. The healing of a crack usually involves the application of a paste to the surface of the defect or, where the crack is sufficiently wide, braze alloy can be packed into the defect. Where the former occurs healing of
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10.9 Photomicrograph illustrating a fully repaired crack following a transient phase restoration process.
the defect relies on capilliary action of the molten braze alloy taking place – hence the requirement for meticulous cleanliness. Once the alloy has been applied in the various forms, the parts are heat treated within a vacuum environment, again to maintain cleanliness. Heat treatment Optimisation of the heat treatment is again important. Too long at a temperature and too high a temperature can cause the braze alloy to become less viscous and run all over the component. Too short a time at a temperature and too low a temperature has the opposite effect, causing the braze alloy to remain solid, resulting in no fill or no joint. The tolerance window in which to work is extremely narrow and an adequate temperature and time control are vital to the overall quality of the end product. The heat treatment process is often a two stage process. Stage one occurs at a high temperature ensuring the braze alloy becomes molten for sufficient time to allow capilliary action to occur. Stage two occurs at a much lower temperature but for an extended time. This ensures that the small amounts of boron and/or silicon that can cause hard intermetallic phases are diffused away from the joint into the parent material. Blending Post-braze clean-up removes any excess braze material from around the joint, defect or surface build-up and re-profiles the component back to its true condition. Often any porosity that may be in the braze joint will rise to the surface and this can then be removed during the final metal removal process.
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Inspection A final check is then required to ensure that joint integrity is achieved. In the earlier section inspection techniques were discussed thoroughly and the same processes are applied as part of the in-process inspection technology. Figure 10.9 illustrates the structure of a crack following repair detailing the fine intermetallics formed and the limited effect of the braze diffusion layer.
10.5
Coating and finishing technology
Once repaired a final protective layer needs to be applied to the critical components. This coating system is dependent on a number of critical factors: 1. 2. 3. 4. 5. 6.
Inlet temperature Fuel chemistry Operating cycle Component weight and geometry Time between overhauls Wet or dry operation
Coatings act as a sacrificial layer protecting the parent material from the effects of hot corrosion and oxidation. Coatings can be categorised into three distinct groups: 1. 2. 3.
Diffusion coatings Overlay coatings Thermal barriers
10.5.1 Diffusion coatings These systems are created through the diffusion of an enriching element into the surface layer of the material. In the majority of cases this is aluminium but chromium has also been used where hot corrosion is the driving degradation mechanism. These alumina (Al2O3) or chromia (Cr2O3) forming coatings are traditionally known as either first generation diffusion coatings or simple diffusion coatings. They are applied by three main techniques: 1. 2.
Pack cementation Gaseous transfer (a) Chemical vapour deposition (b) Gas phase coatings
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10.10 Schematic representation of a chemical vapour deposition system.
Pack cementation In pack cementation a coating powder incorporating either aluminium or chromium, and in some cases both, is poured over the components contained in a metal box or retort. The powder also comprises typically a bulk carrier compound of aluminium oxide and a halide reactant, which combines with the aluminium and chromium during the reaction to allow vapour transportation to occur. Once covered and shaken the box is placed into an atmospheric furnace where the parts are purged through with argon and heated to high enough temperatures and held for long enough times to allow the diffusion process with the base material to occur and to optimise the coating thickness.
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Chemical vapour deposition (CVD): external gas generation In this process the components are placed directly into a furnace and do not come into direct contact with the coating media. The coating gas itself is generated outside the main reaction vessel and is piped into the main chamber in the gaseous halide form. The reactive gas diffuses into the surface of the component, creating a highly uniform coating system. Manifold systems can be applied so that the gas can be channelled into the core passages of the critical turbine components, thus coating the internal surfaces of the part during the same process as coating the external surfaces. Being a gaseous process, any surface that is exposed ends up having a coating applied. It is therefore critical that the masking process is adequate to protect major contact faces such as the turbine fir tree serrations and the machine rails on static parts. Figure 10.10 illustrates a schematic representation of a chemical vapour deposition system. CVD: internal gas generation In this instance the gas is generated within the reaction vessel . A powder mixture or alloy briquette is placed within the vessel and a small amount of a reactive halide is applied. With the application of heat a chemical reaction occurs, similar to that within the pack cementation process generating the aluminium and/or chromium halide reaction gas. This circulates throughout the reaction chamber either through a direct flow mechanism or through a pressure differential process. With the correct manifold systems both internal and external surfaces can be coated using this system but the control and application engineering is significantly more complex than with the pure gaseous CVD system.
10.5.2 Modified diffusion systems As operational regimes changed and different degradation mechanisms became apparent, the simple aluminides mentioned above were modified to enhance the component performance. The modifications included the codeposition of the main coating elements with other elements, namely the precious metals; silicon and in some cases chromium and aluminium are codeposited as a system in their own right. These coatings are classified as the second generation coatings or modified aluminides. Often these systems are applied as a two stage application but some can be applied as a true codeposition in a one stage process. The processes used here are: 1. 2. 3.
Electroplating plus aluminising/chromising Spray coatings Pack cementation
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Precious metal aluminising This is a multideposition system. In the initial stage precious metal is deposited on to the surface through electrolysis. Solutions of precious metals are created (typically these are platinum based but other precious metals including rhodium and palladium have been used over the years). The parts are robustly cleaned prior to the electroplating process to ensure all contaminants are removed. Typically precious metals have a fairly low deposit efficiency and the parameter control is paramount to ensure that the coating is not deposited too quickly or too thickly. Once applied the deposit is diffused prior to the application of a further coating treatment. This additional process will be either a gaseous process or pack process depending on the coating system being applied, the complexity of the component geometry and the properties of the coating required. Spray coatings Here a metallic based slurry is applied to the surface of the component through a standard atmospheric spray application. The process is extremely dependent on operator training and like all of the application processes, part cleanliness is important. Several layers are applied to allow sufficient thickness to be created. Drying and curing must occur between each layer to prevent the coating running. Once sufficient coating has been applied these coatings are heat treated through a high temperature process in an inert atmosphere to diffuse the alloying elements into the parent material.
10.5.3 Overlay coatings The overlay coatings were developed during the 1980s and have, in the majority of large industrial engines, replaced the diffusion systems as the standard coating system of choice. They use a combination of elements alloyed into a powder which is then accelerated towards the component where it adheres through a mechanical bond. Initially these coatings were applied by a standard atmospheric plasma process but the quality of the coating applied through this method was not adequate enough to withstand the harsh environments experienced on many of the hot section critical components. As a result other application technologies were developed. These were as follows: 1. 2. 3. 4. 5.
Vacuum plasma Low pressure plasma High velocity oxy-fuel Electron beam physical vapour deposition Electro-co-deposition
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Vacuum plasma Vacuum plasma takes the same principles as the standard atmospheric system but applies itself within a large vacuum container. A robot and turntable are contained within the vacuum system and a batch of components are located on to a large fixture which is inserted on to the turntable. The large door is closed and the whole system evacuated. Once the batch of parts has had the required thickness applied the system is completely broken down and the process initiated once more. Although the results from this process were good, the production throughput was relatively slow and technology being developed in the USA became the technology of choice through the 1990s for the application of these coating systems. This was low pressure plasma. Low pressure plasma spray Although these systems were significantly more expensive to purchase (by a factor of 3), the throughput and therefore payback was better. The system, as with vacuum plasma, applies the coating in a vacuum chamber but in this system the main chamber remains at the required vacuum for the duration of the production batch. Attached to the sides of the main chamber are transfer chambers segregated from the main chamber by large valves. The transfer chambers are sized to take the largest component produced within the facility. Each component is loaded on to a transfer arm and the transfer chamber is evacuated to the same pressure as the main chamber. Once equalisation has occurred, the valve opens and the part is translated into the main chamber where it can translate and rotate. Here the parts are heated to a cherry red condition and cleaned by reversing the transfer arc and pulling off material from the surface rather than depositing material on to the surface. Once clean, powder is fed in from hoppers into the plasma nozzle where it is accelerated towards the component within the plasma stream. The plasma gun can be manipulated on an additional three axes, providing adequate part coverage. On completion of the coating cycle the part is returned to its existing chamber and the valve closes. Once clear of the chamber, the next part from the opposite side can be rotated through the cycle, improving the overall process efficiency. Figure 10.11 illustrates a schematic representation of the systems, illustrating the key process elements. Figure 10.12 illustrates the structure of the coating produced, identifying the dense duplex formation of beta nickel aluminide in a gamma matrix.
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10.11 Schematic representation of a low pressure plasma system.
10.12 Photomicrograph of an MCrAlY coating applied by LPPS.
High velocity oxy-fuel (HVOF) Initially high velocity oxy-fuel (HVOF) coatings were used for the application of wear resistant coating to the oil industry on component faces that were liable to see major wear. During the late 1990s the powder particle size on the oxidation and hot corrosion resistant coatings was
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10.13 Schematic representation of a high velocity oxy-fuel coating system.
10.14 Robotic application of the coating applied to a turbine component.
refined so as to be able to use this technology as an alternative to the more complex and capital intensive LPPS systems discussed above. This technology combines many of the advantages of the two processes discussed above to deliver a high quality coating that has now become the system of choice for many coating suppliers. It is achieved under standard atmospheric conditions using a six-axis computer numerical controlled robot and associated turntable, similar to that used for standard plasma spraying. The HVOF gun and the fuel used to accelerate the particles towards the component surface, however, are significantly different to standard plasma spraying and the coating deposited is similar to the quality of that seen in, a VPS or LPPS coating. Figures 10.13 and 10.14 illustrate the application
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processes. Over recent years a debate has ensued about the benefits of LPPS coatings over HVOF applied coatings. It is abundantly clear that, dependent on the parameters developed and the complexity of the component being coated, each of the systems have their advantages and disadvantages. Electron beam physical vapour deposition (EBPVD) These systems were developed initially for the application of metallic coatings but found their niche in applying ceramic coatings to complex geometry components within the aero industry. The capital costs of this type of plant is extensive and runs to several million dollars, resulting in coatings that, although technically superior in terms of metallurgical and mechanical properties, were too expensive to produce. The coating is applied within a complex vacuum chamber similar to that discussed in the section above on LPPS. Contained within the main chamber are electron beam guns that scan across the surface of ceramic or metallic ingots, creating a cloud of the material being deposited. The vapour cloud deposits on the surface of the component in a very controlled uniform manner, developing a coating system characterised by its columnar grain structure within the ceramic systems. The strain tolerance of this coating system is far superior to other coating systems discussed above but the commercial model for industrial gas turbines is difficult to justify. Metallic coatings applied in this manner do not traditionally have a columnar grain appearance but form very dense, uniform thickness coatings similar to that seen in the VPS, HVOF and LPPS systems discussed above. The coating costs, however, are significantly different, primarily due to the huge variation in capital plant and operational costs. It is the commercial issues that have seen a decline in the metallic EBPVD system in favour of the other application systems illustrated. Electro-co-deposition This system has had fairly little market exposure due to the overall time taken to apply a coating of similar thickness and quality to the application methods stated above. Solutions of the complex MCrAlY chemistries are formed and suitably charged through an electroplating technique. This is a true co-deposition process where a complex chemistry is applied to the surface of a component as a single stage process, developing a dense, high quality system. The overall process time, deposit efficiencies, controls and manufacturing space have again limited the technology to specific niche markets.
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10.15 Photomicrograph of a thermal barrier ceramic coating applied over a metallic corrosion resistant coating.
10.5.4 Coating chemistry Metallic systems The overlay coatings are known by the generic term MCrAlY coatings where M equates to a metallic element, being either cobalt or nickel or some combination of the two. The other elemental constituents are relatively straightforward. Cr equates to chromium, Al equates to aluminium and Y equates to yttrium. Additional alloying elements are used depending on the coating being applied and some of these include tantalum, silicon and rhenium. All of these additions stabilise the protective oxide scale that is formed and improve the performance of the coating. Varying the coating chemistry allows optimisation of the coating performance for different operating environments. A higher percentage of chromium and cobalt are used if the coating is to be applied for a hot corrosion environment where heavy sulphur bearing fuels predominate and alternatively higher percentages of nickel and aluminium are present when a more oxidation resistant coating is required. Thermal barrier systems The thermal barrier systems use various ceramic constituents to create a barrier to the temperatures experienced in operation. They do not form a barrier to hot corrosion and oxidation products but do act as an insulation layer reducing the effects of thermal conductivity. These coating chemistries surround zirconia as the main ceramic, being either partially or fully stabilised with other ceramics and elements such as magnesium in the early
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systems and now predominantly yttrium. These coating systems are applied primarily by atmospheric plasma spray but EBPVD applied ceramics are utilised, especially in the military and commercial flight markets where, because of component size, the technical benefit significantly outweighs the high cost of application. Figure 10.15 illustrates the typical structure of an air plasma sprayed thermal barrier coating applied over a bond coat of a metallic composition.
10.6
Final repair operations
Once coated, and dependent on the material, the parts are processed again through a final heat treatment stage to optimise both the coating and material properties. This heat treatment is often a two stage process, ensuring that the material is once again solution annealed to ensure uniformity and then aged or precipitation hardened to optimise mechanical properties. On completion of this final process a thorough inspection of the component must occur, utilising some or all of the processes described in the inspection section earlier in the chapter. Heat treatment can initiate cracking through a number of metallurgical phenomena and it is not uncommon to have to re-work parts at this late stage of the repair process. Further finishing operations that are applied include iron and ceramic shot peening to induce compressive stresses into the root serrations of turbine blades, vibro-polishing to optimise the surface finish of the component and moment weighing of blades and area harmonic analysis of static vane assemblies to optimise positional arrangements during engine build. A final inspection report detailing all of the repair processes, inspection techniques and quality control processes will ensure a part is delivered that meets all the stringent engine criteria and continues to operate through to its next major overhaul inspection.
10.7
Quality control and first article inspection
Developing any repair and coating for a new component requires in-depth analysis and creation of robust engineering protocols and procedures. Evaluation of the key features and degradation mechanisms enables the engineering teams to create repair drawings, tooling and methodology to produce a quality repaired component. Depending on the complexity of the component being developed, complex design technology procedures may be required to fully appreciate the degradation mechanisms being noted. Computational fluid dynamics and complex finite element analysis are now tools regularly used by the repair engineer to determine and evaluate repair options. Once all of the understanding has been completed and the draft procedure created, verification and first article processes are applied to
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establish optimised parameters at each of the critical stages. This analysis can be re-visited at regular intervals to ensure that process parameters remain optimised and have not varied, resulting in a poor quality product. Each process and product is supported by comprehensive and approved paperwork systems detailing operational method statements together with critical quality and health and safety data to be applied at each stage of the process. This provides the operation and also the customers with the confidence that the product is being repaired to the highest standards with the highest controls and integrity.
10.8
Part life extension and optimisation
Any part, whether related to the gas turbine or not, has a finite life. Product designers work to relatively conservative limits ensuring that parts are retired from service well before they fail. This is commendable and ensures that engines have improved in reliability, optimising the generating potential of the power utilities. However, the high cost element in the through life calculation of the gas turbine is that of the capital parts. Technology has increased and components have become more complex, both in terms of geometry and materials, and as a result the costs of the components have escalated. Optimising the overall life of the parts is now being questioned in order to minimise total lifetime costs. As a result, repair companies are reviewing whether the original design intent is correct and whether the parts can operate for another overhaul life safely. Figure 10.16 illustrates a stress analysis model of a turbine blade that has been developed to appreciate fully the degradation mechanisms associated with the original blade design. This process is full of assumptions and relies on sound information being provided, competent and credible metallurgical and mechanical assessments and excellent co-operation between the customer and supplier. Historical information is vital in understanding how the part has performed to date. Information required includes: 1. 2. 3. 4.
Operational hours Previous repair history Previous duty cycle Previous fuel chemistries
Furthermore, similar information regarding the parts future is also helpful in indicating fuel, environment, load and duty cycle. The evaluation is destructive in nature and as a direct result an additional part must be made available and annotated accordingly to ensure that it is identified at a future date. A full photographic profile is created of the set and the component illustrating the most typical damage pattern is identified for analysis. This component is initially fully visually, dimensionally and non-
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10.16 Stress analysis model of a typical turbine blade.
10.17 Schematic representation of a typical metallographic sectioning plan for life assessment.
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destructively inspected, and all results are then incorporated into the final report. Metallurgical samples are then taken, some to be used for metallographic analysis and others to be used for mechanical testing. Enough samples are taken in various locations to ensure that results can be obtained in the as-received condition and also after any heat treatments have been undertaken, in order to understand whether the properties can be recovered during a repair procedure. Figure 10.17 illustrates a typical sectioning plan used to determine both metallurgical and mechanical data during the life extension process. The results of these tests combined with the initial inspections undertaken and the historical evidence gathered allow comprehensive assumptions to be evaluated, leading to conclusions being drawn about the safe operation of the parts for a further overhaul life. There are many instances where these decisions, underpinned with robust testing and analysis, have saved customers many millions of dollars in revenue, improving profitability to the operation and reducing overall costs to the consumer.
10.9
Future trends
The industry continues to evolve and has changed significantly from the early Whittle engine. Over the last 20 years improvements in efficiency, availability and reliability have been driven by higher fossil fuel prices, complex commercial contracts, environmental concerns and also changes in World demographics. Early industrial engines that operated at inlet temperatures around 1000 oC with power outputs of 30 MW or less were fundamentally not efficient or powerful enough for today’s markets and have been taken out of service. This trend has continued and today, although engine technology that was considered state of the art 20 years ago and considered mature technology 10 years ago is still utilised, few or no new engines of this vintage are being manufactured. They are simply seen as too inefficient, expensive and environmentally unfriendly to operate. The technology that was developed by the manufacturers as direct replacements to overcome the problems and to drive the industry forward has again now become mature in its life cycle and even more advanced technology engines have now been and are being installed in greater quantities. This new breed of technology utilises advanced technology materials, cooling designs and coating systems to deliver clean and efficient power to the market. The challenges over the next 20 years will be to optimise this technology having the ability to operate to the same standards on a variety of fuels while driving down emissions levels of environmentally harmful gases. Developing these engines with all these advanced designs is only part of the story. This is the front-end design and considers the initial capital expenditure (CAPEX), but unless these engines can operate to the same
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standards for several decades the technology becomes non-cost-effective. The operational expenditure (OPEX) plays a significant part in the overall decision making of an owner operator and a fundamental part of this is the ability to maintain the engine and repair the components. Improvements in repair technology has therefore had to keep up with the advancements in design, taking into account not only the materials, coatings and cooling systems but also having to deal with the operational variations and degradation mechanisms that occur due to locally driven power agreements. Operators are always looking to reduce these OPEX costs and therefore increasing the life of the components. Increasing the maintenance intervals and improving the operational integrity of the overall plant are just some of the key topics constantly under review.
10.10 Conclusion The maintenance and repair cycle are extremely complex in their own right and are totally integrated as a system. Overextending maintenance intervals in order to reduce costs has a negative impact on the optimised life of the component and also leads to increased repair costs as degradation levels are increased. Alternatively, removing parts prematurely can reduce the repair costs and optimise the component life but increase the maintenance costs due to more frequent maintenance intervals being required. Getting the correct balance is not easy and it is incumbent on the customer and supplier working in partnership to achieve the optimised solution.
10.11 Further reading 1. Balevic D, Burger R, Forry D, Heavy-Duty Gas Turbine Operating and Maintenance Consideration, GE Energy GE3620k. Available from www. geenergy.com. 2. Bernstein H L (1999) High Temperature Coatings for Industrial Gas Turbine Users: Proceedings of the 28th Turbomachinery Symposium, Texas A&M University. 3. Boyce M.P (2002) Gas Turbine Engineering Handbook. 2nd Edition. Available from http://www.gulfpp.com. 4. Crimi P V, Brezner J K, Grossman T R, Nash A G, Porisch A J, Gas Turbine Repair Technology, GE Global Apparatus Services Department, Houston, Texas, GER 3957. Available from www.geenergy.com. 5. Strang A, Banks W M, McColvin G M, Oakey J E, Vanstine R W (2007) Power Generation in an Era of Climate Change. Proceedings of the Seventh International Charles Parsons Turbine Conference. 6. Strang A, Banks W M, Conroy R D McColvin G M, Neal J E, Simpson S (2000) Advanced Materials for 21st Century Turbines and Power Plant: Proceedings of the Fifth International Charles Parsons Turbine Conference. 7. Wood M I (1999) Developments in Blade Coatings: Extending the Life of Blades? Reducing Lifetime Costs? CCGT Generation, IIR Ltd.
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11 Steam oxidation in steam boiler and turbine environments G . R . H O L C O M B , National Energy Technology Laboratory, U.S. Department of Energy, USA
Abstract: Steam oxidation is important to the successful long-term operation of steam boilers and turbines. The detrimental consequences of oxide scale formation are section loss, reduced heat transfer, and the possibility of oxide spallation. Subjects important to the understanding of steam oxidation are presented, including a description of the environment found in steam boilers and turbines, oxidation thermodynamics, oxidation kinetics, scale morphologies, and scale spallation. In most cases the discussion is focused on existing power plants, but due to interest in advanced ultra-supercritical (A-USC) power production, additional information will also be given for these systems. Related topics will be introduced on steam oxidation management, spallation models, and future trends. Key words: oxidation, scale, spallation, steam, boiler, turbine, superheater, reheater, ferritic steels, austenitic steels, Ni-base superalloys, advanced ultra-supercritical.
11.1
Introduction
There are two distinct high-temperature environments in coal-fired boilers – the flue gas environment, with solid and gaseous combustion products, and the water or steam environment inside the boiler tubes, boiler pipes, and the steam turbine. These environments are respectively termed fire side and steam side. Steam side oxidation is the subject discussed here. The materials of construction of boiler tubes, steam pipes, and steam turbine components are reactive enough to the low oxygen activities in the steam to form oxide scales. Dense and adherent oxide scales can be beneficial in reducing the rate of further oxidation, the oxidation rate being generally inversely proportional to the oxide thickness. The detrimental 453 © Woodhead Publishing Limited, 2011
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consequences of oxide scale formation are section loss, reduced heat transfer, and the possibility of oxide spallation. Section loss results in a corresponding reduction in component strength, and so can eventually lead to component failure. Reduced heat transfer, due to the insulative nature of the oxide scale, can lead to increased fire side boiler tube temperatures to maintain constant steam temperatures – and so indirectly lead to increased fire side corrosion. Oxide spallation is detrimental in three ways. First, it leads to increased oxidation rates (oxidation is generally proportional to oxide thickness). Second, the spalled oxide can erode downstream components, such as turbine blades. Lastly, if spalled oxide collects within a boiler tube (such as at the bottom bend of a superheater pendant), then steam flow can be blocked or reduced, which can lead to excessive metal temperatures downstream of the blockage and failure of the tube. The sections that follow will describe the environment found in steam boilers and turbines, oxidation thermodynamics, oxidation kinetics, and scale morphologies and spallation. In most cases the discussion will focus on existing power plant conditions, but due to interest in advanced ultrasupercritical (A-USC) power production, additional information will also be given for A-USC. Then related topics will be introduced on steam oxidation management, spallation models, and future trends (ultra-supercritical steam, co-firing with biomass, and oxyfuel combustion).
11.2
Steam boiler and turbine environments
The temperatures and pressures of steam boilers and turbines have been increased to improve efficiencies in steam and power production. Improvements in materials properties such as high-temperature strength, creep resistance, and oxidation resistance have enabled this increase. From 1921 (1.9 MPa and 293 8C, North Tess station, Newcastle Electric Supply Company in northern England (Stultz and Kitto, 1992, p. 7)) to 1960 (34.5 MPa and 649 8C, Eddystone Unit 1, Philadelphia Electric, Pennsylvania, (Seth, 1999)) there was an average increase in steam pressure and maximum steam temperature of 0.84 MPa per year and 9.1 8C per year. The first commercial boiler with a steam pressure above its critical value of 22.1 MPa (3208 psi) was the 125 MW Babcock & Wilcox (B&W) Universal Pressure (UP) steam generator in 1957 at the Ohio Power Company’s Philo 6 plant (Stultz and Kitto, 1992, p. 9). A UP boiler can operate at both subcritical and supercritical conditions. It delivered steam at 34.1 MPa and 621 8C with two reheats of 566 8C and 538 8C. The pressure and temperatures of the primary, first reheat (if present), and second reheat (if present) are designated by a nomenclature such as 34.1 MPa/621 8C/ 566 8C/538 8C for the Philo 6 plant. In 1960, Eddystone 1 reached a world record in efficiency of 40 % operating at 34.5 MPa/649 8C/565 8C/565 8C
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(Seth, 1999; Viswanathan et al., 2003; Henry and Ward, 2006). The Eddystone 1 plant, and others of its generation, soon reduced operating temperatures and pressures because of thermal fatigue issues within the boiler (Seth, 1999) and coal ash corrosion in the superheater and reheater tubing (Henry and Ward, 2006). Eddystone 1 continues to operate at 32.4 MPa/610 8C (Viswanathan et al., 2003). Since the 1970s, advances in the high-temperature strength of ferritic steels have allowed increases of operating temperatures and pressures without the thermal fatigue issues of the austenitic steels that had to be used to obtain the required high-temperature strengths in the early 1960s. Ferritic steel, as used here, refers to the equilibrium structure. In practice, a martensitic or partially martensitic structure is obtained from heat-treating. Currently, the most efficient fossil power plants operate at about 24 to 28 MPa and 600 to 620 8C. Current research programs are aimed at increasing the operating conditions to as high as 760 8C and 35 MPa (Viswanathan et al., 2005). General usage (Viswanathan et al., 2005) of the terms subcritical, supercritical (SC), ultra-supercritical (USC), and advanced ultra-supercritical (A-USC) for coal power plants is as follows: subcritical at below 22 MPa, SC at above 22.1 MPa and 538 to 565 8C, USC at above 565 8C, and A-USC above temperatures where nickel-base superalloys must be used, ~675 8C.
11.2.1 Water properties Several physical properties of water undergo large changes at or near its critical point (374 8C and 22.1 MPa) (Kritzer et al., 1999). Above 374 8C water vapor can no longer be compressed into a liquid. Density and dielectric constant decrease with increasing temperature, and then fall sharply at the critical temperature. The ionic product reaches a maximum at about 300 8C and also falls sharply at the critical temperature. Supercritical water is a low-density fluid that is a non-polar solvent with high solubilities for organic compounds and gases, and low solubilities for salts.
11.2.2 Steam cycle A simplified water path in a large supercritical utility boiler, with a single reheat, is illustrated in Fig. 11.1. Input water from the boiler feed pump is preheated in the economizer and then up the water walls and into the separator. Steam from the separator is fed into the primary and secondary superheaters. From there the steam goes to the high-pressure turbine, then to the reheater, the intermediate-pressure turbine, the low-pressure turbine, and then to the condenser. From the condenser, water goes back to the boiler feed pump to then repeat the cycle.
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11.1 Simplified water path in a large supercritical utility boiler, with a single reheat. Adapted from Stulz and Kitto (1992, pp. 18–19).
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Supercritical boilers are typically once-through systems, but with the capacity to separate water from steam in the separator (particularly during startup). Boilers that operate at constant pressure in either subcritical or supercritical steam conditions (but not both) are universal pressure (UP) boilers. However, due to economic considerations, UP boilers are generally only used in supercritical steam conditions, while drum type boilers are used in subcritical steam conditions. UP boilers are economical for base load operation and load cycling operations (Smith, 1998). Benson (sliding pressure) boilers can operate in both subcritical and supercritical conditions and are widely used in Japan and Europe for frequent load cycling (with nuclear power providing more of the base load power). Benson boilers can accommodate both single- and two-phase flow (Vitalis, 2006). The Rankine cycle, in terms of temperature and entropy, is shown in Fig. 11.2 for a subcritical system and in Fig. 11.3 for an SC system. In these figures, the critical temperature is at the maximum in the saturated-liquid and -vapor curves. Under the saturated lines is a wet-mixture; to the left of the saturated-liquid line is compressed liquid and to the right of the saturated-vapor line is superheated vapor. In Fig. 11.2, liquid water is compressed from points a to b. The temperature rise (a–b) is exaggerated in Fig. 11.2, and is typically on the order of 0.6 8C. The heating that takes place
11.2 Rankine cycle for a subcritical boiler. Adapted from Stultz and Kitto (1992, pp. 2–13).
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11.3 Rankine cycle for a supercritical boiler. Adapted from Stultz and Kitto (1992, pp. 2–20).
from points b to c occurs first in the economizer, then in the water walls (the constant temperature evaporation step), and then in the superheater. Expansion takes place within the turbines from points c to d. Lastly, unavailable heat is rejected to the atmosphere from points d to a. The SC cycle in Fig. 11.3 is similar (points b to e), but does not include the evaporation step, and so energy is not needed to overcome the latent heat of evaporation. Also shown in Fig. 11.3 is a reheat step through points f to g. The unavailable heat is the area under the h–a step. The thermal efficiency in converting heat into work in reversible Rankine cycles is the ratio of the temperature–entropy area within the cycle to the combined areas of the cycle and unavailable heat. Supercritical systems allow the upper temperature to increase, which increases the cycle area and thus the thermal efficiency. The curves in Fig. 11.3 show how reheating also increases the thermal efficiency by raising the average maximum temperature.
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11.2.3 Materials requirements The components exposed to supercritical water in SC and USC power plants include high-pressure steam piping and headers, superheater and reheater tubing, water wall tubing in the boiler (Viswanathan and Bakker, 2001a; Blum et al., 2003; Staubli et al., 2003), and high and intermediate pressurerotors (HP and IP), rotating blades, and bolts in the turbine section (Viswanathan and Bakker, 2001b; Staubli et al., 2003). A synopsis of the materials requirements for each of these is presented below. Steam piping and headers These require high-temperature creep strength. They are heavy section components and are particularly subject to fatigue from thermal stresses. Ferritic steels have lower thermal expansion coefficients than austenitic steels, and so are preferred with respect to thermal fatigue. Current ferritic steels are limited to 620 8C, while the theoretical limit is thought to be about 650 8C (Viswanathan and Bakker, 2001a). Superheater and reheater tubing These require high-temperature creep strength, thermal fatigue strength, weldability, fire side corrosion resistance, and steam side corrosion resistance. Ferritic steels are preferred due to their thermal fatigue resistance. However, high-temperature creep strength limits these alloys currently to 620 8C (650 8C theoretical limit). Fire side corrosion resistance further limits ferritic steels to about 593 8C, which corresponds to a steam temperature of about 565 8C (Viswanathan and Bakker, 2001a; Staubli et al., 2003). Water wall tubing These have similar issues to superheater and reheater tubing, but at lower temperatures so that lower alloyed materials are typically used. High-pressure (HP) and intermediate pressure (IP) rotors The HP and IP rotors are large forgings that carry the rotating blades and transmit the mechanical energy. They are subject to centrifugal loads during operation and to thermal stresses during startups and shutdowns. The key materials properties are creep strength, low-cycle fatigue strength, and fracture toughness (Viswanathan and Bakker, 2001b; Staubli et al., 2003).
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Rotating blades These require high-temperature strength, creep resistance, and a coefficient of thermal expansion similar to the rotor (Viswanathan and Bakker, 2001b). In addition, the alloys must be able to be peened. In each row of blades, a circular cover is used to couple the blades together and to act as a seal. The tenon part of the blades protrudes from the cover and is peened into heads, which attach the blades to the cover (Viswanathan and Bakker, 2001b). Bolts These require high-temperature strength, creep resistance, notch sensitivity resistance, and a coefficient of thermal expansion similar to the rotor. The bolts must remain tight between scheduled shutdowns (20 000 to 50 000 hours) (Viswanathan and Bakker, 2001b).
11.2.4 Boiler alloys The alloys used (or proposed to be used) in SC and USC boilers can be broadly classified as ferritic steels, austenitic steels, and Ni-base superalloys. Ferritic steels are currently limited to temperatures below 620 8C, with a theoretical limit of about 650 8C; austenitic steels are thought to have applications from 620 to 675 8C; and Ni-base superalloys above 675 8C (Viswanathan et al., 2003). A list of nominal compositions of alloys of interest are given in the Appendix in section 11.10 at the end of the chapter. Major contributions to the development of ferritic steels for boiler applications have been made by Abe et al. (2000) and Masuyama (2001).
11.3
Oxidation thermodynamics and kinetics
11.3.1 Thermodynamics The oxidation of metals in steam follows the general reaction: xM þ yO2 ¼ Mx O2y
½11:1
Key thermodynamic issues used to describe this reaction follow.
11.3.2 Oxygen partial pressure The available O2 for equation [11.1] is either from oxygenation treatments or from the dissociation of water: H2 O ¼H2 þ 12 O2
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For utility boilers, it is unclear how much O2 from oxygenation is retained (not scavenged) by reaction with tube walls at any given point in the steam cycle. The dissociation of water in equation [11.2] gives a lower bound on the oxygen activity and is generally the expression used to estimate the oxygen activity. The free energy change of equation [11.2], ΔG28 , is given by (Kubaschewski and Alcock, 1979) DG2 ðJ=molÞ ¼ 246 440 54:8T ¼ RT ln K2
½11:3
In equation [11.3], K2 is the equilibrium constant for equation [11.2] and is given by 1=2
K2 ¼
pH2 pO2 pH 2 O
½11:4
where pi is the partial pressure, in atm, of i. Following the work of Young (2008), a good estimate of the value of pO2 in terms of the total pressure, PT, can be given by 2=3
2=3
pO2 ¼ 2K2 PT
½11:5
As equations [11.3] to [11.5] show, the partial pressure of O2 increases with both increasing temperature and pressure. Expected oxide phases The phase stability diagram for iron in water, as functions of temperature and oxygen partial pressure, is shown in Fig. 11.4. The hashed line ‘pure H2O’ refers to partial pressures of oxygen as calculated by methods similar to equation [11.5]. The hashed line ‘H2O + 200 ppb O2’ corresponds to an oxygenated treatment of 200 ppb of dissolved oxygen (Paterson et al., 1992). This diagram helps to explain the scale morphologies found on low alloy boiler steels. Typically, the scales consist of several layers that correspond to the decreasing oxygen activity found within the scale closer to the metal. Hematite (Fe2O3) may form as an outer layer. A duplex magnetite layer forms under the hematite. The outer magnetite layer is Fe3O4 and grows out from the original metal surface. The inner magnetite layer is (Fe, Cr)3O4 and grows inward from the original metal surface. The interface between the two magnetite layers corresponds to the original metal interface (Cory and Herrington, 1987). Wu¨stite (FeO) may form under the magnetite layer. Wu¨stite formation is a function of Cr content, temperature, and water vapor pressure (Grobner et al., 1980). formation is associated with an increase in oxidation kinetics and may be responsible for the transition from parabolic to linear kinetics that may occur after many thousands of hours of service (Sarver et al., 2003).
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11.4 Phase stability of iron as a function of temperature and oxygen partial pressure. Included are calculations of boiler conditions predicted from feed waters of pure H2O and H2O with 200 ppb oxygen. Adapted from Gaskell (1973) with information from Paterson et al. (1992).
Hematite formation is associated with increased spallation in austenitic alloys during shutdowns due to its large thermal mismatch with magnetite. In addition to the excess oxygen conditions of Fig. 11.4, hematite can also be formed due to a lowering of the amount of Fe arriving at the outer scale, which has been attributed to the voids that form at the base of the scale after many tens of thousands of hours of exposure (Dooley, 2003). Austenitic steels with 18Cr, such as 304 and 347, are similar, but form a thin layer of Cr2O3 under the (Fe, Cr)3O4 layer (Otsuka and Fujikawa, 1991). The Cr2O3 layer is quite protective and substantially reduces the oxidation kinetics. Above 20–25 % Cr, the morphology changes to just Cr2O3 on the surface of the alloy (Otsuka and Fujikawa, 1991; Wright et al., 2007).
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Chromia evaporation For scales that form a primarily chromia (Cr2O3) scale, additional considerations need to be considered based on chromia evaporation. For conditions in a steam boiler or turbine, the primary Cr-containing gas species is CrO2(OH)2, and forms by reactive evaporation: 1 2 Cr2 O3 ðsÞ
þ H2 OðgÞ þ 34 O2 ðgÞ ¼ CrO2 ðOHÞ2
½11:6
There has been much recent interest in chromia evaporation for detrimental effects in solid oxide fuel cells (Gindorf et al., 2001; Holcomb, 2008), aerospace (Opila et al., 2007), and steam boilers (Young and Pint, 2006; Holcomb; 2008, 2009a). The details of these investigations are beyond the scope of this chapter. A direct effect of the loss of chromia scale is subsequent faster (and linear) oxidation kinetics. An early description of this change in kinetics was developed by Tedman (1966), where a point in time is reached where the rate of scale thickness gain from oxidation matches the rate of scale thickness loss from evaporation. This results in linear oxidation kinetics. The faster kinetics can also lead to depletion of Cr within the alloy, which can result in breakaway oxidation (Asteman et al., 1999; 2000; Holcomb, 2009a). The rate of chromia evaporation in steam boilers and turbines is expected to be quite low, due to the low partial pressures of oxygen and the fact that most of the alloys do not form pure chromia scales. However, it might still be an issue in A-USC turbines, with the high temperatures, pressures, and steam velocities all combining to result in enough evaporation to be a concern (Holcomb, 2009a).
11.3.3 Kinetics The growth of oxide scales in steam is generally controlled by ionic transport within the oxide scale. Diffusion controlled oxidation is inversely proportional to the scale thickness and so gives rise to parabolic kinetics: dX kp ¼ dt X
½11:7
where X is scale thickness, t is time, and kp is the parabolic rate constant. Integration gives kp as kp ¼
X2 2t
½11:8
Note that in some cases in the literature, the ‘2’ in equation [11.8] is included with kp to give a slightly different definition. The integration to obtain equation [11.8] neglects the integration constant, which in terms of oxidation corresponds to the initial oxidation when the kinetics might be
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different from parabolic. In the long-term exposures typical of boiler and turbine systems, the amount of initial oxidation is usually neglected. The temperature dependence of kp is usually given as a thermally activated process with an activation energy, Q, and an Arrhenius constant, A: Q kp ¼ A exp ½11:9 RT Industrial service measurements of kp usually involve scale thickness measurements. However, mass change measurements are much more commonly reported in laboratory measurements. In a similar fashion as in equation [11.8], the parabolic rate constant from gravimetric measurements, kw, is given by 1 DW 2 kw ¼ ½11:10 2t SA where W is the weight or mass and SA is the surface area. Translation between kp and kw is possible if the scale composition and porosity is known. A comprehensive (and publicly available) report by Wright et al. (2007) on oxide growth and scale adhesion has collected, compiled, and to some extent analyzed available industrial service and laboratory steam oxidation data. Table 11.1 shows selected results from the compiled oxidation data from sources that used scale thickness measurements. Table 11.2 shows similar data from sources that used mass change measurements. These data are plotted in Figs 11.5 and 11.6. The scale thickness data in Table 11.1 and Fig. 11.5 for T22 (industrial service) and 2.25Cr–1Mo (laboratory data) show consistently higher corrosion rates for industrial service than in the laboratory, yet the similar activation energies point to the same oxidation mechanisms occurring in both cases. Industrial service factors such as the thermal gradient from heat flux, high pressure, and thermal cycling may act to increase the corrosion rate. The lines for 9Cr and 12Cr are from a compilation (Montgomery and Karlsson, 1995) of steam oxidation data, and compared to other results compiled by Wright et al. (2007) show an upper limit to the corrosion data – useful as a conservative estimate of oxidation rates. The gravimetric data in Table 11.2 and Fig. 11.6 show the wide variation in oxidation behavior of different alloy classes. This will later be shown to be related to the type of scales that form. The 0–2Cr results show a change in activation energy at 550 8C. This corresponds to the formation of Wu¨stite (FeO) Fig. 11.4. Wu¨stite formation is suppressed by Cr additions, and so this change in oxidation behavior at 550 8C is not observed for 9Cr alloys or for 300 series austenitics.
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11.5 Selected data adapted from the compilation in Wright et al. (2007) for steam oxidation kinetics based on scale thickness measurements, with parameters listed in Table 11.1.
11.6 Selected data adapted from the compilation in Wright et al. (2007) for steam oxidation kinetics based on mass change measurements, with parameters listed in Table 11.2.
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Table 11. 1 Selected results adapted from the compilation of Wright et al. (2007) for oxidation data in steam from scale thickness measurements. Data are from industrial service except where noted. Originally, the T22 data were from Paterson et al. (1992) and the 9Cr and 12Cr data were from an earlier compilation (Montgomery and Karlsson, 1995) Alloy or alloy type Temperatures (8C) A (μm2/h) Q (kJ/mol) Notes T22 2.25Cr–1Mo 9Cr 12Cr
482–621 550–700 500–750 525–775
6.2261020 326 2.0561020 327 3.9761011 197 7.9661010 188
Pressure of 14.6 MPa Laboratory data
Table 11 2 Selected results adapted from the compilation of Wright et al. (2007) for oxidation data in steam from mass change measurements. Data are from laboratory measurements. Data are compiled from a best fit from several sources Alloy or alloy type Temperatures (8C) Pressures (MPa) A (g2/cm4/s) Q (kJ/mol) 0–2Cr 0–2Cr 9Cr 300 series Austenitics Alloy 800 Ni-base alloys
<550 550–700 450–700 538–760 482–760 650–800
0.1–24.1 0.1–24.1
0.66 2.846108 1.18 1.646102
172 309 195 229
10.5–24.1 1.7–24.1
1.766103 1.286107
186 122
Pressure Paterson et al. (1992) used in-plant measurements to compare the oxide scale thickness from superheater tubes and reheater tubes. The oxide scales were 45 % thicker on the superheater tubes. Using the bulk steam pressures (P) experienced by the two types of tubing, it was calculated that the oxide growth rate was proportional to P1/5. It was surmised that the pressure effect was due to differences in the partial pressure of oxygen at the different pressures. For example, using the STANJAN computer code (Reynolds, 1986), the oxygen partial pressure of steam at 593 8C derived from 200 ppb dissolved oxygen feed water was calculated to be 1.916105 and 2.756105 at 17.2 MPa and 24.8 MPa, respectively. Grain size It is well established that many alloys develop and maintain more protective oxide scales when the metal has a smaller grain size. The grain size can be smaller due to heat treatment or from cold work (for example shot-peening of the surface). At the temperatures of interest, less than 800 8C, much of the diffusion of Cr within the metal occurs along grain boundaries rather than
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as lattice or transgranular diffusion. In terms of Arrhenius diffusion behavior, the activation energy for grain boundary diffusion in single-phase alloys is typically half that of the activation energy for diffusion through the lattice (Kaur et al., 1995); thus grain boundary diffusion predominates at lower temperatures. Increased Cr diffusion allows for scales with more Cr oxides with lower ionic diffusion kinetics for oxygen and metals. Silicon Silicon is always present in the ferritic alloys used in steam boilers. It is generally thought to improve the oxidation resistance. In a study on Fe– 9Cr–1Mo–xSi, Fe–12Cr–1Mo–xSi, and Fe–xCr–1Mo–0.4%Si alloys with tight controls on alloying additions by Fukuda et al. (1995), Si was found to about a factor of four more protective than Cr. A regression analysis of their results for Q and log A found: Q ¼ 333:2 17:4½Cr 69:0 ½Si
½11:11
log A ¼ 15:9 1:1½Cr 4:5 ½Si
½11:12
The major difference in the scales that formed on the low Si (0.2%Si) and the high Si (0.8% Si) alloys was that a continuous Cr- and Si-rich sublayer formed on the high Si steels.
11.4
Scale morphology and spallation
The morphologies of the scales that form play an important role as they can reveal, to a certain extent, the oxidation mechanism, oxidation kinetics, and the type of and tendency towards scale spallation.
11.4.1 Ferritic steels The oxides that form on ferritic steels in steam typically have a duplex structure, with the inner scale being Fe–Cr spinels and the outer scale being essentially pure Fe3O4. The inner scale is typically fine grained, while the outer scale is typically columnar grained. The inner and outer scales are usually similar in thickness. The interface between the two layers represents the location of the original steel surface, although with the volume expansion associated with oxidation, it may be shifted outwards. Sometimes a thin or discontinuous layer of Fe2O3 is found at the scale– gas interface. This type of structure is shown in Fig. 11.7 for T-22 (Dooley and McNaughton, 1995). Much of the scale structure arises from differences in the oxygen activity
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11.7 Schematic drawing of a typical oxide morphology formed on T22 in steam. Adapted from Dooley and McNaughton (1995).
required for oxidation of Fe and Cr. Chromium is much more reactive to oxygen than Fe, so it oxidizes at a much lower oxygen activity. At the scale– gas interface, Fe is oxidized to Fe3O4 (neglecting Fe2O3 formation for the moment) at the modest oxygen activity within the steam that arises from water dissociation (equation [11.2]). Iron is transported outwards through the scale to participate in that reaction. This can cause porosity to occur in the outer scale, near the boundary between the inner and outer scales. Oxygen is transported in through the scale where it reacts with Cr at the metal–oxide interface or as internal oxidation within the metal. The amount of Cr in the inner scale increases with time (as well as with alloy content), as internally oxidized Cr is engulfed within the scale as the metal recedes. Initially, the transport of Fe outwards is fast enough to sustain Fe3O4 formation at the scale–gas interface. However, as the Cr content in the outer scale increases, the transport of Fe slows. At some point, it may slow enough to allow the formation of Fe2O3 at the scale–gas interface. Ferritic steels containing up to 9Cr (with the exception of T91 and other newer 9Cr ferritics) Industrial exposures (with a heat flux and rapid heating and cooling during startups and shutdowns) can modify the above description to give rise to multiple oxide layers near the oxide–metal interface (Manning and Meadowcroft, 1980; Rehn, 1981; Wright et al., 2007). Essentially, a new duplex scale forms at the metal surface. Presumably this occurs after separation of the scale from the metal, possibly as a result of stress relief. This process may repeat, forming multiple duplex layers. The boundaries of these duplex layers are the primary locations for scale spallation. Flakes of spalls are typically small and invariably found to consist of multilaminated layers of Fe3O4 and Fe–Cr spinel (Wright et al., 2007). The spalled flakes of scale can be an erodent to downstream components such as turbine blades.
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T91 The multilaminate scales described above are rarely encountered in T91 (Dooley, 2003). The scales are generally found to be duplex scales with an inner Fe–Cr spinel layer and an outer Fe3O4 layer. The higher Cr level in the alloy results in more Cr in the Fe–Cr spinel, which results in slower outward diffusion of Fe and a thicker Fe2O3 layer. Many times the inner Fe–Cr spinel layer will also have multiple partial layers of Cr-rich oxide. Laboratory exposures in steam produce very similar microstructures as inservice exposures, but many times will lack the Fe2O3 layer. An example of this is shown in Fig. 11.8 for T91 after 4000 hours of exposure in steam at 1.7 MPa at 650 8C, where it shows the typical inner Cr-rich spinel layer and the outer scale of Fe3O4. Internal oxidation of Cr is clearly seen in the higher magnification cross-section. Delamination or spalling of T91 usually occurs along the inner oxide–outer oxide boundary, or rarely at the metal–scale interface (Wright et al., 2007).
11.8 Cross-sections of scales formed on T91 after 4000 hours in steam at 1.7 MPa at 650 8C (micrograph by courtesy of Bruce A. Pint, Oak Ridge National Laboratory) (Viswanathan et al., 2007). Typical of ferritics steels are the inner Cr-rich spinel layer and the outer scale of Fe3O4. Internal oxidation of Cr is clearly seen in the higher magnification cross-section. (Continues on next page)
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11.8
(continued)
12Cr ferritics Scales on 12Cr ferritics can show wide variations, ranging from very thin and protective scales to those similar to what is found on T91. This variation reflects the closeness of the Cr content to the threshold to form Cr2O3. Minor alloying elements and environmental conditions can influence which type of scales are formed. When scales on 12Cr ferritics are similar to those on T91, they typically have even thicker Fe2O3 layers (Wright et al., 2007). Spallation tends to occur along the inner scale–outer scale boundary, and is associated with regions containing thicker Fe2O3 layers than nearby regions than have not spalled. A cross-section of the alloy MARB-2 (essentially 9Cr–0.7Si–2.5W–CoV, the high Si content effectively puts it into the 12Cr category) after 4000 hours of exposure in steam at 1.7 MPa at 650 8C is shown in Fig. 11.9. In contrast with Fig. 11.8, Fe2O3 was formed even in a laboratory exposure. Other than the Fe2O3 layer, the microstructure seen in Fig. 11.9 is very similar to that in Fig. 11.8.
11.4.2 Austenitic steels The oxides that form on austenitic steels in steam typically have a duplex structure, but with a morphology that is different from that of ferritic steels. The morphology is illustrated in Fig. 11.10. The inner scale is an Fe–Ni–Cr
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11.9 Cross-section of scale formed on MARB-2 after 4000 hours in steam at 1.7 MPa at 650 8C (micrograph courtesy of Bruce A. Pint, Oak Ridge National Laboratory) (Viswanathan et al., 2007). Note the presence of hematite (Fe2O3), as indicated by the lighter shade of oxide along the outer oxide edge.
spinel that is usually compact and with an irregular shaped interface with the metal. More protective Cr2O3 scales form at grain boundaries, and so more metal recession occurs between the grain boundaries; this leads, in turn, to an irregular and convoluted interface. As with ferritic steels, the outer portion of the scale is mainly Fe3O4 and Fe2O3 (the Fe2O3 located at the scale–steam interface) with voids and a columnar grain structure. The interface between the two layers represents the location of the original steel surface. There are also cases in which large portions of the surface remain protective, but nodules of two-layered oxides grow as isolated islands. An example crosssection is shown in Fig. 11.11 for Super304H after 4000 hours of exposure in steam at 1.7 MPa at 650 8C. Nodules of two-layered oxides can grow laterally to cover larger portions of the surface. Fine grain (grain size smaller than ASTM 8) austentic steels exhibit much greater oxidation resistance through the formation of a continuous layer of Cr2O3 (Matsuo et al., 2005). The irregular oxide–metal boundary confers some additional adhesion strength to the oxide, so the usual delamination location is at the boundary between the inner and outer scales, or along the nearby porous layer in the outer scale. Spallation usually occurs as large flakes that consist of the entire
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11.10 Schematic drawing of a typical oxide morphology formed on 300 series austenitic steels in steam. Adapted from Dooley and McNaughton (1995).
11.11 Cross-section of scale formed on Super304H after 4000 hours in steam at 1.7 MPa at 650 8C (micrograph courtesy of Bruce A. Pint, Oak Ridge National Laboratory) (Viswanathan et al., 2007). Note the contrast between regions where a protective Cr-rich spinel forms and where its protection has been lost.
outer layer (Fe3O4 + Fe2O3), many times with Fe3O4 on one side of the flake and Fe2O3 on the other side of the flake (Wright et al., 2007). These large flakes can collect at low portions of superheater and reheater arrays, reduce the amount of steam flow, and lead to local overheating of the tube. When austenitic steels spall, it usually occurs during cooling. Austenitic steels with greater than 20–25 % Cr can form a continuous surface layer of Cr2O3 (Otsuka and Fujikawa, 1991; Wright et al., 2007). This type of scale is very protective and results in much lower oxidation rates.
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11.4.3 Ni-base superalloys Nickel-base superalloys proposed for use in A-USC boiler and turbine systems typically have upwards of 20–25 % Cr, which means that they form essentially pure Cr2O3 scales. Minor amounts of Mn, Si, and Ti are also found in Cr2O3 scales. Most of these alloys also contain Al and Ti, which internally oxidize. The internal oxidation of Al and Ti can occur at much greater depths (especially along grain boundaries) than the metal lost in forming the scale, so reports of oxidation behavior should include both scale thickness and internal oxidation depths. An example for Inconel 740 is shown in Fig. 11.12 after 4000 hours of exposure in steam at 1.7 MPa at 800 8C. It has been speculated that should chromia evaporation be significant in A-USC turbines, as one model suggests (Holcomb, 2009a), then the outer oxide surface may become enriched in TiO2 and Cr–Mn spinels as Cr2O3 is lost. The enrichment of the outer oxide surface with TiO2 and Cr–Mn spinels would lower the activity of Cr2O3 and thus serve to lower the evaporation rate – at least for alloys with significant Ti and Cr. It should be noted that oxides of Ti can be found both as internal oxides and at the scale–steam interface, facilitated by the multiple oxidation states of Ti. Nickel-base superalloys with significant Al (above about 4%Al), may have sufficient Al to form Al2O3, which is very protective. However, alloys with significant Al are notoriously difficult to weld and fabricate, so are generally not proposed to be used.
11.12 Cross-section of scale formed on Inconel 740 after 4000 hours in steam at 1.7 MPa at 800 8C (micrograph courtesy of Bruce A. Pint, Oak Ridge National Laboratory) (Viswanathan et al., 2007). Internal oxidation predominates, in terms of effective metal loss, over scaling.
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11.4.4 Spallation models One of the most useful approaches to model spallation is that of Armitt et al. (1978), which mapped the scale integrity–spallation boundary with respect to scale thickness and strain; Fig. 11.13. As scale thickness increases, the total amount of elastic strain that can be accommodated decreases, thus narrowing the band of scale integrity. Thickness measurements of scales that showed significant spallation are superimposed on to the Armitt diagram in Fig. 11.14. Ferritic steels spall in tension, while austenitic steels spall in compression. Scales on austenitic steels spall at much lower scale thicknesses than do ferritic steels. The role of hematite (Fe2O3) is clearly shown in Fig. 11.14. Hematite adds compressive stresses, which tend to help the ferritic steels and hurt the austenitic steels. Differences in thermal expansion between the steel and oxide scales are the source of much of the stresses and strains that lead to spallation. Coefficients of thermal expansion (CTE) are shown as a function of temperature in Fig. 11.15. Ferritic steels such as T22 have CTEs between that of Fe2O3 and Fe3O4 or FeCr spinels, while austenitic steels such as 316 have CTEs largely above that of the oxides. The Armitt type approach is in the process of being expanded (Wright et al., 2004, 2007; Schu¨tze et al., 2010) by replacement of the scale thickness with a parameter, ω, that summarizes all of the physical and chemical changes that occur to the scale during operation including relevant plant operation data.
11.13 ‘Armitt diagram’, a map of the scale integrity–spallation boundary with respect to scale thickness and strain. Adapted from Armitt et al. (1978).
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11.14 Thickness measurements and strains of scales that showed significant spallation superimposed on to the Armitt et al. (1978) diagram. Adapted from Dooley and Paterson (2003).
11.15 Coefficients of thermal expansion (CTE) for relevant oxides and alloys. The rapid increase in CTE for Fe3O4 is due to a spinel structure transformation. Adapted from Osgerby and McCartney (2002).
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11.5
Steam oxidation management
Management of systems exposed to steam is essential for successful longterm operations. Many of the failure modes of boiler tubes, steam piping, and steam turbines do not directly involve steam oxidation. Hightemperature creep, thermal fatigue, erosion, and fire side corrosion are all examples of this. Management of systems with respect to steam oxidation can be divided into groups: items associated with outage maintenance and items associated with operations.
11.5.1 Outage maintenance During an outage, inspections of superheater tubes, reheater tubes, steam pipes, and turbine sections may include many types of examinations. The two main non-destructive tests relevant to steam oxidation are visual examinations and ultrasonic testing. Operational data of system upsets or other problems may focus additional scrutiny on effected components. Visual examination Visual examination may pick up many different issues, one being steam side surface deposits (Stultz and Kitto, 1992, p. 45-1), which can cause local overheating and rapid oxidation or loss of mechanical strength. Detection of surface deposits could trigger cleaning actions or modifications to steam chemistry monitoring or control. Ultrasonic testing (UT) In UT of tubing, a piezoelectric transducer is placed in contact with the outside diameter of the tube, which causes an elastic sound wave that moves through the material (Bar-Cohen and Mal, 1989). This sound wave is reflected by discontinuities it encounters. The reflection is received back at the transducer and the travel time measured. For metal thickness measurement, a high-frequency (2 to 5 MHz) signal is used that is sent perpendicular to the surface (Stultz and Kitto, 1992, p. 45-2). For oxide thickness measurement, a higher frequency is required to differentiate between the inside metal surface and the oxide surface (Stultz and Kitto, 1992, p. 45-2). Only relatively thick oxides can be detected, so UT is generally not used to measure oxide thickness on stainless steels. Surface deposits can also be detected with UT, again with a high-frequency signal. However, in this case signal attenuation is measured instead of travel time (Stultz and Kitto, 1992, p. 45-2).
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11.5.2 Operations While operating (and including startups and shutdowns), much attention is given to maintaining operations within the design envelope. Monitoring of temperatures, pressures, and water chemistry can all be important with respect to steam oxidation. Temperature and pressure are important both in terms of absolute values, but also in terms of rates of change and their effect on the stress level and stress relaxation within the oxide scale. Steam cycle chemistry The control of water chemistry in the steam cycle of a power plant is important for corrosion control, deposition prevention, and higher cycle efficiency (Jonas, 2000). The amount of dissolved oxygen in the steam can define the upper bound of oxygen activity within the system at temperatures below about 700 8C (the lower bound being given by the dissociation of water, equation [11.2]). There are three basic steps in controlling the water chemistry: removal of impurities (in makeup water, feedwater, and condensate), control of feedwater pH, and boiler water treatments (for once-through systems these are feedwater treatments). Impurity removal is by boiler blowdown, mechanical deaeration in the condenser and deaerator, and oxygen scavenging (usually with hydrazine) (Jonas, 2000). For once-through systems, demineralization (condensate polishing) is also performed. Feedwater pH is normally controlled by injection of ammonia or amines to maintain feedwater pH between 9.2 and 9.6, as measured at room temperature (8.8 to 9.2 when copper alloys are also present in the feedwater system) (Jonas, 2000). High pH values help protect carbon steel components at low temperatures. Ammonia and amines are weak bases, so have little effect in raising the pH at high temperatures. In the absence of strong bases, such as phosphate or hydroxide, the pH at elevated temperatures is close to neutral. In subcritical boilers, a variety of boiler water treatments (usually containing phosphates or hydroxides) are used that are not employed in higher pressure systems. For higher pressure once-through systems, either an all-volatile treatment (AVT) or oxygenated treatment (OT) is used (Jonas, 2000). In the AVT, the pH is controlled as described above, with ammonia or volatile amines. There are no phosphates or buffering salts, so there is little tolerance for contamination and thus condensate polishers are required. The AVT was first used in supercritical boilers. In OT, ultra-pure feedwater is required (< 0.15 μS/cm cation conductivity). Oxygen, either in the form of O2 gas or hydrogen peroxide, is added to the feedwater at a controlled rate. The oxygen promotes the formation of a dense and protective iron oxide film on the boiler and feedwater piping
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(Jonas, 2000). The use of OT minimizes the dissolution of iron from the steam circuit, thus preventing much of the deposition of insulating ironoxide layers on heat transfer surfaces that would otherwise occur. Condensate polishers are required in OT programs. Oxygenated treatment can be used alone, termed neutral water treatment (NWT), with oxygen controlled to ~200 ppb, or in combination with ammonia, termed combined water treatment (CWT), with oxygen controlled between 50 and 200 ppb and with ~200 ppb ammonia.
11.6
Future trends
There are many current and future developments related to fossil fuel power production, most of them involving the reduction or elimination of greenhouse gas emissions such as CO2. These include co-firing with biomass, oxyfuel combustion, and A-USC boilers and turbines. All are elements in worldwide research into carbon capture and storage (CCS).
11.6.1 Advanced ultra-supercritical (A-USC) Improved efficiencies, which result in less CO2 being generated per unit of energy produced, are the goal of A-USC boiler and turbine development. For reduced CO2 emissions, calculations by Booras et al. (2003) indicate that a subcritical 37 % efficient plant 500 MW plant burning Pittsburgh #8 coal would produce about 850 tons of CO2 per GW h. Ultra supercritical plants at 43 % and 48 % efficiency would respectively produce about 750 and 650 tons of CO2 per kW h. Less CO2 produced by improved efficiencies will reduce subsequent CCS costs. The increased temperatures and pressures required for A-USC operations will have a direct impact on steam oxidation in terms of higher operating temperatures, but the larger indirect impact will be the use of Ni-base superalloys, more use of austenitic alloys instead of 9–12Cr ferritics, and more use of 9–12Cr ferritics instead of 0–2Cr ferritics. Most of the driving force behind the use of more highly alloyed steels is the need for improved high-temperature creep strength. In many cases the oxidation behavior may be improved for the same type of component operating at a higher temperature but with a more high-temperature creep-resistant alloy. Recent and on-going international materials research has led to numerous new alloys for steam boilers and turbines and an increase in steam temperatures. Much of the efforts aimed at temperatures up to 650 8C are by improving on the ferritic steel 12Cr–1MoV. Major industrial research and development efforts in Japan, the USA, and Europe are briefly outlined in Table 11.3.
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Table 11.3 Major international research and development efforts (Staubli et al., 2001; Blum and Hald, 2002; Viswanathan et al., 2007; Fukuda, 2009). When two temperatures are listed, they refer to superheater and reheater temperatures Research effort
Time span Targets
Notes
EPDC, Japan 1981–2000 30.0 MPa Materials development and Electrical Power 630 8C/630 8C component manufacture Development with 50 MW pilot plant Company operations NIMS STX-21, Japan 1997–2006 650 8C Ferritic steel development National Institute for with a focus on fine Materials Science grained structures Part of the Cool Earth– NIMS A-USC, Japan 2008–2016 35 MPa 700 8C Innovate Energy Technology Program EPRI, USA 1978–2003 Boiler and turbine thickElectric Power walled components, Research Institute standardization, and trial of components in service. Validated NF616 (ASME P92) and HCM12A (ASME P122) U.S. DOE/OCDE USC 2001–2012 35 MPa Government and industry Boiler Consortium, 760 8C boiler consortium for USA materials development and U. S. Department of qualification Energy/Ohio Coal Development Office COST 501, Europe 1986–1997 530 8C/565 8C Turbine and boiler materials Co-Operation in the development for all major Field of Science and components. Technology COST 522, Europe 1998–2003 30 MPa Turbine and boiler materials 620 8C/650 8C development for all major components. THERMIE AD700, 1998–2013 35 MPa Materials development and Europe 700 8C/720 8C qualification, component design, and demonstration plant. Source: Adapted and updated from Halcomb (2006).
11.6.2 Co-firing with biomass The addition of a renewable fuel, such as biomass, to the feedstock of coal power plants can reduce the net amount of generated CO2. Most of the corrosion issues involved with co-firing with biomass are fire side corrosion related. Biomass tends to have different levels of corrosive species than coal has, some higher and some lower, depending on the coal and the biomass. Potassium is typically much higher in biomass fuels than in coal. In a similar manner as in A-USC, the effects from co-firing with biomass
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will largely be indirect. Different alloys may be required for fire side corrosion resistance, and thus the steamside oxidation will be changed.
11.6.3 Oxyfuel combustion Another approach towards CCS is the use of oxygen instead of air as the oxidant for burning coal (or natural gas). The CO2 in the resulting flue gas should be much easier to purify without the N2 found in air. For oxyfuel combustion of coal, the elimination of most of the N2 increases the activity of corrosive gas species such as SO2/SO3 and HCl, the increase arising from a lack of dilution from N2. This should increase fire side corrosion rates. However, there are design decisions that may lower the activities of corrosive gas species back towards that seen in air combustion. In most schemes to refit existing coal boilers for oxyfuel combustion, a certain fraction of CO2 rich flue gas is recirculated back into the boiler. This is to maintain similar heat transfer characteristics within the boiler as was designed and built for air-fired systems. If the recirculation occurs prior to flue gas desulfurization (FGD), then the corrosive gas species will remain what they were without flue gas recirculation.1 If the recirculation occurs after FGD, then the concentration of corrosive gas species in the boiler would be reduced by the addition of ‘clean’ flue gas and could approach that in air-fired systems. However, recirculation after FGD comes with a much larger efficiency loss as the FGD process produces cold and wet gas, which will use up energy to heat back up and to dry. In a similar manner as in co-firing with biomass, the effects from oxyfuel combustion in refitted boilers on steam oxidation will largely be indirect. Different alloys may be required for fireside corrosion resistance, and thus the steam side oxidation will be changed. Emerging designs for newly built oxy fuel combustion power stations (as opposed to the refit designs discussed above) may eliminate steam tubes and pipes altogether. One such design (Anderson et al., 2009) uses the oxy fuel combustion flue gas to drive turbines directly. The resulting flue gases, at least with natural gas as the fuel, are largely H2O, CO2, and O2. Oxidation in this environment may be more aggressive than in steam, CO2, or air alone. For example, it was found that T91 and T92 oxidized about a factor of 5 faster in H2O–10%CO2–0.2%O2 than in steam at 630–650 8C (Holcomb 2009b). 1 A common misunderstanding is to believe that the recirculation of pre-FGD flue gas increases the concentration of corrosive gas species, such as SOx, in the boiler. However, the main reaction within the boiler is C + O2(g) = CO2(g) which has the same number of moles of gases as products and reactants. Therefore recirculated pre-FGD flue gas contains the same concentration of SOx as is already found in the oxyfuel boiler – the SOx concentration is increased compared to air-firing because of the lack of dilution from N2, not due to the recirculation of flue gas.
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Steam oxidation in steam boiler and turbine environments
11.7
481
Conclusion
Steam oxidation is important to the successful long-term operation of steam boilers and turbines. The detrimental consequences of oxide scale formation are section loss, reduced heat transfer, and the possibility of oxide spallation. Background information was given on the environment found in steam boilers and turbines. The thermodynamics of steam oxidation were described, with emphasis on the phases that are predicted to form. Industrial service and laboratory derived kinetics of steam oxidation were reviewed. Much more information is available for widely used ferritic steels than for austenitic steels or Ni-base superalloys. Scale morphology development was examined in terms of how it relates to oxidation kinetics and scale spallation. Scale spallation models were introduced. Additional information was given on steam oxidation management and future trends.
11.8
Sources of further information and advice
Power plant operations: . .
Stultz, S. C., and Kitto, J. B. (1992), Steam, 40th ed., Barberton, Ohio: Babcock & Wilcox. Fryling, G. R. (ed.) (1967), Combustion Engineering, 2nd ed., New York, Combustion Engineering.
Alloy development for power plants: .
Masuyama, F. (2001). History of Power Plants and Progress in Heat Resistant Steels, ISIJ International, 41(6), 612–625.
High-temperature oxidation and corrosion (in general): . .
Kofstad, P. (1988). High Temperature Corrosion, New York: Elsevier. Young, D. (2008). High Temperature Oxidation and Corrosion of Metals, San Francisco: Elsevier.
Oxidation in steam: . . .
Fry, A., Osgerby, S., and Wright, M. (2002), Oxidation of Alloys in Steam Environments – A Review, NPL Report MATC(A)90, Teddington, England: National Physical Laboratory. Wright, I. G., Tortorelli, P. F., and Schu¨tze, M., (2007), Program on Technology Innovation: Oxide Growth and Exfoliation on Alloys Exposed to Steam, Report 1013666, Palo Alto, Colifornia EPRI. Wright I. G., and Dooley, R. B. (2010), A Review of the Oxidation Behavior of Structural Alloys in Steam, International Materials Reviews, 55(3), pp. 129–167.
© Woodhead Publishing Limited, 2011
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Power plant life management and performance improvement
Advanced ultra-supercritical (A-USC) boiler development: .
.
.
.
. .
Abe, F., Igarashi, M., Wanikawa, S., Tabuchi, M., Itagaki, T., Kamura, K., and Yamaguchi, K. (2000), R&D of Advanced Ferritic Steels of 650 8C USC Boilers, in Proceedings of the 5th International Charles Parsons Turbine Conference, Strang, A., Banks, W. M., Conroy, R. D., McColvin, G. M., Neal, J. C., and Simpson, S. (eds) London: IOM Communications Ltd., pp. 129–142. Holcomb, G. R. (2006), Corrosion in Supercritical Water – Ultrasupercritical Environments for Power Production, in: ASM Handbook Volume 13C: Corrosion: Environments and Industries, Cramer, S. D., and Covino Jr., B. S. (eds), Materials Park, Ohio: ASM International, pp. 236–245. Oakey, J. E., Pinder, L. W., Vanstone, R., Henderson, M., and Osgerby, S. (2003), Review of Status of Advanced Materials for Power Generation, Report No. COAL R224, DTI/Pub URN 02/1509, London: Department of Trade and Industry. Staubli, M., Mayer, K.-H., Kern, T. U., Vanstone, R. W., Hanus, R., Stief, J., and Schonfeld, K.-H. (2001), COST 522 – Power Generation into the 21st Century: Advanced Steam Power Plant, in Proceedings of 3rd Conference on Advances in Material Technology for Fossil Power Plants, Published as Institute of Materials Book 770, London: Institute of Materials. Viswanathan, R., Sarver, J., and Tanzosh, J. M., (2006). Boiler Materials for Ultra-Supercritical Coal Power Plants – Steamside Oxidation, Journal of Materials Engineering and Performance, 15, pp. 255–274. Viswanathan, R., Coleman, K., Shingledecker, J, Sarver, J., Stanko, G., Borden, M., Mohn, W., Goodstine, S., and Perrin, I. (2007), Boiler Materials for Ultrasupercritical Coal Power Plants, Phase1. Final Summary Report, October 1, 2001–December 31, 2006, U.S. Department of Energy Grant Number DE-FG26-01NT41175 and Ohio Coal Development Office Grant Agreement D-00-20.
11.9
References
Abe, F., Igarashi, M., Wanikawa, S., Tabuchi, M., Itagaki, T., Kimura, K., and Yamaguchi, K. (2000), R&D of Advanced Ferritic Steels for 6508C USC Boilers, in: Proceedings of 5th International Charles Parsons Turbine Conference Cambridge: England, pp. 129–142. Anderson, R. E., Viteri, F., Hollis, R., Hebbar, M., Downs, J., Davies, D., and Harris, M. (2009), Application of Existing Turbomachinery for Zero Emissions Oxy-Fuel Power Systems, in Proceedings of ASME Turbo Expo 2009: Power for Land, Sea and Air, GT 2009-59995 (Orlando, Florida, June 8–12, 2009), New York: ASME. Armitt, J., Holmes, R., Manning, M. I., Meadowcroft, D. B., and Metcalfe, E.
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Steam oxidation in steam boiler and turbine environments
483
(1978), The Spalling of Steam-Grown Oxide from Superheater and Reheater Tube Steels. FP-686, Palo Alto, California: EPRI. Asteman, H., Svensson, J. E., Johansson, L. G., and Norell, M. (1999), Indication of Chromium Oxide Hydroxide Evaporation During Oxidation of 304L at 873 K in the Presence of 10% Water Vapor, Oxidation of Metals, 52(1–2), pp. 95–111. Asteman, H., Svensson, J. E., and Norell, M., (2000), Influence of Water Vapor and Flow Rate on the High-Temperature Oxidation of 304L; Effect of Chromium Oxide Hydroxide Evaporation, Oxidation of Metals, 54(1–2), 11–26. Bar-Cohen, Y., and Mal, A. K. (1989), Ultrasonic Inspection, in, ASM Handbook Volume 17: Nondestructive Evaluation and Quality Control, 9th ed., Metals Park, Ohio: ASM International, pp. 231–277. Blum, R., and Hald, J. (2002), Benefit of Advanced Steam Power Plants, in Materials of Advanced Power Engineering, Vol. 21, Part II, Lacomte-Becker J. (ed.) Ed., European Commission and University of Lie´ge, pp. 1007–1015. Blum, R., Vanstone, R. W., and Messelier-Gouze, C. (2003). Materials Development for Boilers and Steam Turbines Operating at 7008C, in Parsons 2003, Proceedings of the Sixth International Parsons Turbine Conference, pp. 489–510. Booras, G. S., Viswanathan, R., Weitzel, P., and Bennett, A. (2003), Economic Analysis of Ultra Supercritical PC Plants, in Pittsburgh Coal Conference Proceedings, 55.1 (Pittsburgh, September 2003), Pittsburgh: University of Pittsburgh. Cory, N. J., and Herrington, T. M. (1987), Kinetics of Oxidation of Ferrous Alloys by Super-heated Steam, Oxidation of Metals, 28(5–6), 237–258. Dooley, B. (2003), The Importance of Oxide Growth and Exfoliation, in EPRI International Conference on Materials and Corrosion Experience for Fossil Power Plants Proceedings (Isle of Palms, South Carolina, November 2003), Palo Alto, California: EPRI. Dooley, R. B., and McNaughton, W. P. (1995), Boiler Tube Failures: Theory and Practice, Vol. 1: Boiler Tube Fundamentals, EPRI TR-105261, Palo Alto, California: EPRI. Dooley, R. B., and Paterson, S. J. (2003), Oxide Growth and Exfoliation in Steam: Plant Experience, in: Proceedings of the EPRI-NPL Workshop on Scale Growth and Exfoliation in Steam Plant, Teddington, England: National Physical Laboratory. Fukuda, Y., Tamura, K., and Suzaki, K. (1995), Effect of Cr and Si Contents on the Steam Oxidation of High Cr Ferritic Steels, in: International Symposium on Plant Aging and Life Prediction of Corrodible Structures (Sapporo, Japan, May 15–18, 1995) JSCE and NACE, pp. 835–840. Fukuda, M. (2009), Advanced USC Technology Development in Japan, in: 3rd Symposium on Heat Resistant Steels and Alloys for High Efficiency USC Power Plants (Tsukuba, Japan, June 2–4, 2009), Tsukuba, Japan: National Institute for Materials Science. Gaskell, D. R. (1973), Introduction to Metallurgical Thermodynamics, New York: McGraw-Hill. Gindorf, C., Hilpert, K., and Singheiser, L. (2001), Determination of Chromium Vaporization Rates of Different Interconnect Alloys by Transpiration Experiments, in: Solid Oxide Fuel Cells (SOFC VII), Yokokawa, H. and Singhal, S. C. (eds.), PV 2001-16, Pennington, New Jersey Electrochemical Society, pp. 793–802. Grobner, P. J., Clark, C. C., Andreae, P. V., and Sylvester, W. R. (1980), Steamside
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Oxidation and Exfoliation of r-Mo Superheater and Reheater Steels, 80172 from CORROSION/80 (Chicago, Illinois, March 1980), Houston: National Association of Corrosion Engineers. Haynes International (2002), Pocket Guide to High Temperature Properties of AgeHardenable High Performance Alloys. H-3127B, Kokomo, Indiana: Haynes International. Haynes International (2008a), Haynes 230 Alloy. H-3135C, Kokomo, Indiana: Haynes International. Haynes International (2008b), Haynes 282 Alloy. H-3172A, Kokomo, Indiana: Haynes International. Henry, J. F., and Ward, C. T. (2006), Lessons from the Past: Materials-Related Issues in the Ultrasupercritical Boiler, Eddystone Unit 1, Energy Materials, 1(2), pp. 88–97. Holcomb, G. R., (2006), Corrosion in Supercritical Water – Ultrasupercritical Environments for Power Production. in: ASM Handbook Volume 13C: Corrosion: Environments and Industries, Cramer, S D., and Covino J, B. S, (eds.), Materials Park, Ohio: ASM International, pp. 236–245. Holcomb, G. R. (2008), Calculation of Reactive-Evaporation Rates of Chromia, Oxidation of Metals, 69(3–4), pp. 163–180. Holcomb, G. R. (2009a), Steam Oxidation and Chromia Evaporation in UltraSupercritical Steam Boilers and Turbines, Journal of the Electrochemical Society, 156(9), pp. C292–C297. Holcomb, G. R. (2009b), Oxidation in Environments with Elevated CO2 Levels, in: 23rd Annual Conference on Fossil Energy Materials (Pittsburgh, Pennsylvania, May 12–14, 2009), Pittsburgh, Pennsylvania: National Engineering Technology Laboratory. Jonas, O. (2000), Effective Cycle Chemistry Control, in: ESAA Power Station Chemistry Conference Proceedings (Rockhampton, Queensland, Australia, May 2000), Melbourne: Energy Supply Association of Australia. Kaur, I., Mishin, Y., and Gust, W. (1995), Fundamentals of Grain and Interphase Boundary Diffusion, 3rd ed., New York: John Wiley & Sons, Inc. Kritzer, P., Boukis, N., and Dinjus, E., (1999), Factors Controlling Corrosion in High Temperature Aqueous Solutions: A Contribution to the Dissociation and Solubility Data Influencing Corrosion Processes, Journal of Supercritical Fluids, 15(3), pp. 205–227. Kubaschewski, O., and Alcock, C. B. (1979), Metallurgical Thermochemistry, 5th ed., New York: Pergamon Press. Manning, M. I. and Meadowcroft, D. B. (1980), Effects of Tube Creep Strains on Laminated Scale Formation in Ferritic Pressure Tubing, Report No. RD/L/ R2012, London: CERL. Masuyama, F. (2001), History of Power Plants and Progress in Heat Resistant Steels, ISIJ International, 41(6), 612–625. Matsuo, H., Nishiyama, Y., and Yamadera, T. (2005), Steam Oxidation of FineGrain Steels, in: Advances in Materials Technology for Fossil Power Plants, Viswanathan, R., Gandy, D., and Coleman, K. (eds.), Materials Park, Ohio: ASM International, pp. 441–451. Montgomery, M., and Karlsson, A. (1995), Survey of Oxidation in Steamside Conditions, VGB Kraftswerkstechnik, 75(3), 235–240. Opila, E. J., Myers, D. L., Jacobson, N. S., Nielsen, I. M. B., Johnson, D. F., Olminsky, J. K., and Allendorf, M. D. (2007), Theoretical and Experimental
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485
Investigation of the Thermochemistry of CrO2(OH)2(g), Journal of Physical Chemistry A, 111(10), 1971–1980. Osgerby, S. J., and McCartney, L. N. (2002), Steam Oxidation of 9–12 Cr Martensitic Steels: Characterisation and Modelling the Spalling of Oxide Scale, in Materials for Advanced Power Engineering: Part III, Lecomte-Beckers, J., Carton, M., Schubert, F., and Ennis, P. J. (eds.), Ju¨lich, Germany: Forschungszentrum Ju¨lich GmbH, pp. 1613–1620. Otsuka, N., and Fujikawa, H. (1991), Scaling of Austenitic Stainless Steels and Nickel Base Alloys in Temperature Steam at 9738K, Corrosion, 47(4), pp. 240– 248. Paterson, S. R., Moser, R. S., and Rettig, T. W. (1992), Oxidation of Boiler Tubing, in Interaction of Iron-Based Materials with Water and Steam, Dooley, R. B., and Bursik, A. (eds.), Report No. TR-102101, Palo Alto, California: EPRI, pp. 8-1 to 8-25. Rehn, I. M. (1981), Corrosion Problems in Coal-Fired Boiler Superheater and Reheater Tubes: Steam-Side Oxidation and Exfoliation, Report No. CS-1811, Palo Alto, California EPRI. Reynolds, W. C. (1986), Implementation in the Interactive Program, STANJAN Version 3, Palo Alto, California: Department of Mechanical Engineering, Stanford University. Sarver, J., Viswanathan, R., and Mohamed, S. (2003), Boiler Materials for Ultra Supercritical Coal Power Plants–Task 3, Steamside Oxidation of Materials – A Review of Literature, Topical Report, U.S. Department of Energy Grant Number DE-FG26-01NT41174 and Ohio Coal Development Office Grant Agreement D-0020. Schu¨tze, M., Tortorelli, P. F., and Wright, I. G. (2010), Development of a Comprehensive Oxide Scale Failure Diagram, Oxidation of Metals, 73, 389–418. Seth, B. B. (1999), US Developments in Advanced Steam Turbine Materials, in Advanced Heat Resistance Steels for Power Generation, Palo Alto, California: EPRI, pp. 519–542. Smith, J. W. (1998), Supercritical (Once Through) Boiler Technology, BR-1658, Barberton: Babcock & Wilcox Company. Special Metals Corporation, (2004a), Inconel Alloy 740, SMC-090, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2004b), Nimonic Alloy 90, SMC-080, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2004c), Udimet Alloy 720, SMC-106, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2005), Inconel Alloy 617. SMC-029, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2006), Inconel Alloy 625. SMC-063, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2007a), Inconel Alloy 718. SMC-045, Huntington, West Virginia: Special Metals Corporation. Special Metals Corporation, (2007b), Nimonic Alloy 105, SMC-081, Huntington, West Virginia: Special Metals Corporation. Staubli, M., Mayer, K. H., Kern, T. U., Vanstone, R. W., Hanus, R., Stief, J., and Scho¨nfeld, K. H. (2001), COST 522 – Power Generation into the 21st on the Century; Advanced Steam Power Plant, in Advances in Materials Technology for
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Power plant life management and performance improvement
Fossil Power Plants, Viswanathan, R., Bakker, W. T., and Parker, J. D. (eds.), University of Wales and EPRI, pp. 15–32. Staubli, M., Scarlin, B., Mayer, K. H., Kern, T. U., Bendick, W., Morris, P., DeGianfrancisco, A., and Cerjak, H. (2003), Materials for Advanced Steam Power Plants: The European COST522 Action, in Proceedings of the 6th International Charles Parsons Turbine Conference (Dublin, Ireland, September 16–18, 2003,) pp. 305–324. Stultz, S. C. and Kitto, J. B. (1992), Steam, 40th ed., Barberton, Ohio: Babcock & Wilcox. Tedman, Jr, C. S. (1966), The Effect of Oxide Volatilization on the Oxidation Kinetics of Cr and Fe–Cr Alloys, Journal of the Electrochemical Society, 113(8), 766–768. Viswanathan, R., and Bakker, W. (2001a), Materials for Ultrasupercritical Coal Power Plants – Boiler Materials: Part 1, Journal of Materials Engineering and Performance, 10(1), 81–95. Viswanathan, R., and Bakker, W. (2001b), Materials for Ultrasupercritical Coal Power Plants–Boiler Materials: Part II, Journal of Materials Engineering and Performance, 10(1), 96–101. Viswanathan, R., Armor, A. F., and Booras, G. (2003), Supercritical Steam Power Plants – An Overview, in Best Practices and Future Technologies, New Delhi, India: National Thermal Power Corporation’s Center for Power Efficiency and Environmental Protection (CenPEEP) and the US Agency for International Development (USAID). Viswanathan, R., Henry, J. F., Tanzosh, J., Stanko, G., Shingledecker, J., Vitalis, B., and Purgert, R. (2005), U.S. Program on Materials Technology for Ultrasupercritical Coal Power Plants, Journal of Materials Engineering and Performance, 14(3), 281–292. Viswanathan, R., Coleman, K., Shingledecker, J., Sarver, J., Stanko, G., Borden, M., Mohn, W., Goodstine, S., and Perrin, I. (2007), Boiler Materials for Ultrasupercritical Coal Power Plants, Phase1: Final Summary Report, October 1, 2001–December 31, 2006, U.S. Department of Energy Grant Number DEFG26-01NT41175 and Ohio Coal Development Office Grant Agreement D-0020. Vitalis, B. P. (2006), Constant and Sliding-Pressure Options for New Supercritical Plants, Power, January/February, 2–7. Wright, I. G., Schu¨tze, M., Paterson, S. R., Tortorelli, P. F., and Dooley, B. (2004), Progress in Prediction and Control of Scale Exfoliation on Superheater and Reheater Alloys, in International Conference on Boiler Tube and HRSG Tube Failures and Inspections (San Diego, November (2004), Palo Alto, California: EPRI. Wright, I. G., Tortorelli, P. F., and Schu¨tze, M. (2007), Program on Technology Innovation: Oxide Growth and Exfoliation on Alloys Exposed to Steam, Report 1013666, Palo Alto, California: EPRI. Young, D. (2008), High Temperature Oxidation and Corrosion of Metals, San Francisco, California: Elsevier. Young, D., and Pint, B. (2006), Chromium Volatilization Rates from Cr2O3 Scales into Flowing Gases Containing Water Vapor, Oxidation of Metals, 66(3–4), pp. 137–153.
© Woodhead Publishing Limited, 2011
© Woodhead Publishing Limited, 2011
9.0
9.0
9.0 9.0
Bal
Bal
Bal Bal
0.25
0.8
Ni
Co
0.94 0.5
1.0
2.0
1.0
0.1
1.0
0.5
Mo
Si
Ti
0.2
0.3
0.12 0.07
0.10
0.07
0.12
0.2 0.06
0.4
0.3
0.6
9Cr ferritic steels
0.06
0.12
2Cr ferritic steels
0.5
114Cr ferritic steel 0.15
C
Al
40
30
B ppm
0.51 0.45
0.45
0.45
0.45
0.45
0.45
0.45
Mn
0.9 1.8
1.6
W
0.20 0.20
0.20
0.25
V
0.06 0.05
0.08
0.05
Nb
0.06 N 0.06 N
0.05 N
Other
2 For economic reasons it is not unusual for actual chemical compositions of materials delivered for use in power plants to be near the lower end of the specification. For some alloys (such as the 9Cr ferritic steels) the oxidation behavior significantly changes over the specification range.
9.0
Bal
2.25
Bal
T9 STBA26 HCM9M STBA27 T91 STBA28 E911 T92 STBA29 NF616
2.25
Bal
T22 STBA24 T23 STBA24J1 HCM2S
1.25
Cr
Bal
Fe
T11
Alloy designation
Source values that were ranges are listed as the midpoint of the range.2 Source values that were maximums are listed as half the maximum. Only Nb is listed for sources that gave a value or range for Nb + Ta. Source values for Pb, P, and S are omitted. Ferritic and austenitic steel compositions are from Viswanathan and Bakker (2001a). Nickel-base superalloy compositions are from Haynes International (2002, 2008a, 2008b) and Special Metals Corporation (2004a, 2004b, 2004c, 2005, 2006, 2007a, 2007b).
11.10 Appendix: nominal alloy composition for alloys of interest Steam oxidation in steam boiler and turbine environments 487
TP304H SUS304HTB Super304H SUS304J1HTB TP321H SUS321HTB TP316H SUS316HTB TP347H SUSTP347HTB
18.0
18.0
16.0
18.0
Bal
Bal
Bal
Bal
11.0 11.0
Bal Bal
18.0
12.0 12.0
Bal Bal
Bal
12.0
12.0
Bal
Bal
12.0
Bal
HT91 20CrMoV121 HT9 20CrMoWV121 HCM12 SUS410J2TB TB12 T122 HCM12A SUS410J3TB NF12 SAVE12
Cr
Fe
Alloy designation
Appendix (cont.)
© Woodhead Publishing Limited, 2011
10.0
12.0
10.0
9.0
8.0
0.1
0.5
0.5
Ni
2.5 3.0
Co
2.5
0.2
0.50 0.4
1.0
1.0
1.0
Mo
Si
Ti
0.2 0.3
0.05 0.1
0.3
0.4
0.08
0.08
0.08
0.10
0.08
0.6
0.6
0.6
0.2
0.6
0.5
Austenitic steels
0.08 0.10
0.08 0.11
0.10
0.20
0.4
12Cr ferritic steels 0.20
C
Al
40
30 30
B ppm
1.6
1.6
1.6
0.8
1.6
0.50 0.20
0.50 0.60
0.55
0.60
0.60
Mn
2.6 3.0
1.8 2.0
1.0
0.5
W
0.20 0.20
0.20 0.20
0.25
0.25
0.25
V
0.8
0.40
0.07 0.07
0.05 0.05
0.05
Nb
N Ta N Nd
3.0 Cu 0.10 N
0.05 0.07 0.04 0.04
0.05 N 0.06 N 1.0 Cu
0.03 N
Other
488 Power plant life management and performance improvement
© Woodhead Publishing Limited, 2011
1.5 0.35 0.75 1.5 2.5 Bal 0.7 1.5
0.5
Haynes 230 Haynes 263 Haynes 282 Inconel 617 Inconel 625 Inconel 718 Inconel 740 Nimonic 90
Nimonic 105
Udimet 720Li
23.0
Bal
16
14.85
22 20 19.5 22 21.5 19 25 19.5
21.0
25.0
Bal
Bal
25.0
TP310 SUS310TB TP310NbN HR3C SUS310J1TB 800H NCF800HTB HR6W
Bal
Cr
Fe
Alloy designation
Appendix (cont.)
Bal
Bal
Bal Bal Bal Bal Bal 52 Bal Bal
43.0
32.0
20.0
20.0
Ni
14.75
20
2.5 20 10 12.5 0.5 0.5 20 18
Co
3
5
2 6 8.5 9 9 3.05 0.5
Mo
0.4
0.5
0.4
0.6
Si
0.08
0.5
Ti
0.4
Al
0.015
0.085
0.1 0.03 0.06 0.1 0.05 0.04 0.03 0.065 0.5
0.4 0.2 0.075 0.5 0.25 0.175 0.5 0.5
5
1.2
1.2 2.1 0.3 0.2 0.9 1.8 2.5
2.5
4.7
0.3 0.3 1.5 1.15 0.2 0.5 0.9 1.5
Nickel-base superalloys
0.08
0.08
0.06
0.08
C
150
65
100
30
50 30
75
30
B ppm
0.5
0.5 0.25 0.175 0.3 0.5
0.5 0.3
1.2
1.2
1.2
1.6
Mn
1.25
14
6.0
W
V
3.65 5.125 2.0
0.18
0.45
Nb
0.1 Cu 0.075 Zr 0.075 Zr 0.1 Cu 0.0375 Zr
0.15 Cu
0.15 Cu 0.25 Cu
0.02 La
0.2 N
Other
Steam oxidation in steam boiler and turbine environments 489
12 Steam boiler component loading, monitoring and life assessment J . T A L E R a n d P . D U D A , Cracow University of Technology, Poland
Abstract: Thermal and strength analysis for thick-walled pressure components of steam boilers will be presented. The effects of different ways of conducting start-up and shut-down operations on loading of steam boiler components will be analysed. Total damage of a component owing to creep and thermal fatigue is calculated using the Palmgren– Miner rule of damage summation. An analysis of obligatory standards will be conducted and new methods of calculating and monitoring stress at components of complex geometry will be proposed. Identification of transient temperature and thermal stress distributions on the basis of temperature measurements using the proposed methods will be described. Key words: inverse method, diagnostic system, stress, steam boilers, monitoring.
12.1
Introduction
In order to raise the efficiency of electric energy production in steam power stations, the properties of the live steam are also increased, achieving pressures in supercritical power blocks of the order of 30 MPa and temperatures of 600 8C (French, 1993; Bendick et al., 2001; Ja¨ger and Theis, 2001; Baur et al., 2003; Abe et al., 2008). This has resulted in increases in the wall thickness of the pressure components, despite the introduction of new steel grades (Blum et al., 2001; Ennis and Quadakkers, 2001; Kern and Wieghardt, 2001; Zabelt et al., 2001; Rauch et al., 2004). High thermal stresses occur in these thick-walled components during power boiler operation (Kim et al., 2005), especially during boiler start-up and shutdown. The greatest loads are generated in pressure elements such as superheater headers, high-pressure pipelines and casings of steam gate valves. There is a danger of water accumulation in the lower part of the element while superheated steam flows through the upper part. This leads to 490 © Woodhead Publishing Limited, 2011
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considerable temperature differences over the element circumference and these can sometimes exceed 200 8C. This phenomenon is not observed in the computer systems for thermal stress monitoring used so far (EN 12952-3, 2001; TRD, 2002) because they measure only the temperature difference through the wall thickness. Measuring the temperature at the top and at the bottom of the boiler component is also unreliable because the value of the thermal stress not only depends on temperature measurements at these two points but also on temperature changes versus angle on the outer surface of the horizontal element. Very high thermal stress can occur during the injection of cold water into the thick-walled component, as in the case of the attemperator. In some cases, the thick-walled boiler components are subjected to a thermal shock, especially when steam condensation occurs on their inner surface or when fresh water at a lower temperature enters the installation. The components that are critical to the rate of temperature change in the whole system are boiler drums, superheater headers and attemperator headers. Because the thick-walled boiler components are subjected to irregular and fast temperature changes, it is necessary to monitor their operation (Laire and Eyckmans, 2001; Hofsto¨tter et al., 2003; Kahlert et al., 2003; Kim, 2005; Abe et al., 2008), especially during boiler start-up and shut-down.
12.2
Analysis of different ways of conducting start-up and shut-down operations and their influence on thermal and total stress loads in critical pressure components
Thermal and strength analysis was conducted for the T-pipe presented in Fig. 12.1. The T-pipe was installed in a steam pipeline of a BP1150 boiler with a steam capacity of 1150 t/h. The T-pipe is made of 13HMF steel. ANSYS software based on the finite element method (FEM) was used for the calculation of time–space temperature and stress distributions. Assuming symmetry, one-quarter of the T-pipe area was divided into eight-node finite brick elements, as shown in Fig. 12.2. This is the optimal mesh size in terms of accuracy and solution time. The effects that different ways of conducting start-up and shut-down operations have on the loading of power block elements will be illustrated using the example of heating and cooling a thick-walled T-pipe (Fig. 12.1). The temperature and pressure histories considered are presented in Figs 12.3 and 12.4. The heat transfer coefficient on the T-pipe inner surface, h, is 500 W/m2 K. In the first method for boiler start-up and shut-down operations, it was assumed that heating from a steam pressure p = 0 to pressure pnom =
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12.1
T-pipe made of 13HMF steel.
12.2
Division of T-pipe into finite elements.
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12.3 The first method of boiler start-up and shut-down operations; steam temperature and pressure histories during T-pipe heating and cooling operations: (a) heating, (b) cooling.
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12.4 The second method of power boiler start-up and shut-down operations, steam temperature and pressure histories during T-pipe heating and cooling operations: (a) heating, (b) cooling.
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18 MPa is conducted with saturated steam, p = pn(Tm), where pn is the saturation pressure at temperature Tm. After reaching pressure pnom, T-pipe heating continues at a constant rate of temperature change of vT = 7 K/min. When the steam temperature reaches Tm = 540 8C, the operation continues at constant medium temperature (Fig. 12.3a). In the first part of the cooling process, pressure remains constant and temperature decreases at a constant rate of temperature change of vT = 7 K/min (Fig. 12.3b). When the medium temperature reaches the saturation temperature at pressure pnom = 18 MPa, further cooling continues with saturated steam according to the equation p = pn(Tm). The second method for boiler start-up and shut-down operations involves keeping the pressure change constant during heating and cooling. The saturated steam temperature is calculated from the equation T = Tn(pm), where Tn is the saturation temperature at pressure pm. After reaching x = 1 a further temperature increase is conducted at a constant rate of temperature change of vT = 7 K/min until the temperature Tnom = 540 8C is achieved (Fig. 12.4a). In the first part of the cooling process, the temperature decreases to 540 8C at a constant rate of temperature change of vT = 7 K/min. After reaching x = 1, a further temperature decrease is calculated from the equation T = Tn(pm). The calculation results for the heating process are shown in Fig. 12.5 and the results for the cooling process are shown in Fig. 12.6. Both figures present equivalent stress histories at point P1, calculated by the Huber hypothesis. In the first method for the power boiler start-up operation, i.e. with a linear steam temperature increase, the thermal stresses presented in Fig 7.5a reach 210 MPa at time 3300 s. Pressure-induced stresses have the opposite sign to the thermal stresses and they reduce total stresses. The highest total stress occurs at time 1300 s and is 90 MPa. The second method for the power boiler start-up operation, i.e. with a linear steam pressure increase, causes a greater steam temperature increase in the first part of the heating process. This leads to a faster increase in thermal stresses, which rise to 188 MPa at time 700 s. The steam pressure at this time is only 5.7 MPa and lowers total stresses to 118 MPa. The maximum total stresses caused by pressure and temperature distribution are smaller in the first method of power boiler start-up operation. The second method for power boiler cooling gives a higher steam temperature decrease at the end of the process than occurs in the first method. For this reason, the first method for power boiler start-up and shutdown operations is safer than the second one.
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12.5 Equivalent thermal and total stress histories at point P1 during Tpipe heating operation: (a) first heating method, (b) second heating method.
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12.6 Equivalent thermal and total stress histories at point P1 during Tpipe cooling operation: (a) first cooling method, (b) second cooling method.
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12.3
Monitoring of remnant lifetime of pressure components
A linear summation rule developed by Palmgren and Miner will be used for determination of the remnant life of monitored pressure components.
12.3.1 Determination of usage of an element caused by lowcyclic fatigue Owing to rapidly heating/cooling down pressure elements of boilers and turbines at the locations where larger stresses arise, e.g. hole edges, plastic strains and, after longer performance times, fatigue crack begin to occur. Knowing the difference between maximal and minimal stress 2σa during one cycle, it is possible, using the fatigue curve in Fig. 12.7, to determine the number of cycles for which fatigue usage of an element occurs. The fatigue curves were obtained for ferritic steels in the German boiler TRD rules. Figure 12.7 shows the results of fatigue investigations carried out by the Institute of Nuclear Researches in Poland using a 15HM steel (low alloy) at a temperature of 550 8C. In order to determine the degree of fatigue usage of an element as a result of low-cyclic fatigue, it is necessary to count the number of stress cycles. The analysis has to take into consideration not only the main cycle, i.e. start-up–shut-down, but also small cycles caused by
12.7 Fatigue investigation chart for the 15HM steel at a temperature of 550 oC.
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instantaneous changes of medium temperature and pressure. The method of counting cycles that gives the most accurate results is the ‘rainflow method’. Computer programs have been developed to calculate the permissible numbers of cycles, Nd, to count the number of cycles using the rainflow method and to determine the degree of fatigue usage of an element due to thermal fatigue.
12.3.2 Determination of usage of an element caused by creep For the pressure components of a boiler operating at temperatures over 400 8C, creep effects should also be considered. Failure time, Td, refers to elements that have simple shapes, e.g. pipes, and may be expressed using the creep strength curve for a precise temperature. Figure 12.8 shows the determination of Td for the element made of 15HM steel and operating at a temperature of 550 8C. From the calculated value of the circumferential stress σφ = pdm/2g, the failure time Td is determined; dm and g denote the mean diameter and the thickness of the cylindrical component respectively. The fact that the operating element temperature may change should also be taken into consideration. The total damage of an element owing to creep and thermal fatigue is
12.8 Creep life prediction for the cylindrical element.
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given by the Palmgren–Miner equation X q n X n t þ D Nd j k¼1 Td k j¼1
½12:1
where D = creep and fatigue allowable usage n = number of cycles under the jth conditions Nd = number of allowable cycles of load changes under the jth conditions t = load time under the kth conditions Td = allowable time of static load under the kth conditions In equation [12.1], which is used by European (EN 12952-3, 2001; TRD, 2002) and American (Rao, 2006) boiler standards, the first term on the left side represents thermal fatigue damage and the second term represents creep damage.
12.3.3 Analysis of obligatory boiler standards and regulations Direct measurement of thermal stresses is very difficult to achieve since the largest absolute value for these stresses usually appears at the inner surface, which is in contact with water or steam under high pressure. The simplest and most frequently used method for restricting thermal stresses in cylindrical pressure elements is prescribed by German Code Standards (TRD, 2002) and European Norm (EN) 12952-3 (2001). It is based on the principle that one should not exceed the limit of heating and cooling rates of pressure elements. The allowable rate of temperature change during heating is determined from am
ðp p0 Þdm vT s 2 þ aT jw jf ¼ smin 2s a
½12:2
The maximum allowable cooling rate is calculated in a similar way: am
ðp p0 Þdm vT s 2 þ aT jw jf ¼ smax 2s a
½12:3
In equations [7.2] and [7.3], the following notation applies: l a = cr =thermal diffusivity (m2/s) dm = rin + rout= mean diameter (m) p = absolute pressure (MPa) p0 = ambient pressure (MPa) s = rout – rin = thickness of cylindrical element (m) vT = rate of temperature change of medium or pressure element wall (K/s)
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Steam boiler component loading, monitoring and life assessment αm αT σmin σmax
501
= pressure-induced stress concentration factor = thermal stress concentration factor = stress limit during start-up (heating) (MPa) = stress limit during shut-down (cooling) (MPa)
Coefficients φw and φf are defined by equations: jw ¼ and
Eb 1n
jf ¼
½12:4
3u2 1 u2 1 4u4 ln u 8ð u2 1Þ ð u 1Þ 2
½12:5
where E = Young’s modulus (MPa) β = linear coefficient of thermal expansion (1/K) ν = Poisson’s ratio u = rrout = quotient of inner surface radius and outer surface radius in Equations [12.2] and [12.3] can also be used to calculate the maximum total stress on hole edges. The heating or cooling rate of temperature change in pressure elements νT can be calculated using the following equation (Taler, 1999; Taler and Zima, 1999): df 1 ð300fi5 294fi4 532fi3 503fi2 vT ¼ ¼ dt t¼ti 514Dt 296fi1 þ 296fiþ1 þ 503fiþ2 þ 532fiþ3 þ 294fiþ4 300fiþ5 Þ ½12:6 where fi are medium or wall temperatures at 11 successive time points with Δt time step. Equations [12.2], [12.3] and [12.6] are not only valid for a quasi-stationary state, but also when the temperature change rate νT is a function of time: vT = vT(t) (Taler, 1999; Taler and Zima, 1999). The concentration coefficients of stresses caused by pressure and thermal load are assumed to be in compliance with this code. The main drawback to TRD codes is the assumption of a one-dimensional temperature field in components of complex geometry. For construction elements with complicated geometry, the stress value in the stress concentration areas can be calculated using the FEM. By determining the so-called influence function with the use of the FEM, one is subsequently able to make an on-line stress calculation with a known heat transfer coefficient and medium temperature transient. First, the influence function F(t) is calculated. This represents the history of a chosen stress
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component at a selected point in the construction element caused by a unit step change of a medium temperature. Next, the measured medium temperature Tm(t) is approximated by a staircase function. Next, the online stress distribution can be calculated. This method has many weak points. First, it is based on the superposition method and for that reason the problem must be linear. Therefore thermal and mechanical properties cannot be a function of temperature; they must be constant and strength analysis can only be linear. In addition, this method requires knowledge of the transient heat transfer coefficient and of moving fluid temperature. Both values can be estimated but usually the accuracy of the estimations is low. Therefore, new methods for stress calculation and stress monitoring at components of complex geometry will be developed.
12.3.4 Monitoring of temperature and stress distributions in pressure components The following section presents temperature and stress transients for selected pressure elements from a steam power boiler with a capacity of 6506103 kg/h, which produces steam with the following parameters: p =13.5 MPa, T = 540 8C. The re-heated steam parameters are p = 2.2 MPa and T = 540 8C. From among many monitored construction elements, four were chosen for this analysis: the boiler drum and three headers (Table 12.1). Limit stresses σmin and σmax are determined using a fatigue strength diagram (EN 12952-3, 2001; TRD, 2002) under the assumption of n = 2000 start-ups from a cold state. Table 12.2 presents the results obtained from the calculation of the allowable heating and cooling rates of the elements determined from equations [12.2] to [12.6]. For the study, 16 time periods of transient power boiler operation, which included start-up and shut-down operations, were recorded and analysed. Temperature and stress distributions in the horizontal elements, such as the boiler drum and outlet headers, were obtained on the basis of temperature readings taken at seven points, located on the outer insulated surface at a uniform distance from each other. The methods for determining transient temperature and stress distributions based on temperature measurements at the outer surface of the pressure component are presented in the following references: Duda (2003, 2005), Duda and Taler (2000), Duda et al. (2004), Taler (1999), Taler and Duda (2006), Taler and Zima (1999), Taler et al. (1999, 2002). It was assumed that a cylindrical element can lengthen and bend itself freely. Figure 12.9 shows the division of half of the cylindrical element’s cross-section into finite elements. Temperature distribution and stresses in the component were determined based on the solution of the inverse heat conduction problem (IHCP).
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0.1385 0.1385 0.285
0.1885 0.1885 0.315
3.51 3.76 3.90 3.34
8.526106 8.016106 7.966106
φw (MPa/K)
8.156106
a (m2/s)
1.356105
1.386105
1.336105
1.356105
β (1/K)
38
40
41
39
k (W/mK)
1.736105
1.986105
1.986105
1.826105
E (MPa)
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Boiler drum 18GMNA (15NcuMNb) Live steam outlet header 12H1MF Live steam attemperator (III) 12H1MF Steam re-heater outlet header 15H1MF
σmax(MPa) 511.310 312.295 315.758 295.373
σmin(MPa) 198.567 192.275 195.737 234.615
46.945
2.468 9.448 9.618
vT heating rate at beginning of start-up (K/min)
59.102
6.355 15.346 15.516
vT heating rate at the end of start-up (K/min)
46.945
2.468 9.448 9.618
vT cooling rate at beginning of shut-down (K/min)
7700
7680
7700
7840
r (kg/m3)
59.102
6.355 15.346 15.516
vT cooling rate at the end of shut-down (K/min)
620
650
625
610
c (J/kgK)
Table 12.2 Calculated allowable stresses and heating and cooling rates of pressure elements using TRD regulations
0.785
rin (m)
0.9
rout (m)
Monitored pressure elements of a power boiler with 6506103 kg/h capacity
Boiler drum (18GMNA) Ø18006115 Live steam outlet header (12H1MF) Ø377650 Live steam attemperator (III) of (12H1MF) Ø377650 Steam re-heater outlet header (15H1MF) Ø630630
Table 12.1
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12.9 Division of the half cross-section of the horizontal cylindrical element into finite elements; temperature is measured at points 1 to 7 (nodes 1, 4, 6, 8, 10, 12, 2).
The analysis of temperature and stress field results will be presented for the outlet header of the steam re-heater. The determined temperature and stress transients are shown in Figs 12.10 to 12.14. The measured temperature transients at points 1 and 7, i.e. at the lowest and highest points, and the pressure transient inside the header are presented in Fig. 12.10. One can observe that while the wall temperature at the lower part of the outlet header is about 150 8C, the upper part has a considerably higher temperature. This is most probably caused by the flow of non-evaporated water from the re-heater’s non-drainable pipe coils into the re-heater outlet header. Water accumulates in the lower part of the outlet header. The temperature transients at the seven points on the inner surface, located opposite the measurement points on the outer surface, are shown in Fig. 12.11. The temperature difference transient ΔT in the inner surface of the wall at the lowest and highest points ΔT = Tg – Td is also presented. Owing to a relatively high temperature difference recorded on the header’s circumference, axial thermal stresses are high (see Fig. 12.13).
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12.10 Measured pressure and temperature transients on the outer header’s re-heater surface.
12.11 Temperature transients on the outlet header’s inner surface (opposite points 1–7) and temperature difference transients between the outlet header’s top and bottom ΔT = TgTd, Tg and Td are the outlet header’s inner surface temperature at the highest and lowest points respectively.
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12.12 Circumferential thermal stress transients on the outlet header’s inner surface at the points located opposite points 1–7.
12.13 Axial thermal stress transients on the outlet header’s inner surface at the points located opposite points 1–7.
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12.14 Equivalent stress transients on outlet head’s inner surface at points located opposite points 1–7.
Therefore, longitudinal tension stresses occur in the header’s lower, colder part. One should, therefore, expect a relatively high thermal stress on the hole edges of downcomers, because of a stress concentration on the hole edges. Pressure-induced maximal equivalent stresses in the header are not, as absolute values, much larger than axial thermal stresses, since the pressure of the re-heated steam is relatively low.
12.3.5 Creep behaviour of pressure components and measurement of creep strains ANSYS software based on the FEM was used for the calculation of time– space strain and stress distributions in a T-pipe made of 13HMF steel. It was assumed that the T-pipe has a uniform temperature distribution – in the first case 540 8C and in the second 560 8C. The circumferential strain history caused by creep is presented in Fig. 12.15. It can be seen that the circumferential strain after 200 000 h at point P1, where stress concentration occurs, is high. The circumferential strain at point P5, where stress concentration does not occur, is low. This was confirmed by measurements on straight sections of pipelines that have been carried out over the last 30 years (Taler, 1988; Taler and Mlynarski, 1988). The measured circumferential strains on the pipeline outer surface did not exceed 0.5 % after
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12.15 Circumferential strains caused by creep at points P1 and P5; the T-pipe made from 13HMF steel loaded by pressure.
150 000 h of operation. Figure 12.15 shows that the influence of temperature on the circumferential strain is much higher at point P1. The creep phenomenon causes stress relaxation at points P1 and P5. Figure 12.16 shows that circumferential stress at 540 8C decreases almost by three times after 200 000 h of operation. Stress relaxation at point P5 is even higher. However, after 80 000 h of operation, a stress increase was observed at 560 8C. This was caused by a redistribution of stresses in the T-pipe because of the creep phenomenon. A similar phenomenon is observed in a straight pipe loaded with an inner pressure. During secondary creep, stresses on the outer surface are higher than those on the inner surface. An analysis was carried out on the influence of creep on total stresses caused by temperature distribution and pressure. It was assumed that steam temperature was suddenly lowered by 200 K; steam pressure equals the nominal value of pn = 18 MPa. The equivalent total stress history at point P1 is presented in Fig. 12.17. During the calculation of stress history, it was assumed that the T-pipe operated for 200 000 h in a temperature of 540 8C. It is evident that creepinduced relaxation of stresses in the T-pipe during the 200 000 h lowers the total stresses during a sudden reduction of temperature in the aforemen-
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12.16 Circumferential stresses caused by creep at points P1 and P5; the T-pipe made from 13HMF steel loaded by pressure.
12.17 Equivalent total stress history at point P1 caused by thermal shock and pressure in the new T-branch and one that has been in service for 200 000 h; the T-branch is made from 13HMF. © Woodhead Publishing Limited, 2011
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tioned element. T-pipe creep causes a decrease in equivalent stresses, which are triggered by pressure and thermal shock. The maximum equivalent stresses in the T-branch after 200 000 h in service are about 40 % smaller in comparison with the maximum stress in the new T-branch.
12.3.6 Measurements of creep strains The general arrangement of the pipelines and the location of measuring points are shown in Fig. 12.18. The boiler design parameters, such as steam pressure and temperature at the superheater outlet, are also presented. These 50-MW drum-type boilers are intended to run in steady conditions with as few start-ups and shut-downs as possible. This mode of operation does not cause high thermal stresses in the thick-walled pressure elements. For units that have been operating for over 20 years, remaining creeprupture life is the limiting life extension parameter. There are two main methods of predicting the remaining life of boiler pressure elements operating under creep conditions: 1.
2
Calculations based on design data and measured actual work conditions. This method has the advantage of requiring only temperature and pressure measurements. Stresses can be calculated from the measured values of pressure and temperature. Then, using the life prediction rules for elements operating under creep conditions in the cycling mode (e.g. the linear summation of damages rule presented by Miner–Palmgren), we can calculate residual life. The disadvantage of this method is that the data must be taken from tables. The actual material properties, creep strength, for example, can vary from published data. The method may be employed for making a rough estimate of accumulated damage or for determining the boiler components that should be subjected to more in-depth control. The method is not adequate for accurate determination of remaining life, particularly when the component in question is approaching the end of life and its residual life value is low. Methods based on measurements of damage. These methods involve measurements of parameters that indicate the damage of the pressure component. Measurements and inspections include: ○ diametral measurements to indicate deformation; ○ non-destructive inspection; ○ residual stress measurements; ○ structural examination (destructive methods and surface replicas).
This method of residual life assessment is more accurate than the previous method. Its disadvantage is that the measurements and inspections must be
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12.18 The general pipeline arrangement with the location of measuring points.
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done periodically. They must be done more frequently as the accumulated damage grows. The residual life of the pipeline can be evaluated by means of the external diameter deformation measurement. Diametral measurements to indicate creep swelling have been carried out for over 50 years. The first measurements were conducted in 1954. Four pins welded to the external surface of the pipeline made it possible to measure the pipeline deformation caused by creep (Fig. 12.19). The measurements are carried out on superheated steam pipelines operating at temperatures above 450 8C. The outside diameter of the pipeline is measured by means of a micrometer gauge. The measured diameter d´p should be corrected to obtain the value dp at the reference temperature 20 8C: 0
dp ¼
12.19
dp 1 þ ap Tp 20 am ðTm 20Þ
½12:7
Pins for diametral measurements to indicate creep swelling.
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Steam boiler component loading, monitoring and life assessment Table 12.3
513
Analysed pipelines data
Point number
d0=2r0 (m)
g (m)
p (MPa)
T (8C)
Steel
1 2 3 4 5
0.108 0.273 0.168 0.325 0.168
0.012 0.035 0.019 0.031 0.019
10.8 10.8 9.5 10 9.5
510 540 500 500 500
15HM 13HMF 12MH 15HM 12MH
where Tp = Tm = ap = αm =
pipeline temperature (8C) micrometer temperature (8C) thermal expansion coefficient of pipeline material (1/K) thermal expansion coefficient of micrometer material (1/K)
The operation parameters of the pipelines in question are given in Table 12.3. The creep curve of the pipeline, εφ,0(τ) for the first- and second-stage creep can be approximated by the function h ti ej;0 ðtÞ ¼ a 1 exp ½12:8 þ e_ m t b where a, b and e_ m are constants calculated from the equation N X
2 ei ej;0 ðti Þ ¼ min
½12:9
i¼1
where εi is the measured strain at time τi, i=1, . . . , N. The creep curves obtained from equation [12.8] are shown in Figs 12.20 to 12.22.
12.20 Creep curves of the pipelines in question and strain measurement results at points 1, 2 and 3.
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12.21 Creep curves of the pipelines in question and strain measurement results at points 4 and 5.
12.22
12.4
Schematic creep curve.
Conclusions
This chapter discusses the problems encountered in thermal stress monitoring and in determining the remaining life of pressure components in thermal power plants. The effects of two different methods of conducting start-up and shut-down operations on the loading of power block elements are illustrated using the example of heating and cooling a thick-walled Tpipe. In the first proposed method for the power boiler start-up operation, i.e. with a linear steam temperature increase, the highest total stress of
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90 MPa occured at time 1300 s. The second method for the power boiler start-up operation, i.e. with a linear steam pressure increase, caused a higher steam temperature increase in the first part of the heating process. This led to a faster increase of total stresses to 118 MPa. The second method for power boiler cooling gave a greater steam temperature decrease at the end of the process than the first method. For that reason, the first method for power boiler start-up and shut-down operations was safer than the second one. Monitoring of the thermal and pressure operating parameters of a boiler makes it possible to identify dangerous loads on the power boiler’s pressure elements during boiler start-up and shut-down operations. The start-up and shut-down operations can be run without exceeding the allowable stress values in the stress concentration areas. All stress transients caused by pressure and thermal loads are the basis for determining the creep and fatigue usage factors of pressure elements. Next, temperature distribution and stresses in the outlet header of the steam re-heater were determined based on the solution of the IHCP. The example presented in this chapter demonstrates high efficiency and flexibility of the monitoring method. Creep behaviour of pressure components was presented for a T-pipe made of 13HMF steel. It was assumed that the T-pipe had a uniform temperature distribution – in the first case 540 8C and in the second 560 8C. It can be seen that circumferential strain after 200 000 h is high at the point P1 where stress concentration occurs. Conversely, circumferential strain is low at the point P5 where stress concentration does not occur. This was confirmed by measurements on straight sections of pipelines that have been carried out over the last 30 years. Furthermore, the chapter also considered the creep deformations of both new pressure components and pressure components that had been in service for a specified time. T-branch creep causes a decrease in the equivalent stress that results from pressure and thermal shock loading. The maximum equivalent stresses in the T-branch after 200 000 h in service was about 40 % smaller in comparison with the maximum stress in the new T-branch. Next, it was shown that pipeline material creep curves can be established based on pipeline strain measurements. The creep curves allow an assessment of the residua1 life of an operating pipeline. Approximation of the creep curve with the functional relationship between tangential strain and time allows the calcu1ation of the time (1 %) at which a limiting strain of 1 % is accumulated in the pipeline. Calculation of the 2 % time is inappropriate because, if the strain is 2 %, the effective strain at the skeletal point is even higher, and the pipeline may then be in the third-stage creep regime.
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12.5
References
Abe F, Kern T-U and Viswanathan R (2008), Creep-resistant steels, Cambridge, UK, Woodhead Publishing Limited. Baur F, Stamatelopoulos G N, Vortmeyer N and Bugge J (2003), ‘Driving coal-fired power plants to over 50% efficiency’, VGB PowerTech, 12, 97–100. Bendick W, Hahn B and Schendler W (2001), ‘Development of creep damage in steel grades X10CrMoVNb9-1 (P/T 91) and X20CrMoV12-1’, VGB PowerTech, 12, 98–100. Blum R, Chen Q, Coussement C, Gabrel J, Testani C and Verelst L (2001), ‘Operational tests of superheater materials at high steam temperatures in a hard coal-fired steam generator’, VGB PowerTech, 10, 86–91. Duda P (2003), ‘Solution of multidimensional inverse heat conduction problem’, Heat Mass Transfer, 40, 115–122. Duda P (2005), ‘Identification of transient boundary conditions in complex-shape bodies’, Arch. Thermodynamics, 26, 17–33. Duda P and Taler J (2000), ‘Numerical method for the solution of non-linear twodimensional inverse heat conduction problem using unstructured meshes’, Int. J. Numerical Methods Engng, 48, 881–899. Duda P, Taler J and Roos E (2004), ‘Inverse method for temperature and stress monitoring in complex-shape-bodies’, Nuclear Engng Design, 227, 331–347. EN 12952-3 (2001), Water-tube boilers and auxiliary installations–Part 3: Design and calculation for pressure parts, Brussels, European Committee for Standardization. Ennis P J and Quadakkers W J (2001), ‘High chromium martensitic steelsmicrostructure, properties and potential for further development’, VGB PowerTech, 8, 87–90. French D N (1993), Metallurgical failures in fossil fired boilers, New York, John Wiley & Sons, Inc. Hofsto¨tter P, Keller H P, Hoppe Th, Protogerakis E and Werden B (2003), ‘Application of the potential drop method for monitoring an outside crack in the housing of quick-acting valve-safe continuation of operation for two years until removal of the housing’, VGB PowerTech, 6, 82–87. Ja¨ger G and Theis K A (2001), ‘Increase power plant efficiency’, VGB PowerTech, 11, 21–25. Kahlert J, Grabig J, Drews D and Beesdo H (2003), ‘Life-cycle monitoring and condition evaluation as a module of a modern database-supported test management system for power plants’, VGB PowerTech, 4, 58–61. Kern T U and Wieghardt K (2001), ‘The application of high-temperature 10Cr materials in steam power plants’, VGB PowerTech, 5, 125–131. Kim H J (2005), ‘Assessment of creep life fraction for in-service high-temperature components’, Engng Failure Anal., 12, 578–585. Kim Y W, Kim D O, Lee J S, Choi S and Zee S Q (2005), ‘Thermo-mechanical simulation for nozzle header of once-through steam generator by experiment and finite element method’, Int. J. Pressure Vessels Piping, 82, 602–609. Laire Ch and Eyckmans M (2001), ‘Evaluating the condition and remaining life of older power plants’, VGB PowerTech, 10, 98–102.
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Rao K R (2006), Companion guide to the ASME Boiler and Pressure Vessel Code, New York, ASME. Rauch M, Maile K, Seliger P and Reuter A (2004), ‘Development of creep damage at components made of X10CrMoVNb9-1(P91)’, VGB PowerTech, 9, 71–75. Taler J (1988), ‘Remaining life assessment of pressure boiler components under creep conditions based on strain measurements’, Energetyka, 42, No. 12, 435–438 (in Polish). Taler J (1999), ‘A new space marching method for solving inverse heat conduction problems’, Forschung im Ingenieurwesen (Engineering Research), 64, 296–306. Taler J and Duda P (2006), Solving direct and inverse heat conduction problems, Berlin, Heidelberg, Springer-Verlag. Taler J and Mlynarski F (1988), ‘Residual life assessment for boiler pressure components based on measurements of creep strains’, Trans. ASME, J. Pressure Vessel Technol., 110, 308–313. Taler J and Zima W (1999), ‘Solution of inverse heat conduction problems using control volume approach’, Int. J. Heat and Mass Transfer, 42, 1123–1140. Taler J, Węglowski B, Zima W, Grądziel S and Zborowski M (1999), `Analysis of thermal stresses in a boiler drum during start-up', Trans. ASME, J. Pressure Vessel Technol., 121, 84–93. Taler J, Węglowski B, Grądziel S, Duda P and Zima W (2002), `Monitoring of thermal stresses in pressure components of large steam boilers', VGB PowerTech, 1, 73–77. TRD (2002), Technische Regelen fu¨r Dampfkessel, Berlin, Carl Heymans Verlag– Beuth Verlag GmbH. Zabelt K, Melzer B, Reuter A and Seliger P (2001), ‘Results of recent investigations for boiler application on austenitic steels to ensure long-term service integrity at high steam temperatures’, VGB PowerTech, 2, 81–85.
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13 Steam turbine materials selection, life management and performance improvement R . W . V A N S T O N E a n d S . O S G E R B Y , Alstom Power, UK
Abstract: Materials technology is key to enabling improvements in steam turbine performance arising from retrofit. Materials selection, the main degradation mechanisms and material issues, together with any special processes required, are discussed for each major component of the high temperature and low temperature cylinders. It is demonstrated that the selection of appropriate materials and special processes allows the efficiency improvements arising from retrofit of steam turbines to be achieved with low risk of component failure. Key words: steam turbine, retrofit, materials.
13.1
Introduction
This chapter describes the issues considered in material selection and life management for retrofit steam turbines. Innovations yielding significant technical and economic benefit are outlined. The approach developed by steam turbine manufacturers to avoid failure by creep, fatigue, stress corrosion cracking, corrosion and erosion are described. There are two main types of steam turbine design – impulse and reaction. However, the differences between the two do not impact on materials issues and details of the respective designs are mainly outside the scope of this chapter. The one relevant aspect is that in impulse designs there is a greater pressure drop across the stationary blades, which leads to the presence of welded diaphragms.
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Design criteria for high temperature power plant components
Creep crack initiation High Creep Creep and cycle Component strain rupture growth fatigue Rotors Casings Blades Bolts
13.2
519
X X X
X X X X
X X
Low cycle fatigue and cyclic hold Fracture Creep endurance toughness relaxation
X X
X X X
X X X
High temperature cylinders
13.2.1 General concerns Materials for application in high and intermediate pressure cylinders of fossil fuel fired plant are required to operate in the creep regime. Integrity is, of course, also required at lower temperatures to deal with transient start-up and shut-down conditions. The permissible stresses in high temperature components are generally determined by either low strain creep strength, as in the case of diaphragms, where permissible stresses are limited by creep deflection, or creep rupture strength, as in the case of valve chests and inner casings. These permissible stresses are calculated to ensure design lives of 200 000 hours. It is therefore essential that the long term creep and rupture strengths of the materials applied are well known. Potential failure mechanisms must also be considered. In components that might contain natural defects smaller than non-destructive acceptance or detection criteria, a knowledge of crack growth rates by creep or fatigue is necessary. In components that are significantly stressed at lower temperatures, perhaps during start-up and shut-down cycles, fracture toughness must be determined. Low cycle fatigue and cyclic-hold strength must also be established to ensure integrity against the thermomechanical strain cycles imposed during start-up and shut-down. A summary of the properties to be established in major components is given in Table 13.1. An additional potential concern is that of solid particle erosion of blades. Oxide particles originating from exfoliation in the boiler have the potential to cause erosion on impact with the turbine blades. The phenomenon requires adoption of certain blade design features to minimise the risk of erosion occurring in steam turbines.
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13.2.2 Material selection Many of the alloys used in retrofit steam turbines have been in use since the 1960s. Material suppliers have many years of experience in manufacturing components in these materials, their long term properties have been established through multicast long term testing programmes, often carried out on a national or international collaborative basis, and service experience has been very successful. Therefore any risks associated with sourcing and with design and operation using these materials are extremely low. The alloys applied to major components are described in Table 13.2. Rotor materials Similar 1%CrMoV rotor alloys are applied by all turbine manufacturers in Europe and in the USA. However, for some manufacturers the alloy has been optimised to give maximum creep strength. Studies carried out in the 1960s established a link between the levels of elements such as Cr, Mo, V and C, microstructure and creep strength [1]. By careful control of the level Table 13.2
High temperature alloys and their ASTM equivalents
Component
Material
Nearest ASTM equivalent
Rotors
10%CrMoVNbN 1%CrMoV Nimonic 80A 9%CrMoVNbB 17%Cr13%NiW 11%CrMoVNb 12%CrMoV Cast 1%CrMoV Cast 214 %CrMo Cast 1%CrMo Cast 12 %Mo 10%CrMoVNbN 17%Cr13%NiW 9%CrMoVNbN 12%CrMoV 12%CrMo 12%CrMoAl 10%CrMoVNbN 9%CrMoVNbN 1 2%CrMoV Nimonic 80A 11%CrMoVNbN 12%CrMoV 1%CrMoVTiB 1%CrMoV
None A470 GR8 None None None A565 GR616 A565 GR616 A356 GR9 A356 GR10 A356 GR6 A217 GR WC1 None None A182 F91 A565 GR616 A473 GR403 A473 GR405 None A182 F91 None None A565 GR616 A565 GR616 None None
Moving blades
Inner casing Outer casing
Stationary blades
Diaphragm rings (impulse technology only) Bolts and shrink rings
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and ratios of these alloying elements the optimum upper bainitic microstructure can be achieved throughout the section of even the largest rotor forgings. Achieving this microstructure leads to higher creep strength and the possibility of applying low alloy materials at temperatures up to 565 oC without design compromise. Nonetheless, there are some applications, especially in reaction designs, where the higher creep and low cycle fatigue strengths of 10%Cr rotors are required. The 10%CrMo(W)VNbN alloy was developed in the collaborative European COST 501 project in the late 1980s and early 1990s [2]. Following long term characterisation of a prototype forging, the first commercial forgings entered service in the mid 1990s. Further long term multicast testing of the first commercial forgings was carried out in a project supported by VGB [3] so the properties of this material are now well established and the material is supported by a body of successful service experience. Depending on the specific turbine size and operating requirements, monobloc or welded rotors may be used. Welded rotors offer the possibility of employing different steels at different locations in the rotor to match the required mechanical properties at those locations. Welded rotors have been in use for over 70 years [4]. Casing materials 1%CrMoV is selected for castings for inner casings, with the slightly less creep-resistant but easier to weld and lower cost Cr–Mo alloys applied where new outer casings are necessary. All of these alloys have been used since the 1960s and 1970s. All these materials have sufficient temperature capability for any foreseeable retrofits so no new high temperature casing materials are currently required. Moving blades Martensitic alloys are used for most stages of blading. 11%CrMoVNbN or 12%CrMoV alloys are applied for most stages but a boron containing steel described in section 13.2.3 has been introduced for the most highly stressed high temperature stages. Where the application is particularly demanding, a stronger Ni-base alloy or austenitic steel is available. The use of coatings to reduce oxidation and solid particle erosion has been considered but has not yet been used in practice.
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Stationary blades In reaction machines stationary blades are essentially similar to moving blades so alloy selection is identical. In impulse machines, however, the stationary blades are incorporated in welded diaphragms so that issues of weldability become important. All of the alloys applied to diaphragms have proven weldability and have been used in steam turbine applications for many years. Bolting All the bolting materials have been in use since the 1960s and are supported by extensive long term material testing. In most cases low alloy or martensitic 9–12%Cr steels are adequate but in some highly stressed applications, Nimonic 80A is used. Shrink rings may be used in HP inner casings. This technology has been successfully applied for many years and shrink ring materials are selected to match the inner casing material.
13.2.3 Innovations in high temperature materials technology The primary objective of alloy selection is reliability. When significant advantages are offered by new materials or processes, these opportunities are taken only after full consideration of all potential risks so that reliability is not compromised. Modifications to alloy selection have come about through the availability of new advanced alloys developed with the primary objective of enabling increases in steam temperature. In Europe the focus for much of this development has been the collaborative COST 501 programme but there have also been US collaborations with EPRI (project 1403) as well as detailed private investigations. These developments have been successfully applied in new high temperature plant operating in Denmark with main and reheat temperatures of 580 oC and in Germany operating with reheat steam of 585 oC. However, the alloys developed for these applications also permit higher stresses at lower temperatures typical of retrofit applications (see Table 13.3). Their use permits greater design freedom and allows greater cylinder efficiencies to be achieved, for example by allowing an additional stage of blading. Modified 9%CrMo and 10%CrMoVNbN diaphragms The introduction of new alloys for diaphragms requires the establishment of long term creep properties coupled with demonstration of successful manufacture. Diaphragms are welded under conditions of high restraint
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Increased creep strength of advanced materials Relative 100 000 hour rupture strength at 538 oC Low alloy
Rotors
Conventional 12%Cr steel
100 %
Blades 100 %
Austenitic
150 % 100 %
Diaphragm rings
Advanced 9–12%Cr steels
135 %
150 %
180–220 %
in geometries with inherent notches. The demonstration of weldability is therefore essential and can best be achieved through the manufacture of fullscale prototypes. This approach has been applied successfully to two advanced high temperature alloys to enable greater design freedom and higher cylinder efficiency. The modified 9%CrMo alloy (9%CrMoVNbN) was developed at the Oak Ridge National Laboratory in the USA in the 1970s and is widely applied to pipework as P91. Its adaptation to other product forms, including castings and forgings, was carried out in the 1980s, in part through the EPRI 1403 programme. Products for diaphragm applications were developed in the early 1990s and were first applied in the Danish 580 oC plants [5]. More recently, the even stronger 10%CrMoVNbN alloy first applied to rotors has been applied to bar and forgings for diaphragms. Early in 2003, a prototype diaphragm in this alloy was manufactured (Fig. 13.1). This will give designers even greater freedom to maximise cylinder efficiency in the future.
13.1 Prototype diaphragm in 10%CrMoVNbN alloy.
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Table 13.4
Design criteria for low temperature power plant components
Low cycle fatigue High and cyclic Stress Flow Water Tensile cycle hold Fracture corrosion Corrosion accelerated droplet Component strength fatigue endurance toughness cracking fatigue corrosion erosion Rotors Casings Blades Bolts
X X X X
X X
X X X
X X
X
X
X X
X
X X
9%CrMoVNbB moving blades The COST 501 programme focused principally on developing improved alloys for rotor forgings and castings. However, the alloys developed for rotor forgings are also suitable for application as bar products. The strongest alloy developed in COST 501 was a boron containing alloy designated ‘Steel B’ [2]. Its boron content poses significant challenges to forgemasters in control of segregation and inspectability when manufacturing large forgings, but it is much more easily applied to bar products suitable for moving blades. This alloy was developed for application at temperatures of 600 oC and above but its high creep strength also gives it advantages at lower temperatures. It was first applied to a retrofit steam turbine in 1998.
13.3
Low temperature cylinders
13.3.1 General concerns Permissible stresses in LP cylinders are controlled by the need to avoid both instantaneous and time-dependent failure mechanisms. A summary of the possible failure modes is given in Table 13.4. The main concerns for low temperature cylinders include: . . .
avoidance of stress corrosion cracking of rotors/discs at all locations; avoidance of flow accelerated corrosion/erosion–corrosion; minimisation of water droplet erosion of last stage blades.
Avoidance of stress corrosion cracking (SCC) Serious attention was first drawn to the problem of stress corrosion cracking in nuclear steam turbines as a result of the in-service disintegration of a shaft-keyed, shrunk-on disc of a non-reheat turbine in the UK [6, 7]. It occurred at the Hinkley Point Magnox station in 1969 on a 60 MW nonreheat unit, and resulted in a major and extensive investigation into the
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extent of the industry exposure to SCC and the reason for the failure [9]. The designs of several UK OEMs existing at the time, and in operation all over the world, were affected. Both round and square keyways were found to be susceptible to SCC, together with a wide range of materials. However, there was one material of the many in use in the different shrunk-on disc designs that was found to be immune to SCC [8]. The reason was by no means clear at that time, but after over 30 years’ long term stress corrosion testing the underpinning science is now available. Prior to the Hinkley Point disc failure, the design of shrunk-on discs had already moved from the shaft key to the hub key and so one potential solution to the shaft-keyed disc problem had already been developed. Some OEMs, particularly in the USA, continued with the shaft-key designs and subsequently experienced problems similar to those in the UK in the mid to late 1970s. The problems in the USA also extended to the root fastenings of both reaction and impulse turbines. Alstom hub-keyed turbine designs began, in the mid 1980s, to exhibit SCC in the hub keyway itself, in the smooth non-keyed bore regions and later in the disc heads in a variety of root fastening designs. As a result of the shrunk-on disc cracking, most OEMs who had previously employed this method of construction abandoned it in favour of integral monobloc rotors for new designs. Nevertheless, shrunk-on discs are still in service in many nuclear units and continually present problems to the host utility. Only one OEM continues to promote the shrunk-on disc in new designs for nuclear LP turbines. During all of this time, another method of rotor construction, the welded rotor, was used exclusively in two companies (namely BBC and CEM) and during the period of the 1970s, 1980s and 1990s, when shrunk-on disc SCC was prevalent in nuclear turbines, no welded rotor ever experienced SCC. Following the first evidence of SCC in shrunk-on discs, steam turbine manufacturers started laboratory testing to evaluate the stress corrosion initiation and growth properties of rotor steels in steam environments. It became very clear that while much of the early testing was in the dosed environment, there was no doubt that SCC initiation can and does occur in pure condensing steam [9], and in the early 1970s a switch to long term testing was made. The rationale is that the short term approach to accelerate tests using non-representative environments has the potential for nonrepresentative cracking mechanisms. Such testing programmes have extended exposures to over 100 000 hours and have led to the conclusion that the main influences are material yield strength, operating stress, temperature and operating environment. For example, Fig. 13.2 shows the clear influence of yield strength on time to initiation of SCC in a pure steam environment [10]. For a given test temperature, it is seen that
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13.2 The effect of yield strength on initiation of SCC in a 312% NiCrMoV steel at high applied stress.
. .
As the material strength is reduced, so the stress required for SCC initiation is increased and, conversely, there is a marked reduction in the stress required to cause SCC initiation as the strength level is increased. For a given material strength, the required stress to initiate SCC is reduced as the temperature is increased.
Therefore, for any combination of material strength and temperature, it is possible to define the stress below which the material is invulnerable to SCC initiation. This concept leads to a family of threshold curves that interrelate material strength, component stress and stage operating temperature, illustrated in Fig. 13.3. The general principles of the threshold curves are: . . . .
The material is vulnerable to SCC when the co-ordinates of effective stress and proof stress lie to the right of the relevant iso-temperature threshold stress line. Temperature T1
The threshold stress represents a lower bound to stress corrosion vulnerability. Where data are available, operating experience shows that
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13.3 Illustration of threshold stress approach for rotor steels.
there is no case of SCC cracking in a normal steam environment that has occurred at an applied stress level below the threshold stress. This position is illustrated in Fig. 13.4. The range of operating temperatures where in-service SCC has been experienced is shown in the figure. However, in order to condense the available information, all stresses have been normalised to equivalent stresses at 95 8C. Above the threshold stress, the rotor is considered to be vulnerable to SCC, although as well as the examples of cracking, there are many examples of service experience where such vulnerability has not led to SCC. It is very clear from laboratory tests and extensive service experience that the lower is the material strength, the higher is the stress that can be applied without concern for SCC. This provides an explanation why shrunk-on discs have been vulnerable to SCC; typically the discs have been made from high yield strength material, with higher stresses than in monobloc and welded rotors. It can now be explained with confidence why one material used for shrunk-on discs in the late 1960s never cracked in stress corrosion; although stressed to the same high level as each of the other materials, it was of the lowest yield strength, in the range 496–627 MPa. Similarly, it explains why the welded LP rotor design has shown unsurpassed resistance to SCC. It is a low stress design and thus permits application of materials with low yield strength (typically 634–731 MPa). These strength levels are lower than those
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13.4 All operating experience where SCC has occurred in a normal plant steam environment is explained by the threshold stress approach.
used for typical monobloc rotors and are significantly lower than most shrunk-on discs. The general absence of SCC in nuclear HP rotors can also be explained by the typically low yield strength material from which such rotors are made, but also by the usually lower stresses, which are a function of particular designs. Where cracking is now occurring in HP rotors, it will be no surprise to find that stresses are higher than the threshold as defined by the threshold stress approach. An important advantage of the way in which welded rotors are constructed from individual forgings is that the strength of any forging can be adjusted by heat treatment at the forgemaster in order to optimise resistance to SCC. At a given stress level, a greater margin of assurance that SCC will not occur arises when a lower strength material is used. This
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13.5 Flexibility of the welded rotor to optimise strength of individual forgings prior to welding.
approach, which is very difficult to manage with monobloc rotors, and probably ill-advised for shrunk-on disc rotors, is illustrated in Fig. 13.5. To a large extent, this finding explains the exemplary and unsurpassed record of the welded rotor design in its resistance to SCC. More than 200 nuclear rotors of the welded type (all with more than 100 000 hours service and some with more than 200 000 hours service) are in operation without any SCC cracks ever having been detected. Cracks have recently been detected in some root fastenings of welded rotors at one BWR station in Sweden during a planned maintenance outage, after more than 130 000 hours of service. However, finite element analysis of the fastenings indicates that the peak stresses in the affected roots are well above the threshold stress for SCC crack initiation, confirming their vulnerability to SCC. These rotors have been retrofitted after redesigning the root fixings to ensure that the peak stresses as calculated by three-dimensional FE are lower than the threshold stress. Avoidance of flow accelerated corrosion/erosion–corrosion The term flow accelerated corrosion (FAC) describes an increase in the rate of corrosion or material dissolution caused by the relative movement between a corrosive fluid and a material surface. It can occur in areas of power plant where carbon steels, or low alloy steels with less than about 1 % chromium content, are used [11]. In nuclear plant, piping, preheaters, valves and turbine casings have been frequently affected. The protection of the steam path against erosion is achieved by giving appropriate attention to the following factors: .
Selection of materials. If the complete steam path in the retrofitted HP
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13.6
. .
Erosion protection ring inlays between stationary stages.
turbine is made from 11–12%Cr steels, FAC is not a problem. Although the steam path in the LP turbine is also of high chromium steels, there is some exposure of the fixed blade carriers, which are generally made from nodular cast iron. In areas where previous experience indicates FAC is a significant concern, protection may be provided by inlaid rings of 12%Cr steel (Fig. 13.6). Another approach adopted in some circumstances is the selection of chromium-alloyed steel castings for the blade carriers. Arrangements for moisture extraction. To minimise FAC, proven methods of water removal may be applied to retrofitted turbines, such as the machining of extraction slots between stages. Application of protection where needed (e.g. stainless steel cladding, metal spray coating). Where previous problems have occurred procedures for repair using stainless steel welding, flame spray, plasma spray and HVOF coatings that have been well proven in service in nuclear units have been developed. These processes have been applied, for example, to cross under pipes, moisture separators and casings since the 1970s. Occasionally, the procedures are used on new components where there is a concern about potential FAC.
Minimisation of water droplet erosion of last stage blades Water droplet erosion is a concern with last stage LP blades, especially in nuclear turbines where the expansion line leads to high wetness. In nuclear turbines the wetness in the steam at the exhaust of the high pressure turbine is removed and the steam is reheated in the moisture separator reheater
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13.7 Hardened zone at the leading edges of moving blades.
(MSR). Nevertheless, the level of moisture in the last stages of the LP turbines is considerably higher than in fossil reheat machines. Measures are therefore introduced in nuclear turbines to reduce the risk of water droplet erosion damage, including the provision of moisture removal before the last stage and optimising the spacing between the last stage fixed and moving blades. In addition, treatment of the blade itself may be carried out if evaluation deems it to be necessary. Two methods are available. The first choice is flame or induction hardening of the moving blade leading edge (Fig. 13.7). The hardening process is automated and fully adapted to the selected blade style. After hardening, the blades are tempered in order to maximise toughness and stress corrosion resistance, while retaining a high hardness in the critical location to provide the erosion protection. In the majority of cases, induction hardening together with the moisture removal measures is sufficient to provide the required erosion protection. In other cases, where additional erosion protection is foreseen, brazed-on Stellite shields are used. In some cases, particularly in BWR applications, cobalt based alloys are prohibited in the steam path and other shield materials are then considered.
13.3.2 Material selection In parallel with the HP cylinder, many of the alloys used in retrofit LP cylinders have been in use for many years. Although collaborative research is less common for LP materials than those used at higher temperatures, OEMs have invested significant resources in ensuring that the long term performance of these materials is appropriate for the application. OEMs
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Table 13.5
LP cylinder alloys and their ASTM equivalents
Component
Material
Nearest ASTM equivalent
Rotors Moving blades
NiCrMoV 12CrNiMo PH-15Cr5Ni Nodular cast iron 12CrNiMo Nodular cast iron 12CrNiMo 12CrNi 12Cr
None None A705 XM-12 A536 60-40-18 None A536 60-40-18 None A240 410 A276 403
1CrMoV
None
Inner casing Stationary blades
Diaphragm rings (impulse technology only) Bolts and shrink rings
also work closely with their suppliers to ensure consistent quality; therefore the risks associated with these materials are minimised. The alloys applied to major components are described in Table 13.5. Rotor materials Most OEMs use similar NiCrMoV steels for LP rotors. The choice of a specific material depends on the rotor design – welded, monobloc or shrunkon disk. In any case the strength of the steel is controlled in a range to meet minimum strength requirements dictated by the need to support the moving blades and maximum allowable stress to avoid SCC. Surface treatment, typically shot peening or cold rolling, of specific areas of the rotor is frequently used as further mitigation against SCC. Casing materials Nodular cast iron is frequently used and can be protected against flow accelerated corrosion by the use of high chromium steel inserts in susceptible areas. In special cases the entire casing may be constructed in cast 12%Cr steel. Moving blades As is the case for the HP cylinder, martensitic steels are used for front stages of blading in the LP cylinder. For the longer blades in the rear stages (typically the final two stages) stronger materials are required and precipitation hardened stainless steel, e.g. PH-15Cr5Ni, is used. Special protection against water droplet erosion is required for the last stage blades, which has been described in section 13.3.1.
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Stationary blades Stationary blades may be cast or fabricated from plate. In the former case either cast iron or cast 12%Cr steel is used. In the latter case hollow blades are fabricated from low carbon 12%Cr steels or a similar material. For impulse turbines diaphragms are also typically fabricated from low carbon 13%Cr steel. Bolting Bolts for LP cylinders require relatively high strength and toughness. A typical material is 21CrMoV5-7 heat treated to a proof stress ~600 MPa.
13.4
Conclusion
The industry’s approach to material selection is founded on the application of materials and processes with proven reliability. Many of the materials applied today were developed in the 1960s and long experience and extensive investigations have enabled optimisation of their properties through control of chemical composition and heat treatment. These materials have up to 40 years of successful service experience behind them. At the same time, where there is clear advantage in exploitation of newer materials, they are introduced. However, this is only done after thorough investigation including the manufacture of prototype components and the long term characterisation of properties. The issue of stress corrosion cracking in low temperature rotors has been a significant source of unreliability for nearly all turbine manufacturers. Based on analysis of service experience and long term laboratory testing in realistic environments, a design assessment procedure has been developed to ensure that incidents of stress corrosion cracking will not occur in new turbines. Additional measures against flow assisted corrosion and water droplet erosion are also available. The reliability of all of these material technologies is ensured through rigorous qualification of the supply route.
13.5 1. 2.
3.
References
J F Norton, A Strang, ‘Improvement of creep and rupture properties of large 1%CrMoV steam turbine rotor forgings’, JISI, February 1969. C Berger, R B Scarlin, K-H Mayer, D V Thornton, S M Beech, ‘Steam turbine materials: high temperature forgings’, in Proceedings of COST 501 Conference on Materials for Advanced Power Engineering 1994, Lie`ge, October 1994. K-H Mayer, T-U Kern, M Staubli, E Tolksdorf, ‘Long term investigation of 24 production components manufactured from advanced martensitic 10%Cr steels
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for 600C steam turbines’, in Proceedings of Parsons 2000: Advanced Materials for 21st Century Turbines and Power Plant, Cambridge, July 2000. 4. J E Bertilsson, G Faber, G Kuhnen, ‘Fifty years of welded turbine rotors’, Brown Boveri Review 12, pp. 467–473, 1981. 5. M Taylor, D V Thornton, ‘Experience in the manufacture of steam turbine components in advanced 9–12%Cr steels’, in Proceedings of IMechE Conference on Advanced Steam Plant, London, May 1997. 6. D Kalderon, ‘Steam turbine failure at Hinkley Point A’, Proc. Instn. Mech. Engrs, 1972, vol. 186. 7. J L Gray, ‘Investigation into the consequences of the failure of a turbinegenerator at Hinkley Point ‘‘A’’ power station’, Proc. Instn. Mech. Engrs, 1972, vol. 186. 8. J M Hodge, I L Mogford, ‘UK experience of stress corrosion cracking in steam turbine discs’, Proc. Instn. Mech. Engrs, 1979, vol. 193. 9. B W Roberts, P Greenfield, ‘Stress corrosion of steam turbine disc and rotor steels’, Corrosion-NACE, 1979, vol. 35. 10. S R Holdsworth, M Nougaret, B W Roberts, D V Thornton, ‘Laboratory stress corrosion cracking experience in steam turbine disc steels’, in Proceedings of Conference on Steam Turbine Stress Corrosion Cracking, Baltimore, March 1997, EPRI TR-108982. 11. EPRI 1999, Turbine Steam Path Damage: Theory and Practice, Chapter 29, ‘Flow accelerated corrosion’.
© Woodhead Publishing Limited, 2011
14 Steam turbine upgrades for power plant life management and performance improvement F . C . M U N D , Alstom Power, UK
Abstract: Retrofitting turbines in order to extend the life of an existing plant to improve performance, reliability and/or reduce risk has become an attractive option for owners of mature steam turbine plants. Extensive retrofit experience enables vendors to combine modern blading technology and material advances to bespoke solutions beyond a like-forlike replacement. Considering the plant and future operation as a whole, turbine type, admission control and low pressure (LP) turbine exhaust area can be selected, potential power uprates considered and implications on the balance of plant evaluated in such a way that the turbines are designed to maximise benefits including an economic evaluation of the upgrade scope. Efficiency levels equal to modern new equipment standards can be achieved in many cases. Key words: steam turbine, retrofit, performance, design, installation.
14.1
Introduction
The worldwide electricity demand has increased significantly since the 1970s, which has been supported by the installation of fossil-fired and nuclear power plants. Figure 14.1 shows that, of the currently installed capacity, one-third is over 30 years old, approaching the end of their original design life. For this ageing installed base there is the option to extend the life of the existing stations beyond their original design life with selected rehabilitation works and/or exchange of equipment for improved performance. The exchange of major internal turbine components such as the rotor and steam path (blades) for an upgraded design is commonly referred to as turbine retrofit. The practice of steam turbine retrofits has been firmly Chapter 14 © Alstom Power Limited, 2011 (which own all intellectual property rights in Chapter 14, including but not limited to copyright)
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14.1
Installed generating capacity versus age (source: Alstom).
established since their introduction on fossil units in the 1980s. This chapter will focus on steam turbine retrofits, discussing drivers, products, performance, mechanical design and, finally, the installation of the turbine cylinders. High steam temperatures for conventional fossil-fired steam turbines led to typical turbine design lives of around 20–30 years. Due to less onerous steam conditions, nuclear steam turbines can have a design life of up to 40 years, in line with the typical 40-year plant licenses. In the course of the current nuclear renaissance many licenses are being extended in time and output (thermal power uprates), creating a significant market for nuclear turbine retrofits. The scope of turbine improvements can range from repair and refurbishment to a turbine upgrade or retrofit. Historically, poor sealing, blade erosion and stress corrosion cracking in LPs were common problems faced by plant operators. Depending on the scope of the problems and funds available, a refurbishment project could include the repair or replacement of individual stages or extend to the replant of the complete turbine shaft line. An increased scope rather than simply addressing reliability issues is frequently justified on an economic basis (return on investment) for high pressure (HP) and high and intermediate pressure (HP-IP) fossil cylinders, due to the significant potential for improved efficiency. Initially, turbine retrofits were considered on a like-for-like basis with the thermodynamic design conditions essentially unchanged. Typically, an HP retrofit would
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involve the replacement of the entire inner module, including rotor, steam path and inner casing. The replacement of the critical components extends the turbine life and better sealing and improved blading technology provide attractive performance benefits. Due to the longer blades, inherently higher efficiency and less degradation, performance improvements for IP retrofits are limited and costs are generally higher. Typically, a mechanical driver or a change of the cycle conditions is required to justify the retrofit of an IP turbine. During recent years, the opportunity of HP turbine retrofits has been used to reconsider the original turbine design point, in particular steam flow capacity. Increased revenue through additional capacity strengthens the economic case for a retrofit. Also, with more demanding environmental regulations and the requirement of operating environmental equipment such as flue gas desulphurisation (FGD), the associated parasitic load required to support environmental equipment can be compensated with power gains from increased steam flow rates. The benefits of turbine retrofits can be summarised in three categories: reliability, efficiency and capacity: . . .
Regarding reliability, the extended life comes with extended maintenance intervals and maintenance and/or operational issues can be resolved and costs reduced. Improved efficiency reduces fuel consumption and CO2 emissions. Increasing the capacity is an effective way to increase power output.
14.2
Drivers
Over the last decade steam turbine retrofitting has become economically attractive. About a third of the installed generating fleet is older than 30 years and major overhauls and turbine retrofits are inevitable. Retrofitting steam turbines of existing plants is a viable economic alternative to new-build due to reduced capital costs and it has become an established business sector with its beginnings in the 1990s. Power shortage predictions for the near-and mid-term future increase the pressure on existing plants to improve the heat rate and recover parasitic losses from environmental equipment such as FGD and scrubbers. As discussed in the introduction, a variety of benefits are associated with retrofits and different market drivers have developed. Extended life, reliability and availability, improved heat rate and capacity all translate into financial benefits for the plant owner. Figure 14.2 shows the typical net income over time for an existing power plant and the potential benefits from retrofits, illustrating the short payback periods for retrofits.
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14.2 Typical net income versus time for power plants before and after a retrofit (source: Alstom).
The typical drivers for implementing a retrofit of steam turbines in power plants are described in more detail below.
14.2.1 Life extension With growing environmental pressures and regulations, changing power generation policies and forecast requirements for power, there is a strong driver in many regions to extend the life of power plants. For nuclear plants, licensing of up to 60 years is becoming more common. Apart from the obvious reliability and availability issues associated with ageing plant, the extension of plant life may also provide a strong and immediate economic case for retrofits to improve performance (output and heat rate), while at the same time addressing any current reliability, availability, maintenance issues and reducing the risk to generation capability.
14.2.2 Availability, reliability and risk management Ageing steam turbines suffer from reliability problems, requiring shorter maintenance intervals and more frequent repairs or component replacement. Even without major operational problems, a unit with high service hours represents a risk to availability and an increased risk of a major
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incident (for example risk of stress corrosion cracking) as it approaches its end of design life. Retrofitting will reset life, solve the maintenance and reliability risks and reduce the risk of forced outages.
14.2.3 Performance improvement Efficiency improvement has become crucial with the deregulation of the power market and increased competition and is one of the key drivers for retrofits. Increasing fuel prices and the introduction of CO2 emission controls provides a strong case for efficiency improvements in fossil-fired stations. In some instances the performance improvement can offset restrictions and emission controls. The nuclear fleet is also driven by the growing demand for power to improve performance. Retrofitting steam turbines in nuclear and fossil-fired plants with state of the art technology can translate into significant output gains and heat rate improvements (see also section 14.4.8), particularly on LP turbines where there is often an opportunity for substantial performance improvement through an increase in last stage exhaust area.
14.2.4 Extended power uprates In order to address the growing demand for more power from the installed base, some plants are being licensed to increase thermal power. Significant uprates for nuclear plants will typically require the retrofit of the complete turbine shaftline, to match the increased steam flow, maintain integrity and achieve high performance. This can yield major improvements, with some uprates increasing output by up to 30 %. For their fossil-fired counterparts the typical uprate is in the region of a 10 % power increase, but the possible increase in steam flow is often restricted by environmental regulations.
14.3
Product selection and specification
This section describes some of the key issues to be considered, by customers and vendors alike, during the planning, preparation and specification of a turbine retrofit project.
14.3.1 Improvement potential Once the interest for a turbine retrofit is initiated, the opportunity should be taken to investigate the best combination of benefits tailored to the immediate and future operation of the plant. This may look well into the future due to the substantial design lives of the turbines of 20–40 years. In order to develop a suitable specification for a retrofit, a far greater
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technical interaction between the vendor and plant owner is required compared to defining new equipment. Once the current and future operating conditions are defined the vendors are then in the position to assess the potential improvement for different retrofit design options and operating conditions. Depending on the complexity and extent of the design changes (terminal conditions, pinch points, steam flows), this can be straightforward or might require some collaboration between the station and the vendors. In some cases, for example large uprate projects, it can be beneficial to perform an evaluation of the wider effects of the uprated conditions on the balance of plant. A working group involving the steam generator or reactor manufacturer, the utility and the turbine retrofit vendor provides the necessary technical and operational expertise to determine the best improvement design point, translating the future vision into technical requirements. Such studies should ideally be based on the existing unit recent performance test information as a starting point, and should run various scenarios, balancing various hardware modification options accommodating the relevant requirements against the required investment. These scenarios may consider wide-ranging adjustments to the steam cycle, including modification of final feedwater conditions, changes to crossover pressures, changes to extraction pressures, changes to the HP and LP capacity, increase of LP exhaust area or changes to the mode of operation and valve points.1
14.3.2 Steam turbine types Two conventional steam path blading technologies are generally applied for large steam turbines: ‘reaction’ type and ‘impulse’ type blading (RTB and ITB, respectively). For impulse stages (low reaction) the stage pressure drop occurs predominantly across the stationary blades, whereas for reaction type stages (nominally 50 % reaction) the pressure drop is distributed evenly across the stationary and moving blade row. This has led to different construction types. Impulse technology is typically associated with disc and diaphragm construction and monobloc rotors, where the fixed blades are mounted in robust diaphragm rings (see Fig. 14.3) and reaction technology with drum rotor construction (see Fig. 14.4). Welded drum rotors facilitate sections of condition specific material and the voids between the sections reduce weight. For impulse technology with limited space between discs for welds, monobloc rotors are used. The quality of modern forgings facilitates HP and IP turbine rotor designs without a centreline bore. Without a bore both welded and monobloc rotors benefit from fewer inspections compared to existing equipment.
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14.3 HP inner module impulse turbine retrofit with disc and diaphragm construction (source: Alstom).
14.4 HP inner module reaction turbine retrofit with drum rotor construction (source: Alstom).
Both technologies are competitive, as most effectively demonstrated by their successful persistence in the market place (particularly for new equipment). Impulse technology is characterised by fewer, robust stages featuring low disk and diaphragm leakages for improved performance and low axial thrust, whereas reaction technology offers more stages (approximately twice the stage count of ITB) to improve performance. For a given set of thermodynamic conditions, a higher stage count will reduce the loading per stage, which lowers velocity and consequently reduces losses, resulting in higher stage efficiency. In practice, stages are designed for certain ratios of heat drop to rotational speed. Reducing velocity is restricted as the rotor speed is constant and root diameters are limited by rotor dynamic stability. Often the maximum number of stages is not defined by the available axial space but controlled by the rotor dynamic limits. The higher stage count would suggest that RTB is more efficient than
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ITB, but in practice ITB is competitive due to reduced leakages. A retrofit provides the opportunity to revisit the selection of blading technology type and turbine construction. However, when it comes to retrofit of a particular turbine cylinder, there is normally an advantage to one of the technologies. This can be affected by the restrictions and duty imposed by the existing turbine and how the solution must be configured with each technology to accommodate this, but other evaluation criteria and customer experience and preference play an important role as well. Due to the significant differences in axial thrust, the potential conversion of an impulse to a reaction type turbine involves thrust considerations at an early stage. For opposed flow cylinders the axial forces are balanced by the opposed direction of steam flow together with the balance piston, and changes in blading thrust can be counterbalanced to a degree by changing the balance piston diameter. Benefiting from the opposed blade thrust, double flow architectures are less restricted by thrust limitations. Where single flow cylinders are arranged on one shaft in opposite flow directions a conversion of both opposing cylinders can be considered, but in some cases (an individual single flow cylinder) a conversion may not be feasible. Details on the differences between impulse and reaction turbine technology are further described in textbooks, such as that of Marlow and Brown.2
14.3.3 Steam admission control The type of steam admission control is an important characteristic of a steam turbine and is chosen based on the current and future load profiles of the plant. The HP turbine is configured either with full-arc admission (FAA) or partial-arc admission (PAA) facilitating the steam to enter the first stage of the turbine via the entire inlet annulus or via partitions of the inlet annulus, respectively. FAA HP turbines are typically controlled at less than full load conditions by throttling the steam to a lower pressure via the HP control valves all operating in unison. This throttling reduces the steam flow through the valves to control the pressure at the turbine inlet and adjusts the steam flow provided by the steam generator to the mass flow to be swallowed by the HP turbine to produce the required total shaft power and associated MWe generation. Also called ‘throttle control’, this mode of operation is associated with large throttling losses at lower loads. This can be offset to a degree by operating with sliding (reduced) boiler pressure, or so-called ‘‘hybrid’’ control, combining sliding boiler pressure with throttling. In HP turbines configured for partial-arc admission, the inlet annulus to the first-stage nozzles is divided into separate nozzle arcs, and each arc is connected independently to its own control valve. When operating with
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constant boiler pressure, load is reduced by closing each valve in sequence (effectively varying the extent of the arc of the first stage that admits steam). With this mode of control, also called ‘nozzle control’, higher levels of cycle efficiency compared to ‘throttle control’ can be maintained while operating at less than full load, because throttling losses are reduced compared to FAA. With one control valve closed and the remaining control valves open, there are no real throttling losses, but the steam flow is reduced because it is only admitted over (nominally) three-quarters of the inlet annulus. Typical configurations are ‘3+1’, where one control valve feeds one arc (nominally one quadrant of the inlet annulus), and three control valves feed the remaining segments (nominally three-quarters of the inlet annulus) or ‘2+1 +1’, and more exceptionally ‘1+1+1+1’ For a ‘3+1’ with three control valves open feeding three-quarters of the arc and one control valve closed, depending on the heat drop of the first stage, the turbine will typically swallow in the region of 90 % of the full capacity for all valves wide open. Additional control for further reduction in mass flow is then delivered by throttling the remaining valves together. For PAA admission HP turbines, the control (first) stage experiences higher steam flow velocities and forces and, hence, must be robustly designed. In particular, the rotating blades must be designed to withstand the cyclic loading associated with closing a complete nozzle sector. A mixing volume (space) should also be provided downstream of the control stage to enable uniform circumferential mixing of the flow before entering the full arc second stage (see Fig. 14.5). On some occasions, the provision of this mixing space can reduce the stage count by one or two in comparison to FAA due to the additional axial space required (compare Figs 14.5 and 14.6). With all valves wide open, a turbine configured for full arc admission will be most efficient, due to the improved efficiency of the first stage. However, this benefit deteriorates as throttling progresses, to the point that the performance for PAA is superior. The crossover point typically occurs in the region of 3–5 % throttling (i.e. when operating at around 95–97 % below the full turbine capacity, defined by VWO; see Fig. 14.7). The decision as to whether to adopt partial arc admission on an full arc admission should therefore be based on economic evaluation based on the expected cycle performance (heat rate and power) over the proposed MW loading profile configuration. PAA is the more common mode currently applied due to the efficient operational flexibility it can provide.
14.3.4 Design conditions and life The specification of design conditions is critical in determining the success of a steam turbine retrofit to meet the plant requirements. Ideally, any available performance test information (or even the original performance
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14.5 PAA eight-stage impulse inner module retrofit and typical firststage moving blade (source: Alstom).
14.6 FAA nine-stage impulse inner module retrofit and typical firststage moving blade (source: Alstom).
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14.7 Variation of heat rate with load FAA versus PAA (Hogg and Stephen3).
acceptance tests) should be made available with the turbine retrofit specification. This will identify current thermal performance of the major components of the balance of plant, which may be significantly different from the original design. It will facilitate clarification of the required swallowing capacity, associated throttling margin and the turbine inlet terminal conditions (pressure and temperature). It also will enable vendors to provide proposals on a consistent and accurate basis, using the as-tested performance as the basis for the thermodynamic modelling of retained equipment. This will, in return, provide the customer with the most accurate assessment of the expected post-retrofit performance on which the economic evaluation can be based. The design life for fossil-fired steam turbines is around 20–30 years, whereas it is up to 40 years for nuclear applications. This can be achieved due to lower steam temperatures in nuclear steam turbines, which relieve the restrictions on lifetime due to creep. In respect to the design life, specific attention must be given to the risk of stress corrosion cracking (SCC) in turbine rotors operating in nuclear power plants, more so than for their fossil-fired counterparts. This is because the wetness levels are higher; in fact
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the steam is wet through the majority of the steam expansion in nuclear plants.
14.3.5 Hot end considerations To make sure that the full potential of the existing assets is exploited, it is important to look at the plant as a whole rather than optimising the turbine on an individual basis. In particular at the hot end, fossil-fired steam generators often have very conservative original design margins. This hidden reserve can be accessed when retrofitting the HP turbine with the swallowing capacity tailored to make use of the boilers steam flow margins. Retrofitting the HP turbine increases the heat input required in the reheater if the hot reheat temperatures are to be maintained. This should be considered as to whether this additional reheat capacity is available. By changing the capacity of the IP turbine, the final feedwater temperature and reheat can be optimised. If steam flow in excess of the existing margins is considered, adding boiler surfaces can be added to the scope of the investigation. These considerations require detailed technical studies where the impact of the modifications on the retained equipment is assessed. More details and examples in operation can be found in Stephen et al.4 and Bartley et al.5
14.3.6 Cold end considerations It is also important to consider the cold end, i.e. the condenser implications. Cooling water temperature and flow and the heat transfer between the cooling medium and the condensate set the condenser pressure. In particular, for LP turbine retrofits, the customer should carefully specify the condenser pressure for the guarantee design point, based on an assessment of the normal annual variation in condenser pressure and considering the seasonal requirements for power. Failure to do this can lead vendors to select an inappropriate last stage blade (i.e. the wrong blade exhaust area) to match the guarantee design condenser pressure, rather than that which is required for optimal performance of the plant. Sometimes it is possible to upgrade the condenser, increase the cooling water flow by upgrading existing pumps or installing additional pumps or modify the source of the cooling water flow. As this reduces the condenser pressure and consequently increases power output and reduces the heat rate, it should therefore be considered early in the optimisation process and specification of retrofit projects. Conversely, conversions to cooling tower operation for environmental reasons will significantly increase condenser pressure and a retrofit with a small last stage blade may provide efficiency improvement.
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14.4
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14.4.1 Steam cycles Steam cycles come in an extensive spectrum of configurations and design conditions and generator output can vary significantly from 100 MW to 1 GW for a single unit. Power plants can be arranged in a single shaft line with single and/or combined cylinders driving a single generator (tandem compound) or for larger plants as ‘cross compound’ with individual shafts driving two or more generators in parallel. Subcritical fossil powered conventional cycles are typically designed for main steam pressures of nominally 160 bar and 540 8C steam temperatures. After an initial expansion through the HP turbine the steam is typically reheated back to 540 8C before expanding through an IP turbine and LP turbine(s). Some fossil powered cycles operate with supercritical main steam pressures of nominally 240 bar but main steam temperature typically remains around 540 8C. Nuclear powered cycles typically operate with lower main steam pressures and temperatures around nominally 65 bar and 280 8C, with the main steam being just below the saturation line and having a small moisture content. After an initial expansion through an HP turbine, moisture is typically separated and the steam is reheated and then expands through an LP turbine. Typical expansion lines for fossil-fired and nuclear cycles including reheat and non-reheat cycles are shown in Fig. 14.8.
14.4.2 Plant conversions Combined cycles (CC) combine the high temperature Brayton cycle of the gas turbine with the low temperature Rankine cycle of the steam turbine by reusing the exhaust heat from the gas turbine in the heat recovery steam generator (HRSG). The HRSG links the cycles together and replaces the steam cycle feedwater train. Temperatures of the steam cycle are similar to conventional fossil-fired cycles but pressures are generally lower (in the region of 65 bar) and vary as they are optimised for best heat recovery. Fuel prices, availability and cycle efficiencies make new CCs an attractive option, and in some cases existing plants are considered for conversions. The HSRG typically provides less heat than a conventional fossil-fired boiler and combinations of two gas turbines to one steam cycle have been put into service. Steam cycles can also be used in combined heat and power generation. Benefits from using a heating system for district heating or hot water as the
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14.8 Typical expansion line fossil-fired and nuclear cycles (source: Alstom).
steam condenser improve the cycle efficiency. Large amounts of steam can also be extracted from the turbines, where multiple extractions increase cogeneration efficiency, making it a viable consideration for plant conversions.
14.4.3 Turbine architecture Modern and more advanced blade materials and construction lift some of the old design limitations and generally allow for longer blade heights. This can facilitate changes in turbine architecture. For example, double flow first stages in HP turbines can be converted to single forward stages benefitting from higher first stage efficiency and reduced stage leakages and losses. Most steam turbines for nuclear (wet) steam cycles have double flow HP modules (operating at half speed), which can potentially be converted to single flow cylinders generally benefitting from efficiency improvements on all stages.
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14.4.4 Performance improvement Efficiency improvement Technology advancements have improved efficiency levels significantly since the 1970s. Improvements of computational methods, supported by model turbine testing, have facilitated three-dimensional blade design to improve blade performance. Advanced sealing reduces losses from leakages and longer blades are in use based on modern material technologies. The conversion of double flow stages or modules to single flow stages or modules can give significant efficiency improvements due to the increase in blade height. Improvements have also been realised by optimising the number of stages and by selecting the appropriate technology type for the application. Walker6 published a comprehensive summary. For fossil-fired cycles or an advanced gas-cooled reactor (AGR) a more efficient HP turbine comes with a reduced HP exhaust temperature and therefore reduces reheater inlet temperature. Slight additional heat input to the reheated steam is therefore required if the IP turbine inlet temperature is to remain at pre-retrofit level. If no additional heat is available it is important to understand that the reduced temperature will carry through the reheater to the IP turbine inlet and reduce the retrofit benefit, impacting on heat rate and power. More details on the individual effects are discussed in section 14.4.8. Adding to the reheater surface can increase the reheat input and retrofit benefit. Capacity change Increasing the swallowing capacity of the HP turbine has a direct impact on power gain, as the additional steam provides more energy to be converted into useful work. Assuming that the steam generator can provide additional flow (from unused design margins, change of operating conditions or hardware modifications), operating ranges of components must be checked for the ability to cope with the additional duty (for example pumps, boiler feedpump turbines and heaters). With the rest of the turbine cycle components unchanged, the additional flow (associated with the increased capacity HP turbine) will increase the HP turbine exhaust pressure and may increase the final feedwater temperature via the relevant feedwater heater extraction. Adjusting the IP turbine swallowing capacity can compensate for the change of HP exhaust pressure. It may be possible to move the heater extraction by a stage (if the final heater steam extraction is from the HP turbine), allowing for controlled feedwater temperature specification and consequently adjusting the main
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heat input in the steam generator (environmental concerns, generator capacity). More details can be found in Bartley et al.5 In rare cases, with ageing steam generators not providing sufficient steam and lacking upgrade potential, it may be considered economically viable to reduce the size of the HP turbine so that throttling losses are removed and the throttle valves can operate in a fully open position. This may also be a consideration for refiring options with biomass, which can reduce steam flows. Effect of LP blade exhaust area At commissioning of most original plant, the last stage blade height of the LP turbines has been limited by the technology available at the time. LP retrofit projects therefore offer the scope of optimising the exhaust area to the specified retrofit operating conditions via an increased blading portfolio. Best performance through reduced exhaust losses is generally achieved with exhaust velocities between 150 and 300 m/s, naturally varying with condenser pressures from summer to winter, respectively (see Fig. 14.9). At high exit velocity, typically above 300 m/s, the performance of the last stage blade is inhibited by ‘limit loading’. This is a physical mechanism in which the LSB exit pressure is too low and the transonic shock system detaches from the blade passage. Further reductions in condenser pressure result in a stronger shock wave and increased leaving loss, but do not increase the
14.9 Exhaust loss versus annulus velocity (source: Alstom).
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power output from the last stage or the turbine. This can be corrected through increasing the LSB area. At low exit velocity, below around 150 m/s, the performance of the last stage becomes inhibited by increased leaving energy (high swirl) and ‘negative root reaction’, which leads to gross flow separation. Negative root reaction is a mechanism whereby the exit pressure from the LSB is actually higher than the inlet pressure to the LSB, i.e. the flow is not expanding over the lower portion of the blade. This can lead to separation of flow in the lower regions of the LSB (towards the hub). Poor performance associated with low exit velocity can be relieved by retrofitting a small last stage blade to reduce the LSB exhaust area (and increase the exit velocity), but care must be taken not to compromise the winter performance due to raising the exit velocities at low condenser pressure. When economically viable and technically feasible, it is preferential to treat the problem of poor performance at high condenser pressure through improvements to the condensing/cold water system in order to reduce the condenser pressure. This provides additional output during summer with a longer LSB and maintains the benefit of improved output associated with a long LSB all year round. It should also be noted that for power uprate projects, significant increases in steam flow also necessitate proportional increases in the LSB exhaust area for optimal performance. Control philosophy With the deregulation of the electricity market and more restrictive environmental regulations, the load profile of many power plants has changed significantly in recent years. Base-load plants may be forced to operate at part-load conditions for extended periods of time. A retrofit project should then consider the type of steam admission. Full arc admission HP turbine configurations have a performance advantage in the region of 0.1 % heat rate or typically 0.5–1 % turbine efficiency over partial arc admission at full load, when control valves are fully open. The absence of a nozzle box, narrower first stage blades and possibly an additional stage, instead of the required mixing space for partial arc configurations, benefits turbine efficiency. However, when operating at lower loads, valve-throttling losses penalise full arc admission types and partial arc configurations exceed their counterparts’ performance (see Fig. 14.7). This crossover typically occurs between 3 and 5 % throttling (i.e. the HP turbine operating at 95–98 % of valves wide open load). For a (load cycling) fossil-fired plant the load profile will assist in selecting the most beneficial steam admission. For nuclear applications operating at base load, flow margins offer
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14.10
Efficiency versus load for FAA and PAA control (source: Alstom).
operational control and future flexibility. A typical margin of 5 % and its effect on PAA and FAA HP efficiencies is shown in Fig. 14.10, illustrating the small difference between the two options at the 100 % load condition.
14.4.5 Guarantees A retrofit project usually involves the replacement of at least one turbine component in a complete cycle. To make the project economically viable, the customer generally has certain expectations of the performance improvement of the retrofitted turbine and this translates into guarantees. In order to define and verify guarantees it is important to understand the component as well as the cycle performance. This is required when the guarantees are verified in a test and the operating conditions are different from the specified guarantee conditions. The ability to distinguish the individual component effects on overall plant performance facilitates scaling and correction of the tested performance to the specified guarantee operating conditions, if required. There are two fundamental types of guarantees, relative and absolute. .
Relative guarantees compare the pre- and post-retrofit performance and the level of improvement is guaranteed. This typically encompasses heat rate or output improvements. To verify the guarantee, pre- and posttests are performed, where there is the potential of reduced test
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Steam turbine upgrades for power plant life management Table 14.1
Guarantee alternatives for fossil-fired cycles
Scope of supply Absolute
Relative
HP IP HP–IP
Recommended Efficiency and swallowing capacity
Discouraged Efficiency or MW
LP turbine
Optional heat rate / MW
Recommended MW/heat rate
Full shaftline
Recommended MW/heat rate swallowing capacity
Discouraged MW/heat rate
Table 14.2
Guarantee alternatives for nuclear cycles
Scope of supply Absolute
Relative
HP
Discouraged
Recommended MW
LP turbine
Optional MW/heat rate
Recommended MW/heat rate
Full shaftline
Recommended MW/heat rate
Optional MW/heat rate
.
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uncertainty if the same equipment is used by elimination of systematic test uncertainties. Absolute guarantees refer to the post-retrofit performance only, and common guarantee parameters are turbine efficiency, turbine swallowing capacity, unit heat rate or unit output, depending on the scope of the retrofit.
Guarantees should be tailored to the scope of supply and are heavily influenced by the ability to carry out accurate testing to verify the guarantees. With absolute guarantees there is a level of uncertainty on the retained components, in particular if their performance cannot be tested accurately. For relative guarantees there is a level of uncertainty on the components that are replaced based on the assumptions made. It is important to find the most equitable guarantee structure for all parties involved. For fossil HPs and IPs it is simple and absolute efficiency guarantees and HP swallowing capacity are commonly used. IEC7 recommends relative guarantees for individual cylinder retrofits (as the contribution to power for an individual cylinder is relatively small) and an absolute overall plant guarantee for the retrofit of the entire shaftline (all cylinders) (see also Tables 14.1 and 14.2). For nuclear plants and LP turbines the performance assessment is more
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complex, as additional to pressure the level of wetness defines the state of the steam. The key uncertainty of an existing nuclear plant is the power split between the HP and LP turbines.
14.4.6 Verification of performance improvement A guarantee verification test or ‘acceptance tests’ is an essential part of a turbine retrofit project. As discussed in section 14.4.5, for a relative guarantee this includes a pre- and post-retrofit test, whereas for an absolute guarantee only the final acceptance test is carried out. In order to measure the performance ‘as new’, the acceptance test should preferably be carried out within eight weeks after installation. This reduces the risk of performance deterioration and damage to the turbine, and typically after that corrections for component ageing are permitted (ASME8). So-called ‘full-scale’ tests provide detailed performance of HP, IP and LP turbine performance with a minimum of uncertainty at the expense of extensive measurements and calculations. Alternative tests rely on fewer measurements and more correction curves, resulting in slightly increased uncertainty in return for lower costs. These may be preferred for nuclear units where relative guarantees are more common (ASME8) and test uncertainty reduces due to the repeated test before and after the retrofit with the same instrumentations. It should be noted that the main steam wetness is usually not measured but usually assumed constant for the pre- and postretrofit test. Flowmeters can have higher uncertainties, normally of a systematic nature. It is therefore important for relative uplift tests to retain the same flowmeter. Fossil-fired stations typically have several locations (condensate flow, final feed flow) where calibrated flow nozzles can be fitted additionally to the plant instrumentation for performance tests. In order to reduce correction errors, the tests should be conducted with the smallest possible deviation from the specified guarantee conditions (ASME8). Most fossil unit retrofit guarantee conditions are specified at throttle valves wide open and design boiler pressure. The most meaningful way to test is therefore with the HP throttle valves wide open with the boiler fired at its maximum capacity. If the supplied steam is not sufficient the pressure should be reduced to ensure that all valves are fully open. A detailed account of the plant operating conditions is needed, for example if heaters are out of service, electrical pumps required, extraction pumps operating or if emergency drains are open. To reduce uncertainty in deriving the throttle flow, any auxiliary flows to and from the cycle should be isolated, minimised or measured. Nuclear units may or may not have HP valves throttling (depending on flow margin). Test tolerances generally include a retrofit contract as they reflect the
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uncertainty associated with testing and measurement. It is important to recognise that if a guarantee is based on zero test tolerance the vendor may reduce the guarantee value to mitigate the risk of liquidated damages due to real test uncertainty (which are positive and negative). If the guarantee takes test uncertainty into account, the vendor is allowed a test tolerance equal to the evaluated (or specified) test tolerance before comparing guarantee values. The most equitable position is to have the test tolerance equal to test uncertainty. It is in the interest of both parties to reduce test uncertainty and it is recommended that a test uncertainty analysis be performed according to the relevant standards, e.g. ASME PTC 19.1-2005.
14.4.7 Monitoring Performance monitoring is the continuous collection and presentation of data to assess plant health. Coupled with software tools, basic and preliminary diagnostics and interpretation of data can be performed automatically. Use of these data can be a good indicator of long-term component deterioration or more sudden damage and is vital when decisions for forced shutdowns have to be made. A detailed periodic review of the data should be performed by the plant operator/engineer. Further reading on plant monitoring can be found in ASME9. Trending can indicate if deterioration exceeds normal expectations and needs further root cause analysis. Until 2008 ASME10 average ageing relations have been widely used as guidance for nuclear and fossil-fired turbines. However, component ageing is influenced by many factors such as plant operation, water chemistry and component materials and technology, leading to a wide range from nearly constant performance to severe and rapid deterioration. Due to this wide spread, the ASME ageing guidelines have been withdrawn. In particular, in the US rapid turbine efficiency deterioration has been experienced through copper deposition on turbine blades. The copper pipework used in feedwater heaters and throughout the plant transports as oxides through the boiler and can deposit rapidly on the HP turbine blading with a detrimental effect on HP turbine performance. It starts immediately from steam admission and can worsen the efficiency of a newly installed HP turbine before the acceptance test. A comparison of the deterioration of the HP efficiency in a 740 MW fossil unit severely affected by copper deposits to the relevant former ASME ageing guideline is shown in Fig. 14.11. It illustrates that if copper is suspected to be present in the cycle, monitoring the performance from the first steam admission on and performing benchmark tests as soon as possible is very important. Severe deposition in particular on the first stage can also cause blockages, reducing the main steam flow. More details on copper deposition are
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14.11
Deterioration in HP cylinder efficiency (source: Alstom).
discussed by Jonas et al.11, reporting capacity reduction in excess of 10 %. Some blade designs are more susceptible to deposition than others and the author’s company has not experienced reduced swallowing capacity caused by deposition on their blade designs. Some of the turbine degradation can be temporarily recovered through various cleaning methods, such as chemical cleaning or glass bead blasting. Recovery varies but can exceed 1 % in HP turbine efficiency. An increase in HP turbine efficiency can also be observed after cold starts as thermal stresses can lead to shedding of some of the deposits. The temporary recovery can range up to 0.5 % HP turbine efficiency.
14.4.8 Typical benefits Retrofit turbines are bespoke solutions tailored to each plant and the actual benefit can vary significantly from plant to plant. This section documents some typical performance benefits. For proper evaluation of the economic potential, utilities should request an assessment from retrofit vendors. The main steam flow has the biggest influence on performance with a ratio of approximately 1:1 for power output (the heat rate largely remains unchanged), making power uprates by increasing the HP swallowing capacity such an attractive solution. Cylinder efficiency as a component improvement contributes as a ratio of the cylinder share in overall power. Due to the different cycle architecture, the effects are different for nuclear and fossil-fired plants. An overview comparing typical nuclear cycles to fossil-fired single reheat units can be
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Table 14.3 Cylinder efficiency effects for typical nuclear and fossil-fired single reheat cycles (Kearney et al.12) Effect of a 1 % increase in Effect of a 1 % increase in turbine cylinder efficiency turbine cylinder efficiency on generated power on heat rate HP turbine (fossil-fired cycle) HP turbine (nuclear cycle) IP turbine (fossil-fired cycle) LP turbine (fossil-fired cycle) LP turbine (nuclear cycle)
0.3
0.2
0.35
0.35
0.17
0.17
0.45
0.45
0.6
0.6
found in Table 14.3. These should only be used as a guideline as they vary depending on cycle architecture. It is not uncommon that upgrades from 1970s and 1980s technology HPs to a modern retrofit can provide 7–10 % efficiency improvement and hence around 2–3 % power gain without increasing the main steam flow. IP improvements in efficiency can range between 2–4% cylinder efficiency, equivalent to 0.3–0.7 % increase in output. A typical gain from a nuclear LP retrofit would be 1–3 % efficiency, translating to 0.6–1.8 % increase in generator output. For LP retrofits the exhaust area can often be increased as longer LSBs are available. The potential improvement heavily depends on the exhaust conditions, but 0.5 % to 1.5 % can generally be realised.
14.5
Mechanical design
14.5.1 Measurement process The fundamental difference between new equipment and retrofits is that the retrofit has to be designed to fit into a well-defined space and connect to the interfaces of the retained equipment. The geometrical details and definition of these are therefore an essential part of the ‘input data’ to start the design process. If the original equipment manufacturer (OEM) is retrofitting an inhouse design, all relevant details should be available from drawings. However, non-OEM projects are common and these data can then only be obtained via measurement surveys. Typical opportunities for measurements are routine maintenance outages, with the cylinder being opened for inspection, but also forced unplanned shutdowns may be a possibility. If geometry data are required prior to an outage, a reduced measuring effort can then be an alternative with a
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minimal degree of dismantling (e.g. outer casing and pedestal covers). This reduces downtime but may also reduce retrofit options when there is a lack of detailed data. In any case, the measurement activities are carefully coordinated with the station teams to ensure access to the required components and minimum impact on the outage program. The tools used in these surveys are listed below: .
.
Conventional hand measuring equipment, such as pi tapes, digital callipers and inside micrometers are adapted specifically for cylinder measurements. Basic features can be assessed with an accuracy in the region of 0.25 mm. Portable coordinate measurement machines (CMMs), as shown in Fig. 14.12, record the measurements directly into a three-dimensional computer aided design (CAD) program, offering a degree of accuracy in the order of 0.13 mm. The generated model can comprise surfaces,
14.12 CMM in use in rotor (source: Alstom).
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14.13 Three-dimensional laser scanning device (source: Alstom).
.
curves, lines and points and provides an accurate record of a component. A three-dimensional laser scanner (Fig. 14.13) generates 3608 point clouds which, following a certain amount of post-processing, can generate CAD models of pipework and structural steelwork to an accuracy in the region of 0.5–1.25 mm. While not as accurate as the CMM data, extensive use has been made in nuclear plants, for complex pipework and condenser surveys where less accuracy is acceptable in return for a high level of detail. A detailed example has been described by Jones and Nelmes13.
Modern and conventional methods can be applied in parallel to create independent sets of measurements for reliability. Based on the accuracy required, CMM or hand measurements may be preferred for specific details (e.g. glands). On the other hand, for large LPs the CMM may not be practical, limited by the extension of the measuring arm. Using both sets of
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measurement data ensures that an accurate and comprehensive model of the existing equipment can be created. In addition to the measurements, a comprehensive portfolio of photographs is taken during the course of the survey. These photographs can be used to clarify any areas of ambiguity and their contribution to the survey content often proves to be of great value. More details can be found in Holmes and Lord14.
14.5.2 Design life The main physical conditions and the selected component material affect the design life. Operating hours and cycles, variations of operating conditions, site conditions and maintenance all influence material life and potential failures such as low cycle fatigue and high temperature creep rupture. Fossil-fired cycles generally have a design life of 20–30 years, equivalent to in the region of 200 000 operating hours. Operating at lower steam temperatures, nuclear cycles typically have design lives in the region of and beyond 40 years. Using more advanced materials at higher costs will in return increase the design life benefitting from higher creep and fatigue limits. Transient conditions such as startups and extensive load cycling reduce the design life (thermal fatigue) and have to be accounted for. Typical values are around 5000 start ups during a 200 000 h design life at full load and design steam conditions. Some components are exposed to more severe conditions and require robust design to meet the above life criteria. PAA admission is particularly demanding. When nozzle sectors are closed at lower loads, the dynamic loading on the first stage increases and higher steam velocities make the blades more susceptible to solid particle erosion. Larger steam temperature swings at lower loads also increase rotor and casing thermal stresses. For the rotor design particular attention is paid to the fillet radii at changes of rotor diameter to minimise surface stress concentration. This limits fatigue life expenditure from thermal cycling or oscillating fault torques. Another important design criteria are peak torques from electrical malfunctions or malsynchronisation.
14.5.3 Maintenance requirements As discussed in the previous section, the design life of components is limited. It is therefore important to monitor and document operating hours, cycles and variations in operating conditions in order to determine the ‘equivalent operating hours’. Maintenance schedules and remaining life assessment go hand in hand, in particular for critical components. For high temperature components, where thermal fatigue associated with machine starting or load
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changing significantly contributes to component life exhaustion, adequate instrumentation should be fitted to assess the actual stress and temperature environment the component was operating in. During a scheduled maintenance outage, the state of material and components should be inspected by various non-destructive evaluation (NDE) methods. Additional to monitoring equivalent operating hours, performance should be monitored. As discussed in section 14.4.7, cycle parameters can be a valuable indicator of system health and assist in maintenance scheduling. Steam chemistry can also have an adverse effect on component performance and life through failure mechanisms such as stress corrosion cracking and should therefore be monitored. Typical major overhaul intervals (covers off) are 100 000 equivalent operating hours, approximately 10–12 years of operation (less for a two-shift operation). It makes sense to consolidate maintenance work and as an example, the outer casing bolt material in modern retrofits is usually selected to have retightening schedules to match. Borescope ports are provided to allow for inspection of the inlet and exit areas of HP and IP elements. Minor inspections are recommended every 25 000 operating hours including visual inspection of the outer casings, borescopes and bearing inspections, as well as checks of the valve function and pedestal alignment. It should be noted that as first retrofits were installed in the 1980s the vast majority of the retrofits in service have not been due for major inspection. Continuous successful service confirms that adequate design philosophies have been applied, but more detailed assessments will be available in the coming years.
14.5.4 Concept design process Unlike the design process for new equipment, where the design starts with the steam path followed by the casings being designed around it, the retrofit design process evolves from the ‘outside in’ and focuses on working from the steam path outwards and the existing, retained hardware inwards. Beyond reliability improvements, replacing inner (and sometimes outer) casings lifts spatial design restrictions and may be worth considering in spite of the additional costs. A successful design approach is to use standard components with proven technology to create a bespoke retrofit. Each retrofit solution is customised with confidence from considerable service history and benefits from design and manufacturing advantages. The starting point of the concept design process is a set of input data, comprising the geometrical details (measurement survey or OEM drawings, as described in section 14.5.1) and thermodynamic data. The thermodynamic data usually in the form of heat balance diagrams (HBD) define the
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design turbine inlet and outlet conditions, but also other requirements specified by the plant owner, such as operational requirements and seasonal variations affecting the design. For a retrofit it is critical that the thermodynamic data are based on actual operation. One key design aspect for retrofits is not only to fit the geometrical requirements but also to have minimal impact on post-retrofit operation and maintenance. Maintaining operating conditions for retained components as close as possible ensures similar interface loads and thermal expansion. The concept design process is characterised by close interaction between the thermodynamic and mechanical disciplines to optimise within mechanical, aerodynamic and thermodynamic constraints. For example, the moving blade width is driven by mechanical requirements, whereas blade height and base diameter are optimised for thermodynamic performance. To be able to create detailed solutions for retrofit proposals, vendors have developed integrated design suites. It facilitates rapid assessments of different design options and provides performance predictions along with layout, design documentation and product costs. In fact, all hardware data are combined to a virtual assembly using CAD. For additional validation the data can be cross-checked with similar type machines and layout drawings are derived directly. A simplified flow diagram of the concept design process can be found in Fig. 14.14. Fundamental decisions on turbine technology and architecture, the type of steam admission, the number of stages and the position of the extractions
14.14 Flow diagram concept design process (source: Alstom).
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are made early on in the design process and the main mechanical considerations are discussed in the following. Thrust The decision to use ITB or RTB technology is sometimes linked with the resulting thrust from the rotor (thrust bearing capacity). RTB turbine sections come with a higher thrust due to the higher pressure drop across the rotating stages (see also section 14.3.2). This is not as relevant for double flow turbines, where the thrust is balanced between the opposing flows. For single flow cylinders the thrust can be balanced by adjusting the geometry of the balance piston. Such changes may also need assessment in terms of rotor dynamics as the geometry of the rotor fundamentally changes and in terms of performance to evaluate the effect of a change in balance piston leakage flow. The different stage count also generally requires replacement of the inner cylinder. Control stage design The cyclic loading on the first stage blades when traversing through open and closed nozzle sections is challenging for mechanical fatigue. A robust first (‘control’) stage design has to be employed to withstand the large pressure and temperature variations. The design features that the author’s company employs for special control stages are the integral shroud and pinned root, as shown in Fig. 14.5, and have decades of non-failure service history. The principle of this so-called ‘alternate torsion’ design is based on pre-twist: elastic stress on the assembly ensures continuous coupling on alternate contact faces for increased stiffness and high damping. For larger HPs some OEMs have employed reverse flow double first stage design to ease mechanical loads at the expense of poorer efficiencies and large axial space requirements. Replacing these with a single forward flow control stage improves efficiencies and creates space for additional stages, but adds substantial size to the first stage blades. Occasionally, advanced materials with higher creep strength have to be applied.
14.6
Installation
14.6.1 Design for installation Installation is an integral part of completing a retrofit project and it is important to incorporate aspects of short and simplified installation into the
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retrofit design. Delays result in lost generation revenue for the plant and are therefore generally associated with liquidated damages for the turbine installer. An efficient design for installation minimises dressing or machining on site to small components with the aim to use tools that are easily available on site. Designing the interface connection features for flexibility and easy machining is a prerequisite for rapid installation. Uncertainties of the existing hardware and manufacturing tolerances of the new components cumulate at installation to potential installation difficulties. For non-OEM projects the data for the existing hardware are associated with measurement uncertainties (in particular when survey data were taken from a sister unit), whereas OEM projects are faced with manufacturing tolerances and in-service wear and distortion. Designing flexibility into the interfaces between the existing and the retrofit equipment to compensate for these tolerances facilitates a good fit for a range of positions. The application and adjustment of the connection features requires detailed documentation, including drawings and tolerances to guide the installation process. Some common flexible design features for interfaces from the author’s company are described by Holmes and Lord14 and are summarised in the following.
14.15 Inlet connection with sleeve (source: Alstom).
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Inlet connection To seal the outer and inner casing or nozzle boxes at the inlet, ‘high/low’ rings are typically fitted. Due to wear and distortion these generally need replacing, providing the opportunity to apply a more flexible design feature. A sleeve and piston ring (see Fig. 14.15) with a machining allowance on the outside facilitates cleaning and correct measurement of the bore and machining of the sleeve to match. Due to the additional material allowance the outside of the sleeve can be machined eccentric to the bore, thus compensating for misalignment. This principle is also applied to ‘bell-seals’ using ‘floating rings’. Alignment keys and packers Alignment is achieved using keys and packers and existing ones are generally retained. For existing direct fit keys (without packers), the new design incorporates packers to axial (Fig. 14.16) and transverse alignment. In the majority of cases, calculations have shown that the top transverse keys can be eliminated, thus significantly simplifying installation.
14.16 Inner casing packers for axial alignment (source: Alstom).
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14.17 Casing baffles (source: Alstom).
Casing baffles The positioning of casing interspace baffles can be simplified by a combination of axial and radial clearances (Fig. 14.17). The clearances have evolved with retrofit experience and can easily be adjusted with machine tools on site. Steam extractions When inner casings are replaced (often necessary for LP retrofits) the steam extractions need to be reconnected. Generally located in the bottom half of the cylinder, the extraction pipes are difficult to survey. Standard practice is therefore to include a short length of pipe that can be cut to length and can be used to accommodate some degree of misalignment. This principle is also applied for blow-down and pressure tappings.
14.6.2 Alignment Suppliers have different principles for shaft alignment. The bearings can be positioned following the natural rotor catenary, i.e. the shaft coupling gaps
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are parallel or the bearings can be set on a level plane. Parallel coupling gaps benefit from reduced bending loads across the coupling, minimising the risk of fretting. Changing the original heights is generally avoided as it implies moving major components such as bearings, pedestals, generator(s) and exciter(s) and this would have a major impact on the installation duration and cost. Welded rotors are much stiffer that shrunk-on disc LP rotors and their natural sag differ significantly. To avoid or minimise any height changes resulting from rotor type changes, welded rotors can be designed close to the original catenary. This has to be incorporated early into the design phase to ensure that rotor stresses and bearing loads are acceptable (see Holmes15).
14.6.3 Technical field advice The retrofit supplier, plant personnel or a third party can carry out the installation of the retrofit. Regardless of who installs the equipment, the turbine manufacturer should provide technical field advisors (TFAs) to support the installation. The role of the TFAs is to provide technical advice on disassembly, inspection, adjustment and re-assembly, such as: . . . . .
advice and assistance with outage planning; providing the technical link between the station and retrofit supplier; analysis of the strip-down measurements and NDE results; approval of all re-assembly settings and clearances; compiling and issuing the final installation report.
Involvement of the TFA starts from the planning phase on, including early site visits to ensure that station personnel are familiar with all aspects of the installation.
14.6.4 Installation planning and management Retrofits require rapid and predictable installations, typically in the region of 20–35 days, timed for planned outages and not impacting on the concurrent station’s maintenance work. As common resources such as laydown space, cranage and access are shared between the maintenance and installation activities, careful planning is required to manage the complex activities within the physical and time constraints. Fundamental for successful timing is robust on-time delivery. In order to keep to the time frame during which a retrofit must be installed, it is essential that all of the activities likely to affect an outage are considered and planned well in advance. Additional complexity is added for nuclear retrofits, as the total radiation exposure has to be managed and minimised.
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This increases the planning and installation efforts significantly in comparison to fossil-fired retrofits. For a retrofit project there is typically a strong financial incentive, driven by contractual installation duration-related penalties/incentives, to use cost– benefit analysis to explore the investment in special assembly tools or rigs in order to reduce assembly time and/or risks to the installation duration. A detailed planning sequence supported by analysis of the critical path in respect to risk and opportunities for accelerating the schedule is required. Over the past years retrofit suppliers have developed installation management processes for the planning, preparation, mobilisation and execution of retrofit installations with the aim of improving the installation quality (on-time installation of the defined scope with an incident-free return to service). Based on ‘best practices’ interdisciplinary experiences are shared and lessons learned are implemented, thus continuously evolving and improving the process framework. Emphasis is placed on contingency management based on scenario planning, improved management of subcontractors, internal communication and reporting, structured communication with the plant owner (to ensure timely clarifications) as well as early identification of standard procedures for plant owner inspection and handover. Specialists assist with the installation specification including the equipment, interfaces between existing and new equipment, necessary machining activities, the scope of the installation services and quotations from internal and external installation contractors. A division of work/ responsibilities is developed, to ensure that all parties are clear on their responsibilities. The backbone of the process is a continued evaluation of risk to the safety of personnel, plant equipment, the outage duration and finally to the installation budget. For particularly complex retrofit projects, a specialist team may also be mobilised to perform a site survey during the tendering phase to highlight and capture any site specified risks and restrictions. For complex installations three-dimensional visualisation systems can be used to simulate the outage and assist the planning and management.16 This facilitates integrity checks but more so familiarises the team with the planned sequence and provokes constructive dialogue within the team. In summary, topics included by the installation planning are: . . . . . .
outage schedule; site logistics – components, manpower and installation resources; knowledge management – quality process and documentation control; environment, health and safety plans; hardware design for installation analysis; contractor management.
Practical aspects helping to reduce installation times are also:
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Factory trial assembly (shop built). Rotors are generally bladed off-site, allowing shorter installation times and high standards of workmanship. Furthermore, after assembly highspeed balancing and over-speed testing is performed before delivery on site. When installing the bladed rotors it is essential to position the rotors such that they rotate on the same centreline as during balancing, i.e. the journal centreline. This ensures low levels of vibration and avoids in situ balancing. Retrofits are generally designed not to impact on the remaining cycle and operating procedures. However, if the customer has specific concerns, such as extended holding periods during startups, the retrofit installation may be an opportunity to address these issues. With additional instrumentation rotor stresses can be assessed and startup times can potentially be reduced.
14.6.5 Potential durations The duration of a retrofit installation depends on the scope as well as the location and site restrictions, but is also an economic choice. Where for nuclear plants it may be critical to fit the installation into a refuelling outage, smaller fossil-fired plants may be able to afford a longer outage with a smaller crew. As discussed in section 14.6.4, pre-assembling off site can reduce the installation time on site, but will have other cost implications. For example, for a nuclear HP retrofit, new extraction pipes can be welded at site to the outer casing lower half, while supported in a tall frame prior to the outage start. This significantly reduces the number of pipe welds that need to be performed on the critical path during the outage but requires the provision of the support frame. Typical installation durations can vary from 20 days to 35 days. Driven by the large cost implications for delays, in general outage durations are in line with the planned durations and delays are minimal in practice.
14.7
Conclusion
The replacement of steam turbine cylinders has developed from a niche product to an established business sector, as retrofits are an attractive option for owners of mature plants and about one-third of the current installed fleet is over 30 years old. Additional to the deteriorating performance, ageing steam turbines suffer from reliability problems, reducing availability and increasing the risk for the plant owner. Modern retrofits generally exceed the original performance levels and can achieve performance levels equal to current new equipment. The retrofit also
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resets the life clock, extending component life and solving reliability and maintenance issues. The performance potential for a steam turbine retrofit project can be realised in three different principle ways: . . .
higher cylinder efficiency due to technology advancements, last stage blade optimisation (e.g. increased exhaust area) and uprating, i.e. adapting the turbine capacity to increase power.
Typical power gains based on efficiency improvements for retrofits are in the region of 2–3 % of unit output for HPs, up to 0.7% for IPs and 1–2 % for LPs compared to pre-retrofit. Improvements for last stage blade design changes are typically in the region of 0.5–1.5 %, but depend heavily on the cold end conditions. Uprating is usually limited by the retained components or environmental limitations, but large nuclear uprates can yield up to 30 % power increase. To ensure that all balance of plant and hot end implications are considered, close collaboration between vendors and plant owners is recommended. The performance improvements are substantiated by guarantees and their verification. Depending on the scope of supply and the ability to carry out accurate testing, guarantees are selected to ensure an equitable position for the supplier and plant owner. Retrofits have to be designed to fit with the existing components. NonOEM suppliers rely on measurement surveys to define the design boundaries and interfaces, which can be performed during outages in parallel with the maintenance work using conventional tools to three-dimensional laser scans. Based on the spatial constraints from the existing equipment, a design process from the ‘outside in’ is applied to achieve the best performance solution. The installation of a retrofit is a critical discipline as delays are a major commercial risk. It is essential to keep installation effort and time on site as low as possible. Therefore, retrofits typically incorporate flexible design features to compensate for uncertainties at interfaces between the existing and the new equipment. Furthermore, from planning to execution, dedicated field experts are on site to assist the installation activities, providing a direct technical link between the supplier and station. To reduce the typical installation times of 20–35 days (depending on scope) even further, planning systems are continuously improved following best practice. The application of three-dimensional visualisation technologies coupled with planning tools can help to optimise the installation process and is an effective tool to familiarise the work force with the planned activities. With the right planning and management, a retrofit can provide improved performance, reliability and new life with minimal disruption to the planned operation.
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References
Stephen, D., Optimised plant retrofits, in ASME (American Society of Mechanical Engineers). International Joint Power Generation Conference, New Orleans, Louisiana, 4–7 June 2001, ASME, New York. Marlow, B. A., Brown, R. D., Steam turbines, in Hall Stephens J (ed.), Kempes Engineers Year Book for 1998, Miller Freeman, Tonbridge, England, pp. 1248– 1296, 1998. Hogg, S., Stephen, D., Alstom-CPS San Antonio retrofit of JK Spruce unit 1 HP–IP turbine – an example of an advanced steam turbine upgrade for improved performance by a non-OEM supplier, in ASME (American Society of Mechanical Engineers), ASME 2005 Power Conference, Chicago, Illinios, 5–7 April 2005, ASME, New York. Stephen, D., Baldwin, G., Kiensle, C. Optimisation of turbine plant as part of an integrated retrofit project, in ASME (American Society of Mechanical Engineers),ASME 2007 Power Conference, San Antonio, Texas, 17–19 July 2007, ASME, New York. Bartley, P., Foucher, J.C., Hestermann, R., Hilton, B., Keegan, B., Stephen, D., Maximizing economic and environmental performance of existing coal-fired assets, in Alstom Power, Alstom Retrofit Conference, Lisbon, Portugal, 16–17 October 2007, Alstom Power, Baden, Switzerland. Walker, P., Performance improvement and calculation methods of retrofit cylinders, in Alstom Power Turbo-Systems, Steam Turbine Retrofit Conference, San Francisco, California, 16–17 September 2003, Alstom (Switzerland) Ltd, Baden, Switzerland. IEC, IEC 60953, Rules for steam turbine thermal acceptance tests – Part 3: thermal performance verification tests of retrofitted steam turbines, International Electrotechnical Commission, Geneva, Switzerland, 2001. ASME, Steam turbines, performance test codes, ASME PTC 6-2004, ASME, New York, 2004. ASME 2010, Performance monitoring guidelines for power plants, ASME PTC PM-2010, ASME, New York. ASME, Guidance for evaluation of measurement uncertainty in performance tests of steam turbines, ANSI/ASME PTC6 Report – 1985, ASME, New York. Jonas, O., Dupree, L., Gilchrist, T., Lawrence, G., Stevens, M., Copper deposition and MW loss: problem and its solutions, in International Water Conference, Pittsburgh, Pennsylvania, 21–23 October 1996, IWC-96-8. Kearney, P. J., Hogg, S. I., Brown, R. D., Performance guarantee and testing of steam turbine retrofits, in ASME (American Society of Mechanical Engineers), ASME Power, Baltimore, Maryland, 30 March–1 April 2004, POWER200452116, ASME, New York. Jones, M., Nelmes, D. J., The steam turbine retrofitting of six boiling water reactor units using innovative engineering and laser scanning techniques, in ASME (American Society of Mechanical Engineers), ASME Power, Alburquerque, New Mexico, 21–23 July 2009, ASME, New York. Holmes, A., Lord, A., Turbine upgrading by design from survey data, in ASME (American Society of Mechanical Engineers), International Joint Power
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Power plant life management and performance improvement Generation Conference, New Orleans, Louisiana, 4–7 June 2001, ASME, New York. Holmes, A., Retrofit installation, in Alstom Power Turbo-Systems, Steam Turbine Retrofit Conference, San Francisco, California, 16–17 September 2003, Alstom (Switzerland) Ltd, Baden, Switzerland. Jones, C., Turbine island modeling – reducing the risk of retrofit outage schedule overruns, in Power Generation International, Las Vegas, Nevada, 8–10 December 2009.
14.9
Appendix: glossary
AGR ASME CAD CC CMM FAA FGD HBD HP HRSG IP ITB LP LSB NDE OEM PAA RTB SCC TFA VWO
advanced gas cooled reactor American Society of Mechanical Engineers computer aided design combined cycle coordinate measurement machine full-arc admission flue gas desulphurisation heat balance diagram high pressure heat recovery steam generator intermediate pressure impulse blade technology low pressure last stage blade non-destructive evaluation original equipment manufacturer partial arc admission reaction blade technology stress corrosion cracking technical field advisor valves wide open
Published by Woodhead Publishing Limited, 2011
15 High-temperature heat exchangers in indirectly fired combined cycle (IFCC) systems: materials management and performance improvement J . P . H U R L E Y , University of North Dakota Energy & Environmental Research Center, USA
Abstract: In order to increase the efficiency of coal-fired power production, ultrasupercritical pulverized coal-fired boilers, indirectly fired combined cycle (IFCC) systems, and integrated gasification combined cycle systems are being investigated. Each type of system benefits from higher working fluid temperatures, which means higher-temperature materials are required for heat exchangers and turbines. In this chapter, a brief introduction to these technologies is provided, followed by a detailed discussion on the material, fabrication, corrosion, and performance issues for high-temperature heat exchangers (HTHXs) aimed at IFCC systems. In particular, oxide dispersion-strengthened (ODS) alloy-based HTHXs are examined, with data on flowing slag corrosion and pilot-scale performance presented. In addition, a new method for joining ODS alloys is described. Key words: ODS alloys, high-temperature heat exchanger, indirectly fired combined cycle (IFCC), MA956, MA754, evaporative metal bonding.
15.1
Introduction
The Energy Information Administration (2009) predicts that world electricity generation will grow by 77 % between 2006 and 2030, an increase of 14 trillion kilowatt-hours a year. Approximately one-half of that growth will be in coal-fired generation, which will increase as a share of total production from 41 % to 43 %. Yet this growth will occur during a time of increased public concern over the emission of greenhouse gases, particularly CO2. Since coal utilization creates more CO2 per unit of power produced 575 © Woodhead Publishing Limited, 2011
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than any other fuel, the growth in coal-fired generation will be accompanied by increased pressure from the public, corporate boardrooms, and lawmakers to convert the coal as efficiently as possible. Currently, the fleet average coal pile-to-busbar conversion efficiency in the United States is 32.5 %, measured on a higher-heating-value basis, and the average efficiency of the top 10 % of plants is only 37.9 % (Stallard and DiPietro, 2009). In coal-fired power systems the need for greater efficiency will require the use of higher-temperature materials for construction because the quality of the heat produced is more important than the quantity in determining efficiency; that is, the heat in a small volume of gas at a high temperature is more efficiently converted to kinetic energy in a turbine than if the same amount of energy were present in a larger volume of gas at a lower temperature. This means that the heat exchangers used to contain the working fluid, the turbine blades against which the fluid pushes, and the materials from which they are made will ultimately limit the efficiency of a power system. Without employing higher-temperature materials than currently used, the efficiencies of the energy conversion systems cannot grow substantially. Ultrasupercritical (USC) pulverized coal-fired boilers, indirectly fired combined cycle (IFCC) systems, and integrated gasification combined cycle (IGCC) systems are three technologies that enable higherefficiency coal-fired generation. In this chapter, a brief introduction to these technologies is provided followed by a detailed discussion on the material, fabrication, corrosion, and performance issues for high-temperature heat exchangers (HTHXs) aimed at IFCC systems.
15.1.1 Ultrasupercritical systems Worldwide research efforts in advanced pulverized coal (pc) combustion technologies are currently focused on USC and advanced USC (A-USC) steam technology. By increasing pressures and temperatures in traditional pc supercritical (SC) power steam boilers and turbines from 540–585 8C and 24–28 MPa toward A-USC steam conditions of 760 8C and 35 MPa, energy conversion efficiencies over 45 % higher heating value (HHV) can be achieved. Despite extensive steel research, 9–12 wt% chromium martensitic steels appear limited to ~620 8C (Masuyama, 2001) and high-chromium austenitic steels workable to ~675 8C, with limitations for thick-section application because of a high coefficient of thermal expansion combined with low thermal conductivity (Shingledecker and Wright, 2006). Above 675 8C, nickel-based alloys are recommended (Blum and Vanstone, 2003; Viswanathan et al., 2007), although some new austenitic alumina scaleforming dispersion-strengthened steels are under development that may also reach 750 8C, with both creep and oxidation resistance (Yamamoto et al., 2007).
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15.1.2 Indirectly fired combined cycle systems To reach even higher cycle efficiencies than USC systems, a different type of energy cycle is required than the standard Rankine steam cycle. One such type of coal combustion system was researched extensively in the 1990s and early 2000s by the U.S. Department of Energy (DOE). The IFCC power plant was developed under the high-performance power system, or HiPPS, Program (Klara, 1993, 1994). An IFCC plant uses a coal-fired boiler where a fraction of the boiling water heat exchanger pipes, near the hottest part of the flame, are replaced with pipes carrying 1 MPa of air heated as high as 1100 8C. This hot air is used as the working fluid in a gas turbine where its high temperature offsets over two-thirds of the natural gas normally burned in the turbine. Figure 15.1 is a schematic of one such system, which is a modified natural gas-fired combined cycle (NGCC) power plant in which most of the heating of the working fluid in the gas turbine occurs via combustion of less expensive coal rather than more expensive natural gas (United Technologies Research Center, 2001). The waste heat from the gas turbine and the coal combustion system is used to produce steam, possibly USC in some designs, to turn a steam turbine. The Brayton (gas turbine) cycle, which makes one-half of the electricity in the IFCC, is inherently more efficient than a Rankine (steam) cycle, because there is no loss of the heat of vaporization of the water to make steam in the Brayton cycle. Therefore, an IFCC system has the potential to be more efficient than a system based solely on the Rankine cycle (Robson et al., 2002). IFCCs have the added benefit of minimizing water usage at coal-fired power plants by dramatically reducing the amount of cooling and makeup water, since only half as much steam is produced per kilowatt-hour of electricity produced as compared to a typical steam plant. The high efficiency of an IFCC system also makes it suitable for oxygen-blown combustion (oxycombustion) in order to make carbon capture and sequestration more economical. In that case, pure oxygen is added to recirculated flue gas to burn the coal, leaving a gas stream comprising mostly CO2 and steam. After water condensation, only carbon dioxide is left in the gas stream, which can then be used industrially or sequestered, leaving near-zero emissions. If the system is cofired with coal and biomass, sequestration of the carbon dioxide would create a net reduction of its concentration in the atmosphere. Oxygen firing also prevents the formation of thermal NOx. In addition, by staging the combustion of the coal, the volume of flue gas could be dramatically reduced, shrinking the overall size and capital cost of the system. A brief economic analysis was performed by Robson (Hurley et al., 2005) to identify the performance and first-order economics of an IFCC compared to typical pc-fired plants, IGCCs, and NGCCs. Two main IFCC firing scenarios were investigated: coal fired with air and coal fired with O2.
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15.1 The United Technologies Research Centre (UTRC) HiPPS concept; GT stands for gas turbine; ST stands for steam turbine.
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Because IFCC systems have a high potential for operation as zero-emission power plants, special focus was placed on firing with O2 and including CO2 recovery. Two main IFCC energy cycles were investigated: (1) an advanced aeroderivative gas turbine using a coal-fired HTHX with 1150 8C (2100 8F) exit and natural gas boost fuel and (2) a commercially available aeroderivative gas turbine using only a coal-fired HTHX with 1150 8C exit, i.e. all-coal IFCC. When an IFCC was fired with O2 in this analysis, it was assumed that flue gas, consisting primarily of CO2 and H2O, would be recycled to replace a portion of the N2 normally found in the combustion air in order to keep the flame temperature from becoming too high. In all cases, it was also assumed that the amount of O2 supplied would be 20 % greater than that required to completely burn the fuels, but that the amount of recycled flue gas would be less than the amount of N2 excluded so that the total concentration of O2 in the flue gas, before fuel combustion, would be approximately 50 %. By using a reduced flue gas recycle rate relative to the N2 replaced, the overall size of the furnace and pollution control devices can be reduced. All systems used an advanced cryogenic air separation unit (ASU) to supply 98 % O2 for combustion. Also, the exhaust for each, consisting essentially of only H2O and CO2, was cooled, the H2O removed, and the CO2 compressed to 140 bar (2000 psia) for pipeline recovery so that the plant would operate with essentially zero emissions. The levelized cost of electricity (LCOE) is widely used to compare alternative power systems. The Electric Power Research Institute (EPRI) developed the most widely used methodology for determining LCOE (Electric Power Research Institute, 1993). This procedure was the basis of the LCOE presented by Parsons Engineering (2002). These values are used in this analysis to compare the 10-year LCOE of pc-fired plants, IGCC, and natural gas-supplemented and all-coal-fired O2-blown IFCCs with CO2 recovery. The fuel costs were estimated by applying the different system heat rates and prorating the value for the coal and gas fractions. The fuel costs are based on $1.04/MMBtu for coal and $3.25/MMBtu for natural gas. Costs for consumables and fixed and variable operating and maintenance (O&M) were assumed the same for the IFCC plants as for the O2-blown pc plants. The results shown in Fig. 15.2 indicate that the LCOE for an oxygenblown IFCC with CO2 recovery (near-zero emissions), whether fired only on coal or with some supplemental natural gas firing, is approximately 25 % less than for a pc-fired plant in which the CO2 is captured and comparable to the LCOE for an IGCC system. The advantage of the IFCC over an IGCC is that its operation is essentially the same as that of a pc-fired system and, therefore, more likely to be adopted by the existing utility industry. An alternative to CO2 recovery would be to pay a carbon tax. The tax increases the COE by increasing the cost of the fuel. For example, a carbon
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15.2 Calculated LCOE for coal-fired power systems employing CO2 recovery.
15.3 Calculated LCOE for standard pc-fired power systems, including carbon taxes as compared to oxygen-blown IFCC systems with CO2 recovery.
tax of $75/ton carbon emitted ($20/ton CO2) would increase the price of coal by the fraction of coal that is carbon multiplied by $75/ton. For the coal used by Parsons, that would be 0.58 metric tons of carbon/metric ton of as-received coal, which would increase the coal cost by $43/ton, a 170 % increase of the assumed coal price of $25/ton (Parsons Engineering, 2002). If the baseline pc plant were to have a carbon tax, it would raise the LCOE, as shown in Fig. 15.3. The figure also shows that it would take a carbon tax
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somewhat under $50/ton carbon to raise the LCOE of the baseline pc to a level higher than the LCOE for the IFCC with CO2 recovery and to a value of around $75/ton carbon to be higher than the all-coal IFCC with CO2. Given the propensity of the utility industry to use trade-off credits rather than invest in new technology, the introduction of CO2 recovery technology becomes dependent on the carbon tax rate, with a value of approximately $50/ton carbon being the threshold. If rules for CO2 similar to the best available control technology for other pollutants are implemented, then CO2 recovery will be required, and the IFCC appears to be an attractive alternative power system.
15.1.3 Coal gasification systems Research on the IFCC concept largely stopped in the United States in the early 2000s when the DOE decided to focus instead on the FutureGen power system concept (U.S. Department of Energy, 2007). Based on coal gasification, FutureGen plants can achieve efficiencies similar to IFCCs but have the added benefit of being able to separate hydrogen from the syngas during off-peak hours for use in the transportation sector. As shown in Fig. 15.4, the hydrogen concentration in the syngas produced in atmo-
15.4 The effects of temperature and pressure on syngas compositions produced from oxygen-blown coal gasification (from Sondreal et al., 2006).
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spheric-pressure oxygen-blown coal gasifiers is increased at lower temperatures (Sondreal et al., 2006). However, in order to fire the syngas in a turbine, it has to be produced at high pressures. The figure shows that in high-pressure gasification, temperatures of over 9008C are required to produce the highest hydrogen concentrations. To increase them even further, the steam gasification reaction must be emphasized in the gasifier: H2 O þ C?CO þ H2
½15:1
Unfortunately, the steam gasification reaction is endothermic, requiring 56 490 Btu per mole of carbon consumed. To produce the heat required, more carbon can be burned in the gasifier by firing with higher oxygen concentrations, but then more carbon dioxide is produced rather than the preferred carbon monoxide. Another option is to inject superheated steam into the system. In this case, an HTHX carrying the steam can be placed within the combustion zone of the gasifier to heat the steam and then inject it into the reducing zone where the steam and other endothermic reactions take place, reducing the need for added oxygen combustion and increasing the overall hydrogen content of the syngas.
15.2
High-temperature heat exchanger (HTHX) construction
Although advanced coal-fired power technologies have many potential benefits, materials requirements for the HTHX and turbines are stringent. Pressures for the air heaters in IFCC systems are much lower than those for USC systems, so the need for creep resistance is much reduced and austenitic steels and nickel-based alloys are not automatically required. However, material temperatures of as much as 1200 8C in IFCC systems require materials exceptionally resistant to oxidation or corrosion by the products of coal combustion. During the 1980s and 1990s, several programs funded in the United States and United Kingdom focused on the development and testing of ceramicbased heat exchangers for use in coal-fired power systems (Laws and Reed, 1986; Parthasarathy et al., 1992; Solomon et al., 1993), including another HiPPS project headed by Foster Wheeler Development Corporation (Shenker, 1993). Although the nonoxide ceramics maintained their strength at very high temperatures, they were prone to active oxidation (Jacobson, 1993) and severe corrosion by the coal slag when surface temperatures were maintained above the solidus temperature of the slag (Watne et al. 1996). Because of the need to maintain material temperatures below the slag solidus temperature, the difficulty and cost of manufacturing ceramic heat exchangers, and their propensity to thermal shock failure, the UTRC HiPPS
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effort shifted to using high-temperature creep-resistant oxide dispersionstrengthened (ODS) alloys. ODS alloys are excellent candidates for this type of environment (Whittenberger, 1981; Turker and Hughes, 1995; Wright and Stringer, 1997; Hurley et al., 2003b). The presence of small, stable oxides (often yttria [Y2O3]) helps to prevent dislocation motion and preserve the high-temperature strength of these materials. Grain boundary sliding and Herring–Nabarro diffusional creep are both retarded by the elongated grain structure typical of ODS alloys. Two ODS alloys manufactured by Special Metals held the primary focus of the UTRC HiPPS work: MA754, a nickel-based chromia-scale former, and MA956, an FeCrAl alumina scaleforming alloy. Each contain approximately 0.5 % yttria particles by weight. The original HTHX design tested at the pilot scale in the UTRC HiPPS Program used a tubes-in-a-box heat exchanger design, shown in Fig. 15.5. It used MA754 tubes to carry pressurized air, protected from the products of coal combustion by corrosion-resistant ceramic plates. This was the configuration of the heat exchanger during nearly all of the HTHX testing under the UTRC HiPPS Program. However, the ceramic panels were expensive, reduced heat transfer, and were prone to slag corrosion and thermal shock. Subsequent testing of a bare-tube design showed that removing the ceramic panels from the heat exchanger increases the heat transfer rate by nearly a factor of five. This means that the heat exchanger could be as small as one-fifth of that originally proposed. Factoring in the price reduction resulting from the smaller size and lack of ceramic panels, the unprotected heat exchanger panel could cost as little as one-tenth as
15.5 The original UTRC tubes-in-a box HTHX design tested at the EERC at the pilot scale.
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much as the original tubes-in-a-box design. However, corrosion of the alloys then becomes the primary concern. Therefore, in order to determine if the alloys could be exposed directly to the molten slag, laboratory- and benchscale corrosion tests were performed under carefully controlled conditions (Hurley et al., 2003a).
15.2.1 Flowing slag corrosion testing The slag corrosion tests of the MA754 and MA956 alloys were performed in the Energy and Environmental Research Center (EERC) bench-scale dynamic slag application furnace (DSAF), which was modified to permit measuring of the corrosion rates of sections of cooled alloy pipe to flowing slag. Figure 15.6 shows a schematic of the system. During a test, the furnace is heated to 1500 8C while the surface of the alloy tube section is cooled to the desired temperature by an internal airflow. The slag is screw-fed into the furnace system as sand-like grains where it drops into a platinum melting pan. The slag was collected from a commercial slagging boiler system burning Illinois no. 6 coal. The major components in the slag were determined by X-ray fluorescence. Reported as oxide weight percents, the composition is approximately 55 % silica, 19 % alumina, 16 % iron oxide, and 7 % calcia normalized to a closure of 100 %. After melting in the platinum pan, the slag drips on to the surface of the cooled alloy tube section at approximately the 11 o’clock position. It then flows around the cooled tube and drips off the alloy and out of the furnace through a hole in its bottom. After 100 h of slag exposure, the DSAF is cooled, the alloy tube is removed and photographed, and its diameter is measured with a caliper. The surface is then sprayed with dilute epoxy to hold the remaining slag in place, and the tube section is cut into pieces, then embedded in epoxy, crosssectioned, and polished for analysis in a scanning electron microscope (SEM). The DSAF testing provides a worst-case-scenario type of test for the alloys because the corrosion mechanism is more severe than would be encountered by a heat exchanger in a power system. It is more severe because the slag is dripped directly on to the surface of the alloy, whereas in a power system, the surface would be coated with a thin layer of less reactive fly ash before building to the thickness at which the slag would become molten. MA754 and MA956 were tested at 1000, 1100, and 1200 8C. Figure 15.7 shows the surfaces of the alloy tube sections after each test. In all cases, the surface recessions resulting from slag corrosion were not measurable with a caliper because they were so small and, in some cases, slag remained attached to the surface. For both alloys, as the surface temperature increased, the slag wetted the surface more broadly and flowed farther around the circumference before
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15.6
The DSAF system modified for flowing slag corrosion testing of air-cooled alloy tube sections.
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15.7 The surfaces of the MA754 (top row) and MA956 (bottom row) tube sections after flowing slag corrosion testing in the DSAF.
dripping off (the surface coating of slag on the 1000 8C MA956 sample was caused by splashing from the slag-melting pan). However, the surface temperatures of the alloys in all tests were below the solidus temperature of the slag, so immediately adjacent to the alloy the slag was frozen into a very high-viscosity aluminosilicate glass with some intermittent crystallization, primarily of iron oxide-rich species. In all tests, the ribbon of slag frozen to the surface fell off upon cooling. In the case of the MA754, the slag ribbon broke from the alloy within the chromia scale, leaving no slag behind and removing a small amount of chromia in the spalled slag. The chromia content of the frozen slag increased as the slag flowed around the tube. During operation of a power plant, temperature cycling would be expected, so the loss of the chromia with the slag upon cooling would be detrimental to alloy lifetime because the chromia layer provides protection and reduces the oxidation rate of the alloy. Figure 15.8 is an SEM image of a cross section of the surface of the MA754 tube directly below the spot on which the slag dropped after the test at 1200 8C. It shows that, at this particular spot, a 10–15 μm-thick layer was left behind on the alloy, although over most of the sample the slag had completely spalled off. Analyses of composition variations within the alloy and spalled slag show that no slag constituents dissolved into the alloy but that a small amount of the chromia did dissolve into the adjacent slag. In contrast, during cooling of the MA956, the frozen slag broke from the alloy surface within the slag itself, leaving behind a slag layer 20 μm thick on the entire surface of the alloy. Figure 15.9 is an SEM image of a cross section of the surface of the MA956 tube directly below the spot on which
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15.8 SEM image of a cross section of the surface of the MA754 tube directly below the spot on which the slag dropped after the test at 1200 8C.
15.9 SEM image of a cross section of the surface of the MA956 tube directly below the spot on which the slag dropped after the test at 1200 8C.
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the slag dropped after the test at 1200 8C. The figure shows the protective alumina layer which is 5–10 μm thick, with the attached slag layer on top. SEM analyses showed that the slag remaining on the surface was significantly enriched in alumina, indicating that some of the alumina protective layer had dissolved into the slag, but that the slag that spalled off was not enriched. The intergrown fingers of alumina into the remaining slag indicate dissolution and recrystallization of the alumina at the interface of the alumina and slag. However, the fact that some of the slag remained attached to the MA956 surface shows that the protective oxide layer is not lost through spallation during thermal cycling and that up to 1200 8C no measurable alumina is lost into the spalled slag. In addition to imaging the slag–scale interface, the SEM was used with energy-dispersive X-ray analysis to measure the depletion of scale-forming elements as a function of depth within the alloys. The analyses were performed in 10-μm-square areas starting at the top of the scale and moving into the alloy perpendicularly to the surface. Figure 15.10 shows a semilog plot of the concentration of chromium within the MA754 samples as functions of depth for the 1100 and 1200 8C samples directly below where the slag was dropped, as well as for an unexposed sample, and just upstream of where the slag was dropped (no slag). The graph shows that, relative to the unexposed alloy, the samples corroded in the DSAF show a significant depletion of chromium down to a depth of approximately 400 μm below the slag at 1100 8C and
15.10 Chromium content in MA754 as functions of depth for an unexposed sample and two samples exposed in the DSAF.
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15.11 Aluminum content in MA956 as functions of depth for an unexposed sample and three samples exposed in the DSAF.
500 μm at 1200 8C. The amount and depth of depletion are greatest at 1200 8C where there is no slag, possibly indicating that the slag layer may protect the alloy by reducing the oxygen activity at the surface of the alloy. Figure 15.11 shows a semilog plot of the concentration of aluminum within the MA956 samples as functions of depth for the 1000, 1100, and 1200 8C samples directly below where the slag was dropped, as well as for an unexposed sample, and just upstream of where the slag was dropped (no slag). The graph shows that relative to the unexposed alloy, the samples corroded in the DSAF have a thicker alumina scale, but below the scale and in the slag there is no significant depletion of aluminum in the alloy in 100 h. In contrast, for the 1200 8C sample upstream of where the slag was dropped, there is a small but measurable depletion of aluminum down to a depth of 100 μm. As was true for the MA754 samples, this may indicate that the slag actually protects the alloy from oxidation by reducing the oxygen activity at the surface of the alloy. However, because there was no loss of the slag layer from the MA956 during temperature cycling, it should have a longer lifetime than the MA754 in a commercial coal-fired power system.
15.2.2 Joining ODS alloys An excellent summary of experimental work done in joining ODS alloys has been prepared by Oak Ridge National Laboratory (ORNL) (Wright et al.,
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2009), so a review of methods will not be presented here. Instead, we will focus on joining work done in support of, and subsequent to, the UTRC HiPPS Program. Brazed MA754 Figure 15.12 shows a cross section of one of the MA754 tubing elbow joints from the original tubes-in-a-box heat exchanger tested at the pilot scale at the EERC in a slagging combustion system as part of the UTRC HiPPS Program (Hurley and Bornstein, 2004). The pilot-scale system and its operation are described later in this chapter. Under the HiPPS testing, as well as under two separately funded programs, the HTHX was used commonly to produce pressurized air at 950 8C and 150 psig. For a short time, air at 1100 8C and 100 psi was produced. The HTHX was composed of three MA754 tubes, each 1.8 m tall and 6.4 cm outside diameter (o.d.). The tubes were joined to the elbows by UTRC using a nickel-based brazing alloy. The tube and elbow, shown in Fig. 15.12, were removed from the HTHX after 2000 hours of testing in the tubes-in-a-box configuration and 300 hours in a bare-tube configuration. Tests typically lasted 100 hours each. Fuels included natural gas, several coals, and several coal–biomass mixtures. The nickel alloy braze used to join the MA754 tubes to their elbows contained 9.5 % Cr, 3.0 % Al, 4.9 % Ti, 7.0 % (Mo + W), 4.5 % Si, and
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One-half of the cross-sectioned MA754 elbow.
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0.7 % B. The nominal melting point of the braze is about 1040 8C, and the brazing process is usually completed within 36 hours. The joint experienced temperatures ranging from 990 to 1025 8C during a typical test. SEM analyses of the joint between the tube and the elbow showed that the alloy had recrystallized near the joint, but not across the joint. The liquid associated with the braze alloy is most likely responsible for the recrystallization in the area. The braze material did dissolve the MA754 into the joint, but the joint line was very rich in Al and Ti relative to the braze material, possibly through crystallization of an intermetallic within the joint. Bolted MA956 Because of the excellent performance of MA956 in liquid slag corrosion testing described previously in this chapter, it was decided to replace the MA754 tube shown in Fig. 15.12, taken from the bare-tube heat exchanger, with a tube and elbows made from MA956. To join the MA956 tubing to the elbow blocks, the tubes were machined to yield a flat collar, and larger plates of MA956 were bolted on to the elbows, forcing the tube down on to a mica-based O-ring. One tubing elbow joint is shown in Fig. 15.13. The joints were pressure-tested at room temperature to 1500 psig and did not leak. The joints also did not leak during several subsequent 100-hour tests in the pilot-scale combustion system. Results of the performance of the
15.13 Bolted MA956 tubing–elbow joint.
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bare-tube heat exchanger with the MA956 tube replacement while operating the combustion system in an oxy-firing mode are given later in this chapter. Transient liquid-phase bonding of MA956 One technique that had been demonstrated at the EERC for joining ODS alloys is transient liquid-phase (TLP) bonding. This method involves sandwiching a thin layer of low-melting-temperature alloy between the sections to be joined. When the alloy is heated above its melting temperature, it dissolves a portion of the aluminum oxide surface layer as well as the FeCrAl itself and then diffuses into the surrounding ODS sections, leaving a continuous joint. In the TLP work (Hurley, 2009), attempts were made to join pipe sections of MA956 to each other. Joining surfaces were prepared by grinding them flat and polishing with diamond paste, ultimately to a shiny surface with 3μm-diameter diamond particles. The pieces were placed in a clamp made of metals with expansion coefficients similar to or less than that of the MA956, with a layer of joining alloy between them, and held in place with screws. After heat-treating to create the joint, a piece was cut off for SEM analysis of joint cross sections in order to determine the microstructures in the vicinity of the joints. The room-temperature ultimate tensile strength of the joints was determined on the remainder.
15.14 Joint between MA 956 pipe sections showing globules of bonding alloy that had moved into the MA956 on either side of the joint.
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15.15 Joint between MA956 pipe sections made using the heating schedule prepared from DICTRA calculations showing greatly reduced globule formation.
Initial joints made with the MA956 showed that large numbers of small globules of the joining alloy remained near the joint but within the MA956, as shown in Fig. 15.14. After computer modeling of diffusion processes using the computer code DICTRA, a new heating schedule was employed that was successful in causing the bonding alloy to dissolve more fully into the MA956, leaving very few globules of bonding alloy remaining. Figure 15.15 shows an area of the joint that is visible in the unetched samples. For the majority of its length, the unetched joint was not easily visible in backscatter electron imaging. Room-temperature ultimate tensile strengths of the joints reached a maximum of 520 MPa, approximately 80 % of the 650 MPa strength of MA956 reported by the manufacturer. Fracture occurred along recrystallized grain boundaries, as shown in Fig. 15.16, looking down on the break. There was also significant grain pullout, but it occurred in some elongated grains and not in others. SEM analyses of the grain surfaces showed large numbers of small oxide particles remaining along the broken recrystallized grain boundaries, as shown in Fig. 15.17, indicating that the material should retain high-temperature creep resistance.
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15.16 Surface exposed near a joint between MA956 pipe sections showing the tensile break was largely along grain boundaries that had formed through recrystallization near the joint.
15.17 Close-up of the surface exposed from a tensile break showing the distribution of yttrium and aluminum-rich oxide particles along the grain boundaries.
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Evaporative metal bonding In addition to the TLP bonding of MA956, the EERC has been working with Siemens Energy Inc. and DOE on the development of a method for joining FeCrAl ODS alloys to CM247LC, an alumina scale-forming directionally solidified nickel superalloy. The process developed, called evaporative metal (EM) bonding, is similar to TLP bonding. In both, an alloy with a melting temperature below that of the structures to be joined is used to dissolve small quantities of the structures into the joint, and the bonding alloy then diffuses into the large volume of the structure. However, in TLP, the remaining bonding alloy, even if dilute, can adversely affect the properties of the bonded structure. In EM bonding, the joining material similarly dissolves small quantities of base alloy into the joint and then diffuses away but, ultimately, it diffuses through and then evaporates from the surface of the structure and condenses in the cold end of the furnace tube (Cavalli et al., 2008). The EM technique has been used to create joints between CM247LC and dispersion-strengthened Kanthal APMT in which the joint has greater creep rupture life at 950 8C than the APMT. Figure 15.18 shows a creep rupture test specimen after testing at 950 8C with a constant load. The sample was made from a CM247LC rod (left) joined to an APMT rod (right). The joint is at the base of the plastically deformed cone on the CM247LC rod to the left. All of the plastically deformed material is made of APMT. In dozens of
15.18 Creep rupture specimen after testing at 950 8C.
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tests with different loads, joints made with plate or rod, or actual turbine parts, all of the creep rupture breaks occurred within the APMT and not at the joint. The room-temperature ultimate tensile strength of the APMT is not affected by the joining process. However, high-temperature creep rupture data are not yet available for the types of material (rod and plate) or orientations of APMT that we have used in the joints to compare with the creep strength of the joined APMT. Electron microscope X-ray analyses indicate why the joints are so strong relative to the APMT. First, the analyses do not detect bonding metal remaining within the structures, demonstrating that it has almost completely evaporated from the structures since its remaining concentration must be below 0.1 %, the detection limit of the SEM. Second, the EM bonding alloy sufficiently dissolves the oxide scales from the parts being joined to allow significant interdiffusion of the CM247LC and the APMT. Figure 15.19 shows a backscattered electron image of a joint and an X-ray map of the same area. The image shows the precipitates that form within the diffusionaffected zones within each alloy. It is believed that these precipitates help to resist creep near the joint so that, ultimately, the failure occurs in the APMT
15.19 Backscattered electron image (top) and X-ray map (bottom) of an EM joint between CM247LC and APMT.
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rather than at the joint. The X-ray map shows that the depth of iron penetration into the CM247LC, as indicated by the zone containing high levels of precipitates immediately above the joint, is approximately 200 μm. Nickel penetration into the APMT, as indicated by the dark precipitates in the APMT near the joint, is obvious to approximately 300 μm but continues in trace amounts to approximately 700 μm.
15.3
Pilot-scale HTHX testing
15.3.1 Design of the EERC slagging furnace system The EERC tested the HTHX in the slagging furnace system (SFS), which was designed, constructed, and operated under the UTRC HiPPS Program and several subsequent programs. The SFS is designed to heat the HTHX under flowing slag conditions so that it can produce process air at 950 8C and 150 psig. A schematic of the system is shown in Fig. 15.20. The SFS is designed for a maximum furnace exit gas temperature of 1600 8C but is typically run at 1500 8C in order to maintain desired slag flow while extending refractory lifetime. It has a nominal firing rate of 2.5 million Btu/h and a range of 2.0–3.0 million Btu/h using a single burner. The EERC oriented the furnace vertically (downfired) and based the burner design on a swirl burner currently used on two EERC pilot-scale pc-fired units that are fired at 600 000 Btu/h. The furnace dimensions are 1.2 m inside diameter (i.d.) by roughly 5.5 m in length. It is lined with three layers of refractory totaling 0.3 m in thickness. The inner layer is composed of an alumina castable, developed by the Plibrico Company in conjunction with the EERC, which has been shown in laboratory- and pilot-scale tests to be extremely resistant to slag corrosion. A key design feature of the furnace is accessibility for installation and testing of a large HTHX test panel for testing material lifetimes and heat exchange coefficients. The HTHX is 0.361.8 m in size. This size was based on manufacturing constraints identified by UTRC, which designed and built it. It is shown in Fig. 15.21 in the bare-tube configuration. The HTHX is composed of three vertically oriented 6.4-cm-diameter ODS alloy tubes. Originally constructed with MA754, a nickel–chromium alloy, the center tube was replaced prior to starting bare-tube configuration testing with an MA956 tube, as described in section 15.2.2. After the MA956 tube was installed, the lower half of the HTHX was encased in new insulating material and high-density refractory to limit the surface area exposed to furnace conditions. In addition, the upper 11 % was encased in fibrous insulation and an insulating board to protect bolting materials used in the mechanical seal employed on both ends of the MA956 tube. As a result, only 39 % of the HTHX alloy tube surface was exposed to furnace conditions
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15.20 Schematic of the EERC SFS used to test the HTHX.
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15.21 The bare-tube HTHX in place in the SFS.
during bare-tube configuration tests. Air is heated from 540 8C to as much as 980 8C as it passes through. A tie-in to an existing nitrogen system was also installed as a backup to the existing air compressor system to prevent the panel from overheating in the event of a power outage. As the hot combustion gases leave the combustor, they pass through a slag screen to remove the entrained ash as a nonleachable slag and reduce deposition on the convective air heater (CAH). It removes approximately 65 % of the particulate matter from the gas stream. As the hot combustion gas leaves the slag screen, it is quenched with recirculated flue gas to 1010 8C in order to make the ash less sticky and reduce deposition on the CAH. The quench zone is the only region in the furnace where hard ash deposits form, but they are easily removed by knocking them into a hopper at the bottom of the zone. The gases then pass over the CAH, which is used to heat the process air to 540 8C. The hot combustion gases then flow through a series of heat exchangers and, finally, a baghouse on the way to the system stack.
15.3.2 HTHX testing during oxy-firing The HTHX was tested in the EERC SFS in the tubes-in-a-box configuration for over 2000 hours while burning natural gas or one of several different coals. It was tested for over 500 hours in the bare-tube configuration while burning a variety of coals and coal–biomass blends. In one test, the EERC operated the SFS in an oxy-fired mode. That test is described in detail here.
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Replacing the primary airstream for O2 blowing involved recirculating flue gas through a quench/condenser to remove as much water vapor as possible, reheating the gas stream to avoid moisture problems associated with fuel feed, and using a booster fan to provide an adequate driving force. The primary gas stream was heated to > 10 8C above saturation temperature to avoid moisture condensation and the potential for coal feed problems related to condensed moisture. Oxygen was added to the secondary gas stream only and limited to 35 vol% (wet basis) or less. Since the primary gas stream conveying fuel to the burner and supporting burner operation was composed of recycled flue gas, it contained <4 vol% O2 on a wet basis or 5–10 vol% on a dry basis. Other modifications to permit oxygen-blown operation included sulfur dioxide control to avoid sulfur dioxide accumulation in the recirculated flue gas as well as temperature control to avoid moisture condensation from the recycled flue gas. HTHX process airflow rates during O2-blown natural gas firing were between 5.7 and 11.2 scm/min. HTHX process air exit temperatures were 750–860 8C. Inlet process air temperatures ranged from 560 to 615 8C. Figure 15.22 summarizes the heat recovery data as a function of furnace temperature for air-blown tubes-in-a-box HTHX tests, air-blown bare-tube HTHX tests, and the oxygen-enriched bare-tube HTHX test. The data are presented on an equivalent-surface-area basis. Heat recovery data for baretube air-blown tests are represented by closed circles (natural gas-fired) and closed triangles (coal-fired). Data resulting from the oxygen-enriched tests
15.22 Summary of the HTHX heat recovery data as a function of furnace temperature for air-blown tubes-in-a-box HTHX tests, air-blown bare-tube HTHX tests, and the oxygen-enriched bare-tube HTHX test.
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are represented by open circles (natural gas-fired) and open triangles (coalfired). The open (coal-fired) and closed (natural gas-fired) squares in the lower right represent heat recovery data from work completed in 1999 and 2000 with the tubes-in-a-box configuration. As previously described, with that design, the HTHX heat recovery did not exceed 120 000 Btu/h (126 600 kJ/h) for furnace temperatures in the range of 1540–1580 8C. However, removing the ceramic panels from the heat exchanger increases the heat recovery rate by nearly a factor of five during air-blown operation. This means that the HTHX could be as small as one-fifth of that originally proposed by UTRC under the HiPPS Program. Factoring in the cost savings by not requiring the ceramic panels, the bare-tube heat exchanger could cost as little as one-tenth as much as the original UTRC design. The measured heat recovered from the HTHX during O2-blown natural gas firing ranged from nominally 118 500 Btu/h to 192 000 Btu/h, respectively. Corresponding furnace firing rates and exit temperatures were 2.9–3.3 MMBtu/h and 1290–1400 8C, respectively. HTHX tube surface temperatures ranged from nominally 560 to 1180 8C. Isolating 61 % of the HTHX from furnace conditions permitted the slagging furnace to be fired at rates sufficient to achieve furnace temperatures comparable to previous work and avoid overheating the HTHX alloy tubes. However, slag screen exit temperatures were higher than the furnace temperatures during the O2-blown test series. Therefore, optimization of slagging furnace operation in an O2-blown configuration will be necessary before it will be possible to document the heat-transfer advantages of employing the HTHX specifically in an O2-blown combustion system. Heat recovery for the natural gas-fired O2-blown test was greater than the air-blown test by 42–68 %, on a comparable surface-area basis at furnace temperatures up to 1400 8C and for similar process airflow rates. These data indicate that at comparable furnace temperatures, a > 50 % increase in radiant heat transfer to the HTHX may occur as a result of operation of the slagging furnace in an O2-blown configuration. Optimization of O2-blown firing in the slagging furnace would likely result in an additional increase in radiant heat transfer to the HTHX because of the high radiant flux from the CO2/H2O-rich flame. Proving this hypothesis will only be possible when sufficient process air can be made available to protect alloy tube surfaces from overheating with the entire HTHX exposed to furnace conditions while operating in air-blown and O2blown configurations at comparable furnace-firing rates and temperatures and in an optimized O2-blown configuration. A significant improvement in furnace flame characteristics and resulting furnace temperature distribution and operating flexibility could be achieved by upgrading the O2 delivery system to permit a 21 vol% (wet basis) O2
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concentration in the primary gas stream supporting the main burner. Based on the combined natural gas- and coal-fired data set, the O2-blown operation at a furnace temperature of 1540–1565 8C could result in an HTHX heat recovery rate of 600 000 Btu/h and possibly greater. This indicates that the O2-blown heat recovery rate could be 50 % greater than the air-blown heat recovery rate in the radiant zone at a given furnace temperature. In addition, by reducing the amount of flue gas recirculation, it was demonstrated that flue gas flow rates could be substantially reduced relative to air-blown operation, conceivably reducing the size of the furnace and pollution control devices. Also, after scrubbing and particulate cleanup, the final flue gas consists almost completely of carbon dioxide and water vapor. Therefore, it would be possible to condense most of the water for reuse in the plant and harvest the remaining carbon dioxide for industrial use or geological sequestration so that the plant could operate essentially emission-free.
15.4
Conclusions
In order to reach higher energy conversion efficiencies, advanced coal-fired power plant concepts such as USC, IGCC, and IFCC are under development. In the USC and IFCC concepts in particular, heat exchangers operating at higher temperatures than those in standard pc-fired plants are being tested. The main focus of EERC work has been on the HTHX concept, which could be used in an IFCC to create relatively low-pressure (1 MPa) but very high-temperature (1100 8C) air to turn an aeroderivative turbine. The IFCC HTHX concept has been demonstrated to be very resistant to corrosion by coal slag if made from alumina scale-forming ODS alloys and operated with a surface temperature below the solidus temperature of the slag. Because of the high metal surface temperatures, the slag layer, which inevitably forms on the tubes, is very thin and is much less of an impediment to heat transfer than sintered deposits that form on USC heat exchangers. IFCCs have the added benefit of minimizing water usage at coal-fired power plants by dramatically reducing the amount of cooling and makeup water, since only half as much steam is needed per kilowatt-hour of electricity produced as compared to a purely Rankine cycle-based plant. The high efficiency of an IFCC system also makes it suitable for oxygen-blown combustion (oxycombustion) in order to make carbon sequestration more economical. Pilot-scale oxycombustion tests have also shown that radiant heat transfer from the flame may be increased because of its high CO2 and H2O contents. One problem with ODS alloy-based heat exchangers is the difficulty in joining the components, since ODS alloys lose a large fraction of their creep resistance if joined by fusion welding. The EERC has demonstrated a
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mechanical joining method and tested it for several cycles in a pilot-scale system. A more robust solution, employing a version of TLP bonding known as EM bonding, has also been demonstrated for joining ODS alloys, although tubes joined by this method need to be tested on the pilot scale.
15.5
Acknowledgments
Funding for this work was provided by the DOE Advanced Research Materials Program, DOE–UTRC HiPPS Program, DOE–EERC Fossil Energy and National Center for Hydrogen Technology Cooperative Agreements, the North Dakota Industrial Commission, Xcel Energy, and Siemens Energy Inc.
15.6
References
Blum, R., Vanstone, R. W. (2003), Materials development for boilers and steam turbines operating at 700 8C, in Parsons 2003 – Proceedings of the Sixth International Charles Parsons Turbine Conference, Institute of Materials, Minerals, and Mining, London, 2003, pp. 489–510. Cavalli, M. N., McNally D. P., Hurley, J. P., Norman Bornstein, N. S. (2008), Joining of difficult-to-weld materials. U.S. Patent Application 2009/0250442 A1, filed December 3, 2008. Electric Power Research Institute (1993), TAG technical assessment guide – electricity supply – 1993. EPRI TR-102276. Energy Information Administration (2009), International energy outlook 2009 [online]. Available at www.eia.doe.gov/oiaf/ieo/index.html [accessed May 2010]. Hurley, J. P. (2009), Joining of ODS alloys. Final report from Task 8: Oxide Dispersion Strengthened Alloys. Phase 1 effort under the UK–US Advanced Materials Program, UK–US Collaboration on Fossil Energy Research and Development 2004–2009, published April 23, 2009. Hurley, J. P., Bornstein, N. S. (2004), Testing of a very high-temperature heat exchanger for IFCC power systems. Presented at the 18th Annual Conference on Fossil Energy Materials, Knoxville, Tennessee, June 2–4, 2004. Hurley, J. P., Kay, J. P., Williams, K. D., Bornstein, N. S. (2003a), Measurements of the corrosion of two ODS alloys by flowing coal slag. Proceedings of HighTemperature Corrosion and Materials Chemistry IV, Electrochemical Society, Paris, April 2003. Hurley, J. P., Weber, G. F., Jones, M. L. (2003b), Tests of high-temperature heat exchangers for indirectly fired combined cycles. In Proceedings of the 20th Annual International Pittsburgh Coal Conference, Pittsburgh, Pennsylvania, 2003. Hurley, J. P., Weber, G. F., Robson, F. (2005), Pilot-scale test and cycle analyses of an oxygen-blown IFCC power system. In Proceedings of the 22nd International Pittsburgh Coal Conference, Pittsburgh, Pennsylvania, September 2005, Session 25-3.
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Jacobson, N. S. (1993), Corrosion of silicon-based ceramics in combustion environments. Journal of the American Ceramics Society, 76(1), 3–28. Klara, J. M. (1993), HiPPS: beyond state-of-the-art: Part I. Power Engineering, December, 37–39. Klara, J. M. (1994), HiPPS can compete with conventional PC systems: Part II. Power Engineering, January, 33–36. Laws, W. R., Reed, G. R. (1986), Ceramic heat exchangers for coal-fired and dirty gas applications. In High-temperature heat exchangers. Hemisphere Publishing, pp. 550–567. Masuyama, F. (2001) History of power plant and progress in heat resistant steels. ISIJ International, 41[6], 612–625. Parsons Engineering (2002), Advanced fossil power system comparison study. Report to U.S. Department of Energy. Parthasarathy, V., Harkins, B., Beyermann, W., Keiser, J., Elliot, W., Ferber, M. (1992), Evaluation of SiC/SiC composites for heat exchanger applications. Ceramic Engineering and Science Proceedings, 13[7–8], 503–519. Robson, F. L., Ruby, J. D., Nawaz, M., Seery, D. J., Jones, M. L., Hurley, J. P. (2002), Application of high-performance power systems (HiPPS) in Vision 21 power plants. In Proceedings of the 2002 American Power Conference, Chicago, Illinois, 2002. Shenker, J. (1993), Development of a high-performance coal-fired power-generating system with a pyrolysis gas and char-fired high-temperature furnace. In Proceedings of the Ninth Annual Coal Preparation, Utilization, and Environmental Control Contractors Conference, U.S. Department of Energy, Pittsburgh, Pennsylvania, July 1993, pp. 349–355. Shingledecker, J. P., Wright, I. G. (2006), Evaluation of the materials technology required for a 760 8C power steam boiler. In Proceedings to the 8th Liege Conference on Materials for Advanced Power Engineering 2006. Forschungszentrum Ju¨lich GmbH, pp. 107–120. Solomon, P. R., Serio, M. A., Cosgrove, J. E., Pines, D. S., Zhao, Y., Buggein, R. C., Shamroth, S. J. (1993), Feasibility study for an advanced coal-fired heat exchanger/gas turbine topping cycle for a high-efficiency power plant. In Proceedings of the Ninth Annual Coal Preparation, Utilization, and Environmental Control Contractors Conference, U.S. Department of Energy, Pittsburgh, Pennsylvania, July 1993, pp. 210–225. Sondreal, E. A., Swanson, M. L., Benson, S. A., Holmes, M. J., Jensen, M. D. (2006), A review of gasification technology for coproduction of power, synfuels, and hydrogen from low-rank coals. In Proceedings of the Symposium on Western Fuels: 20th International Conference on Lignite, Brown, and Subbituminous Coals, Denver, Colorado, Oct 24–26, 2006. Stallard, S., DiPietro, P. (2009), An opportunity to improve coal-fired generation efficiency. Power Engineering, November, 122–125. Turker, M., Hughes, T. A. (1995), Oxidation behavior of three commercial ODS alloys at 1200 8C. Oxidation of Metals, 44, 505–525. United Technologies Research Center (2001), Combustion 2000 Phase II final technical report, DE-AC22-95PC95144. U.S. Department of Energy (2007), Fossil energy [online] (July 14, 2009). Available
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at www.fossil.energy.gov/programs/powersystems/futuregen [accessed May 2010]. Viswanathan, R., Coleman, K., Shingledecker, J., Sarver, J., Stanko, G., Borden, M., Mohn, W., Goodstine, S., Perrin, I. (2007), Boiler materials for ultrasupercritical coal power plants. Final Summary Report for U.S. Department of Energy DE-FG26-01NT41175, Energy Industries of Ohio. Watne, T. M., Hurley, J. P., Gunderson, J. R. (1996), Pilot-scale combustion testing of silicon-based ceramics. In Proceedings of the 20th Annual Cocoa Beach Conference on Composites, Advanced Ceramics, Materials, and Structures, Cocoa Beach, Florida, January 7–11, 1996. Whittenberger, J. D. (1981), Elevated temperature mechanical properties of the iron base oxide dispersion strengthened MA956 bar. Metallurgical Transactions A, 12A, 845–851. Wright, I. G., Stringer, J. (1997), Materials issues for high-temperature components in indirectly fired cycles. In Proceedings of the International Gas Turbine and Aeroengine Congress and Exhibition, Orlando, Florida, 1997. Wright, E. G., Tatlock, G. J., Al-Badairy, H., Chen, C. L. (2009), Summary of prior work on joining of oxide dispersion-strengthened alloys. Oak Ridge National Laboratory, ORNL/TM-2009/138. Yamamoto, Y., Brady, M. P., Lu, Z. P., Maziasz, P. J., Liu, C. T., Pint, B. A., More, K. L., Meyer, H. M., Payzant, E. A. (2007), Creep-resistant, Al2O3-forming austenitic stainless steels. Science, 316, 433–436.
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16 Heat recovery steam generators: performance management and improvement V . G A N A P A T H Y , Boiler & HRSG Consultant, India
Abstract: This chapter discusses the general features of fire and water tube waste heat boilers used in chemical plants, refineries and in cogeneration plants including gas turbine HRSGs. One of the major issues with this heat recovery system is optimization of gas and steam temperature profiles and maximization of energy recovery. The chapter discusses a simple simulation technique by which any plant engineer who is unfamiliar with the design aspects of HRSGs can simulate the thermal performance of unfired and fired HRSGs at various operating conditions of the gas turbine and arrive at the optimum steam system parameters and the number of steam pressure levels with an understanding of only pinch and approach points. Once the heat balance is achieved then one can contact the HRSG supplier for the actual design and supply of the equipment. Key words: boiler, water tube, fire tube, HRSG, single pressure, multiple pressure, efficiency, unfired, fired, circulation, natural, forced, once through, circulation ratio, pinch point, approach point, HRSG simulation, economizer steaming, evaporator, water chemistry, off-design performance, superheater, economizer, emissions.
16.1
Introduction
A heat recovery steam generator (HRSG) is a term generally used in the power and process industry for a waste heat recovery system generating steam (or hot water or hot industrial heat transfer fluids) from hot exhaust gases from diesel engines, gas turbines, incineration systems, effluents from chemical plants and refineries, kilns and furnaces or any hot gas generator. However, for reasons unknown, the term is widely used today in connection with waste heat boilers behind gas turbines. This chapter will discuss waste heat boilers in general and gas turbine exhaust HRSGs in particular. A 606 © Woodhead Publishing Limited, 2011
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notion that existed in earlier years was that waste heat boilers were generally unfired and of water tube type design. However, they can be unfired or fired and they can be of fire tube or water tube type. Fire tube and water tube boilers can also be combined to form a single steam system with a common steam drum and external down comer, riser piping. There are waste heat boilers of hybrid type recovering energy from incinerator exhaust gases and in addition are separately fired with a fuel such as natural gas or oil to augment steam generation. These boilers can generate steam even when the waste gas stream is not available, like a fired steam generator. Thus there are several options for recovering energy from a hot gas source. Waste heat boilers can be classified based on the purpose for which they are used, the nature of the flue gas stream, circulation type, whether fired or unfired, steam pressure levels and configuration, whether fire tube or water tube, as shown in Fig. 16.1. As can be seen, there can be a variety of configurations for waste heat boilers and the type of design is only limited by the imagination and experience of the boiler designer. Waste heat boilers are used in chemical plants such as sulfuric/nitric acid plants and hydrogen plants to cool the waste gas stream from a given inlet gas temperature to a desired exit gas temperature for process purposes, as the flue gas stream may be taken to a reactor or catalyst for further processing. In these applications, energy recovery is not the primary concern. Such boilers are equipped with gas bypass systems for controlling
16.1 Waste heat boiler classification.
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Table 16.1
Typical waste gas data (% volume of gas)
Gas Temperature (8C) P (atm) N2 1 2 3 4 5 6 1. 2. 3. 4. 5. 6.
1000 500 1100 1100 550 900
1 1 1 50 1 1
O2 SO2 SO3 CO2 H2O CO CH4 H2
80 10 10 81 11 1 70 3 0.5 75 15 73 2
7 9 6 3 5
18 37 7 20
8
5.5
43
Raw sulfur gases from sulfur combustor in sulfuric acid plant. Sulfur gases after converter in sulfuric acid plant. Reformer flue gases in hydrogen plant. Reformed gas in hydrogen plant. Gas turbine exhaust. Fume incineration exhaust.
the exit gas temperature at low loads, while in applications such as gas turbine exhaust or fume incineration systems, the primary concern is to recover the maximum amount of energy from the exhaust gas stream compatible with considerations of cost and low temperature corrosion. Low pinch and approach points and/or multiple pressure level steam generation and the use of extended surfaces are features of these water tube boilers. Gas turbine HRSGs can be of single or multiple pressure, adding to the complexity of its performance evaluation. Multiple pressure units are required to recover the maximum amount of energy from the turbine exhaust gases. Reasons for this will be discussed later when we discuss gas turbine HRSGs. Another aspect to be considered in the design of waste heat boilers is the analysis of the flue gas stream, its nature and ash content if any. Table 16.1 shows the typical flue gas analysis from a few common applications: corrosive gases such as hydrogen, hydrogen chloride or sulfur containing gases can cause high and low temperature corrosion problems; some gases such as hydrogen and water vapor are responsible for high heat transfer coefficients and high heat flux due to their high thermal conductivity and specific heat, and hence care should be taken selecting materials for such applications; ash if present in the flue gases, as in municipal waste incineration systems, can cause slagging and fouling problems, which will affect the boiler design, performance and cost significantly. Water treatment and steam purity are also important issues and will affect the fouling on the steam side and the tube wall temperatures. Waste heat boilers can be of fire tube or water tube type. In fire tube boilers, the flue gas stream flows inside the tubes while steam is generated on the shell side. In water tube boilers, it is the opposite.
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Effect of steam pressure on tube thickness
Tube thickness (mm)
External pressure (kg/cm2g)
Internal pressure (psig)
2.7 3.03 3.43 3.81 4.6
40.4 48.2 56.2 64.75 82.4
80.6 94.1 107.8 121.6 150
(Based on sa 178A or sa 192 carbon steel tubes at 371 8C temperature.)
Features of water tube boilers are: . . . . . . .
They are suitable for high steam pressures and temperatures and large gas flows. Provision can be made for cleaning within the gas stream. A superheater can be located anywhere inside the gas path. It is expensive if the gas pressure is high. Heating surfaces may have to be located within a cylindrical shell or pressure vessel. The typical gas pressure handled is 300 to 400 mm wc (water column). Multiple pressure steam generation is feasible. There is a lower water holdup inside the drum and tubes compared to fire tube boilers and hence the response to gas flow variations is faster. Extended surfaces may be used (if the gas is clean) to make the HRSG compact.
Features of fire tube boilers are: . . . . . .
They can handle high gas pressures up to 15 MPa with ease as flue gas flows inside the tubes. Due to smaller heat transfer coefficients and the absence of extended surfaces, they are large in size. The location of a superheater is either at the boiler inlet or outlet, not at an optimum location as in water tube boilers. An economizer and superheater can be added if required. They are limited to low steam pressures and capacities typically less than 50 t/h gas flow. They have a sluggish response to load changes due to a large water holdup in the shell.
Table 16.2 shows the effect of steam pressure on tube thickness per the ASME code based on whether it is applied inside or outside the tubes. It can be seen that a given tube can withstand a higher pressure if it comes from inside the tubes as in a water tube boiler. Hence water tube boilers are ideal for handling high steam pressures and large gas flows. Figure 16.2 shows a water tube boiler behind a fume incinerator. The gas
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16.2
Waste heat boiler for a fume incinerator.
stream is clean and hence finned tubes are used where applicable to make the HRSG compact. The HRSG consists of a screen section, followed by a two stage superheater with interstage attemperation and an evaporator and economizer. A typical flue gas inlet temperature is 800 to 900 8C and hence the screen section and superheater are generally of a bare tube construction. The evaporator, which operates in parallel with the screen section, consists of extended surfaces of varying fin densities in each row along the gas flow direction depending on the tube wall temperature and fin tip temperature. The enclosure can have a membrane wall or be refractory lined. If the gas flow is large, the steam drum may be mounted externally and both the evaporator and screen are connected by external downcomers and risers. The screen section shields the superheater from hot gases and external radiation and is recommended in such applications. If NOx reduction is also a factor, a suitable NOx catalyst may be added at the appropriate gas temperature region. This may require splitting the evaporator. In a few incineration projects, some saturated steam may be imported into the superheater and the steam from the HRSG as well as the import steam may be raised to the desired steam temperature. The designer should also check the superheater design and materials for cases when the import steam is absent. Several gas inlet conditions may occur in such projects and the HRSG performance should be evaluated in all cases, with and without the import steam. Figure 16.3 shows a fire tube boiler behind a hydrogen reformer handling high pressure reformed gases containing hydrogen and, methane, as shown in Table 16.1 above. The steam drum is mounted externally and connected to the boiler using downcomers and riser pipes. There is an internal gas bypass system to control the exit gas temperature at part loads of the boiler. Since the gases are corrosive due to the presence of hydrogen and methane,
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16.3 Fire tube boiler for a hydrogen reformer with an internal gas bypass system.
alloy steel tubes such as T11 or T22 are used for the boiler tubes. The gas inlet temperature is high, about 1050 8C, and hence the heat flux can be quite high inside the tubes and at the tube sheet. Hence refractory and ferrules are used to protect the tube sheet. The inlet and exit vestibules are lined with refractory to protect them as well as to minimize the heat loss. Heat flux at the steam side of the boiler tubes is limited to about 315 kW/m2. With this brief introduction on waste heat boilers and their types and features, let us proceed to gas turbine HRSG design and performance aspects.
16.2
Gas turbine heat recovery steam generators (HRSGs)
Gas turbine power output varies from a few MW to hundreds of MW and hence the exhaust gas flow can also vary from, say, 20 000 kg/h to 2 million kg/h. HRSG sizes can be large. Generally they are shipped in modules such as superheaters, evaporators and economizers, while small HRSGs can be completely assembled in the shop, as shown in Fig. 16.4. In large HRSGs, the modules are shipped separately and assembled in the field. NOx and CO catalysts may be added as required. Figure 16.5 shows a large multiple pressure natural circulation HRSG. In multi-pressure HRSGs, steam is generated at different pressure levels in order to lower the exhaust gas temperature and recover more energy from the turbine exhaust gas and thus to improve the HRSG efficiency. HRSG simulation studies as shown in the following pages can give an idea of how more energy can be recovered using multi-pressure steam generation. Since the emphasis of this chapter is on performance and optimization, constructional aspects will be mentioned only in passing. As discussed earlier, HRSGs can be of natural or forced circulation type. Figure 16.6
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16.4
Two pass HRSG for a small gas turbine (up to 15 MW).
16.5 Multiple pressure fired HRSG for a large gas turbine (above 15 MW).
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16.6 Forced circulation HRSG.
shows a forced circulation unit in which a circulation pump is used for circulating the steam water mixture through the tubes. In forced circulation HRSGs, the boiler water from the steam drum is taken and forced through horizontal evaporator tubes to generate wet steam and is then recirculated to the steam drum – hence the name. Forced circulation designs are common in Europe while in the USA natural circulation designs are the norm. The forced circulation units have an operating cost associated with the circulating pump; allowable heat fluxes inside the tubes are lower due to the horizontal tube design compared to the vertical tube designs in natural circulation units. This issue assumes significance in the case of fired HRSGs. However, they occupy less real estate compared to natural circulation units. Cleaning of selective catalytic system modules may be difficult with forced circulation units as the modules are stacked one below the other, while in natural circulation designs they are side by side. To minimize the circulating pump power consumption, a low circulation ratio (CR) of 3 to 5 is used, depending on the manufacturer’s practice. The circulation ratio is the ratio of total mixture flow to the steam generation. Thus, if 100 000 kg/h of steam is generated in the HRSG, a CR of 5 would mean that 500 000 kg/h of steam–water mixture is flowing
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through the evaporator tubes. The water is separated from the mixture in the steam drum using steam purifiers and dry steam is sent to the superheater. The remaining water is mixed with feedwater entering the drum and sent through the circulating pumps. Once through HRSGs are also available. In these designs, water is converted to saturated or superheated steam and hence phase separation occurs inside the tubes. While these dispense with the steam drum and enable a slightly quicker startup, the steam side pressure drop due to the two phase flow transition is rather high, on the order of 15–20 % of the exit steam pressure. This is a significant operating cost and has to be reckoned with by the end user. In addition, very pure feedwater with zero solids is required to avoid deposition of solids inside the evaporator tubes during the phase change. If deposited in the tubes, solids can increase the tube wall temperatures and lead to tube failures. To overcome this concern some suppliers use inconel tubes, which can withstand high tube wall temperatures and can also run dry, but add to the cost. As mentioned earlier, HRSGs can be of unfired or fired type depending on plant needs. In process or cogeneration plants fired HRSGs are common as steam demands can vary wildly, while in power plants the HRSG designs are rather standardized. Recent HRSG designs for combined cycle applications include reheaters so that the combined cycle efficiency is increased by a few percentage points. However, this adds to the cost and complexity of the HRSG as a turbine steam bypass may have to be provided to protect the reheater during startups. From the data on exhaust gas, it is seen that the exhaust gas temperature from the gas turbine is quite low, on the order of 450 to 6008C. However, it has sufficient oxygen in the gas stream to be supplementary fired without the addition of combustion air. The supplementary fired HRSGs, in which the firing temperature lies in the range of about 650 to 8508C look similar to the unfired units except for the insulation thickness, larger drums and interconnecting piping to handle the higher steam generation. A typical exhaust gas temperature from the gas turbine varies from 450 to 525 8C and contains sufficient oxygen for combustion of fuels such as natural gas. By adding the fuel in burners, the exhaust gas temperature is raised to as high as 8508C in supplementary fired HRSGs. This temperature after the burner is called the firing temperature. The tube materials have to be selected while considering the higher tube wall temperatures due to the higher gas temperatures and heat flux inside the tubes. The steam generation is varied by varying the firing temperature while the exhaust gas flow is typically unchanged. The furnace fired units, on the other hand, are fired to a much higher temperature, approaching 1650 8C which naturally requires a water cooled membrane walled furnace to handle the combustion process. Figure 16.7 shows a furnace fired unit in which the combustion process occurs inside a water cooled furnace. The superheater is
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Comparison of unfired and fired HRSGs
Item
Unfired
Supplementary fired
Furnace fired
Gas temperature (8C) Gas/steam ratio Burner type Fuel Furnace casing
450–600
650–850
850–1600
5.5–7 No No Insulated with 100 mm in ceramic fiber Natural/forced/ once through Single/multiple steam pressure 100–200
2.5–5.5 Duct burner Oil/gas 200–250 mm ceramic fiber or membrane wall Natural/forced/ once through Single/multiple pressure 150–300
1.2–2.5 Duct/register Oil/gas/solid Membrane wall
Convective finned tubes
Convective, finned Furnace, bare tubes tubes mostly
Circulation Configuration Back pressure (mm wc) Design
Natural/once through Single pressure > 300
Note: gas temperature refers to gas temperature entering the HRSG after firing, if any.
16.7 Furnace fired gas turbine HRSG.
shielded from the furnace flame radiation by screen tubes. The convection section follows, which can have a combination of bare and finned tubes. Note that HRSGs can also be fired with combustion air from a separate fan when the gas turbine is not running in order to generate steam. However, this is an inefficient process and should be resorted to as a backup measure only. Furnace fired and supplementary fired units are very efficient, as will be discussed later, and any additional energy input to the burner is 100 % utilized, while efficiency of oil/gas fired steam generators is about 90– 92 %. A few salient aspects of unfired and fired HRSGs are shown in Table 16.3.
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16.2.1 Understanding HRSG performance In order to understand the performance aspects of HRSGs, the gas/steam temperature profiles must be understood. Due to the low inlet gas temperature of HRSGs, the steam generation is very much dependent on the steam pressure. Two important variables, namely the pinch point and the approach point, come into the picture while evaluating the gas/steam temperature profiles. One does not concern oneself with pinch and approach points in typical oil/gas/solid fuel fired boilers or steam generators as the inlet gas temperature is the combustion temperature of the fuel and high enough to ensure that the temperature cross situation does not arise, whereas with gas turbine HRSGs the pinch and approach temperatures determine the gas/steam temperature profiles and duty of the HRSG. The large ratio of gas flow to steam generation also affects the gas/steam temperature profile and particularly affects the economizer performance at low loads, a problem called ‘steaming’ in the economizer. This is discussed later. Figure 16.8 shows the typical gas/steam temperature profile for an HRSG. Pinch and approach points are terms associated with the evaporator. The pinch point is the difference between the gas temperature leaving the evaporator and the steam saturation temperature (tg3ts), while the approach point is the difference between the saturation temperature and the water temperature entering the evaporator or leaving the economizer
16.8 Pinch and approach points, where pinch = (tg3–ts) and approach = (ts–tw2).
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(tstw2). While arriving at the design or thermal performance of the HRSG, one assumes a design pinch and approach point in the unfired mode and estimates the duty and heating surface of each section. If the surface area selection and steam generation are satisfactory, the design is finalized. One then evaluates the off-design performance based on other gas turbine exhaust conditions. Note that the exhaust gas flow and temperature of a gas turbine is impacted by the ambient temperature, altitude, load of the gas turbine and whether steam or water is injected for NOx concerns. Since a change in inlet gas conditions affects the gas/steam temperature profile and steam generation and temperature, several off-design performance calculations are done before the design is firmed up. One must keep in mind the following: 1.
2.
3.
The economizer may generate some steam at low loads but the design should ensure that it is minimized or eliminated. The exhaust gas flow remains nearly constant at all loads but the exhaust gas temperature decreases as the gas turbine load decreases. At low inlet gas temperatures to the HRSG, the steam generation will be reduced. The economizer flow is also correspondingly reduced. However, the gas flow remains nearly the same as in the design case. The overall heat transfer coefficient (U) is high as the gas side flow governs U and hence the energy transferring ability of the exhaust gases is unaffected while the water flow is reduced, resulting in economizer exit enthalpy that exceeds the saturation liquid enthalpy, causing ‘steaming’, which should be avoided if possible by proper selection of the design approach point. If we use a reasonable approach point at the lowest continuous load of the gas turbine, then as the turbine load increases, the approach point can only increase and steaming is avoided. However, this increases the exit gas temperature and lowers the HRSG efficiency. Steaming in the economizer can cause deposition of solids inside tubes and also results in tube vibration or flow maldistribution/blockage inside tubes. The sizing of HRSG heat transfer surfaces should be reasonable and cost-effective. A small pinch point generally pays off in the long run. Designs have been done using as low as 5 to 6 8C pinch. The overall cost of the HRSG, which includes piping, controls, duct work, etc., may not go up in proportion to the evaporator surface area and hence this suggestion should be explored. In multiple pressure units, the selection of pinch/approach is done to ensure that steam generation at each pressure level is acceptable. The pinch point may be very high in some evaporators if the steam requirement is low at that pressure. For example, in a dual pressure HRSG, if the HP steam demand is low and the LP steam demand is
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Power plant life management and performance improvement high, it is prudent to use a high pinch point for the evaporator and see if the gas temperature entering the LP section is high enough to generate the steam required. Also one should spend some time in arranging the heating surfaces so that the gas/steam temperature profile is optimized.
16.3
How pinch and approach points affect HRSG size and steam generation
How and why pinch and approach points affect the HRSG steam generation, size and gas/steam temperature profiles can be explained through the following simplified analysis for a single pressure HRSG. Considering the superheater and evaporator as shown in Fig. 16.8 we have Wg 6Cpg 6ðtg1 tg3 Þ ¼ Ws 6ðhso hw2 Þ
½16:1
Considering the economizer also, we have Wg 6Cpg 6ðtg1 tg4 Þ ¼ Ws 6ðhso hw1 Þ
½16:2
Dividing equation [16.1] by [16.2], we have ðtg1 tg3 Þ=ðtg1 tg4 Þ ¼ ðhso hw2 Þ=ðhso hw1 Þ ¼ K
½16:3
For steam generation to occur, two conditions should be satisfied, namely tg3 > ts
and
tg4 > tw1
½16:4
If pinch and approach points are arbitrarily selected, it is likely that temperature cross can occur (this situation theoretically means that the gas temperature leaving the evaporator can be lower than the saturation temperature or the exit gas temperature from the economizer can be lower than the feedwater temperature!). Table 16.4 shows K values at different pressures and the effect of steam pressure on the exit gas temperature. It is seen that as the pressure increases, the exit gas temperature also increases. Also, with an increase in the degree of superheat temperature, the exit gas temperature increases as the amount of steam generated is reduced and the heat sink (economizer) becomes smaller. Using a pinch point of 11 8C and an approach point of 8 8C, with a gas inlet temperature of 483 8C and feedwater at 110 8C, Table 16.4 can be generated as the right side of equation [16.3] is a function of steam properties. Example 1: Determine the HRSG exit gas temperature using a pinch point of 11 8C, an approach point of 6 8C and gas inlet of 482 8C at 686 kPa (g).
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Heat recovery steam generators Table 16.4
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K values
Saturated K Steam pressure Steam temperature (8C) temperature (8C) (kg/cm2g)
Exit gas temperature (8C)
7 28.1 28.1 42.2 42.2
149 178 186 189 203
170 231 316 255 399
170 231 232 255 256
0.904 0.7895 0.7728 0.74 0.7728
Solution: The saturation temperature at 7 kg/cm2g is 170 8C. The gas temperature leaving the evaporator, tg3 is 181 8C and the water temperature entering is 164 8C. Using a K value of 0.904, we have (482–181)/(482–tg3) =0. 904 or tg3=149 8C. The following example illustrates why pinch and approach points should not be arbitrarily selected. They are better selected at unfired conditions. Example 2: With 871 8C, can we use a pinch point of 11 8C? The steam pressure is 4.11 MPa. Solution: The saturation temperature is 256 8C and tg3=267 8C. Then (871–267)/(871–tg4)=0.7728 or tg4= 89 8C, which is below the feedwater temperature of 1108C. This is not possible. Only a higher pinch point would ensure a feasible gas/steam temperature profile. This is called a temperature cross situation – a thermodynamically impossible case. Example 3: At 2.75 MPa saturation, can we get an exit gas temperature of 149 8C? Other parameters are the same as before. Solution: Using K=0.7895, 0.7895=(482–tg3)/(482–149) or tg3=219 8C, which is less than the saturation temperature of 231 8C. This is a temperature cross situation! Hence, from these examples, we note the following: 1.
Pinch and approach points should be selected in the unfired mode of operation,(even if the HRSG is fired) using design gas flow and exhaust gas temperature. These are called design ‘pinch and approach points’. One should not try to select the pinch and approach points in the fired case as that can lead to a temperature cross situation as discussed above. It is known from experience that in the unfired mode, pinch and approach points of even 5 8C are feasible as extended surfaces are used in these applications. (When one is forced to use bare tubes, such a low
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2.
3.
4.
Power plant life management and performance improvement pinch point may result in an unusually large HRSG and hence larger pinch points as dictated by shipping considerations and cost may have to be used.) In the fired mode, however, if the approach point of 5 8C is selected, the economizer is likely to steam in the unfired mode. The HRSG size will also be manageable if pinch and approach points are selected in the unfired mode. If the gas turbine operates at a low load for a long time (not transient or startup), then it is a good idea to check whether there is economizer steaming at this load condition or even design the HRSG for this limiting condition by using a higher approach point to avoid steaming. Another reason for not arriving at the temperature profile in the fired case has to do with the steam temperature. If we select the steam temperature in the unfired mode, as the firing temperature increases, the steam temperature will increase, which can be controlled by spray desuperheating. On the other hand, if we select the steam temperature in the fired mode, the extent of oversurfacing or spray cannot be determined easily. One has to check the HRSG performance in unfired modes to ensure that the steam temperature is achieved in these cases too, which is not a smart way to design the HRSG! Once selected, the pinch and approach points fall in place for other gas inlet conditions; that is, there is only one design case and several offdesign cases. HRSG exit gas temperatures cannot be arbitrarily fixed. As steam pressure increases, the exit gas temperature also increases, as seen in Table 16.4. The author has seen specifications calling for a desired low stack gas temperature with single high pressure HRSG. It is not thermodynamically possible to obtain a desired stack gas temperature at a given steam pressure even if infinite surface areas are used for the boiler and economizer unless we add additional modules such as water heater/low pressure evaporators and so on. In the fired mode, the exit gas temperature will be lower as the economizer acts as a bigger heat sink with the larger flow of water through it compared to the design, unfired mode.
Once pinch and approach points are selected in the design of the unfired case, the HRSG steam generation, surface areas are fixed. Then for other cases, one merely checks what the steam generation and temperature profiles are. Using the same figure as above (Figure 16.8), let us assume the design pinch and approach points in the unfired mode: Q12 ¼ Wg 6Cpg 6ðtg1 tg3 Þ ¼ Ws 6ðhso hw2 Þ
½16:5
Since tg3 is known through pinch point selection, as also the water and steam
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enthalpies, we can determine the steam generation from the above equation. Once the steam flow is established, we can determine the superheater duty using: Q1 = Ws 6 (hsohs). The gas temperature leaving the superheater can be found from this; that is, tg2 = tg1Q1/(Wg 6Cpg). The evaporator duty Q2 is obtained from Q2 = Q12Q1. The economizer duty Q3 = Ws 6 (hw2hw1). Since Ws is known, the duty Q3 is obtained. The exit gas temperature is obtained using: Q3 = Wg 6 Cpg 6 (tg3tg4). Since tg3 and the other variables are known, tg4 is calculated. In the above equations, the blowdown, variation in gas specific heat with temperature and heat loss were neglected for simplicity. In an actual design or performance evaluation, all these factors can be considered accurately. In addition to estimating the gas/steam profiles, the selection of pinch and approach points also determines the surface areas of the superheater, evaporator and economizer indirectly. The superheater surface =S1 = Q1/ ΔT1/U1. Since Q1 and ΔT1 are known, one can, based on the geometry, obtain the overall heat transfer coefficient U1 and hence the surface area S1. Then DT1 ¼ ½ðtg1 tso Þ ðtg2 ts Þ= In½ðtg1 tso Þ=ðtg2 ts Þ
½16:6
Similarly S2, the evaporator surface, is obtained from S2=Q2/ΔT2/U2 and S3=Q3/ΔT3/U3. Generally a counterflow arrangement is used for superheaters and economizers in order to obtain a higher steam temperature and also lower the stack gas temperature. Now that the design is done, the off-design performance has to be evaluated. Since a gas turbine HRSG operates at different inlet gas flow and temperature conditions, one has to check how the HRSG performs at various gas inlet conditions. For example, at low loads of the gas turbine, the exhaust gas temperature will be lower and hence the steam generation will be reduced. The economizer, which has been sized for a larger steam flow, is also likely to steam. Hence a wide range of operating conditions have to be checked. The logic for off-design performance is quite involved, particularly if multi-pressure HRSG is involved. For a simple case of superheater, evaporator and economizer, the following Fig. 16.9 shows the logic. Several iterations are involved in this process and a computer program helps.
16.4
HRSG simulation
While designing an HRSG, one first arrives at the duty as shown above for each surface and then computes the overall heat transfer coefficient U, the log-mean temperature difference and then the surface area S required. To arrive at U, one should know the tube geometry, tube lengths, spacings, fin geometry, gas velocity, etc. In short, the physical dimensions of the HRSG
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16.9
Logic for off-design HRSG performance calculations.
are arrived at through this process. However, this is typically done by HRSG suppliers. How can one planning a cogeneration project with the gas turbine HRSG arrive at the plant’s steam generation capability or performance priory? HRSG simulation is the answer. In the case of HRSG simulation, one does not compute the overall heat transfer coefficient U. The term US is computed as an entity, US =Q/ΔT, for each surface such as the superheater, evaporator and economizer. Then in the off-design case, the term US is corrected for the effects of gas flow, temperature and analysis, as detailed in reference [1]. Once the corrected (US)c for each surface is known, the NTU (number of transfer units) method is need is applied to arrive at the duty of each heating surface or the equation Q=(US)c ΔT is used. Then using the logic discussed in Fig. 16.9, the complete HRSG performance, gas/steam temperature profiles in various off-design unfired or fired cases is arrived at. Simulation has several advantages and is a valuable tool for consultants and plant engineers who want to do some preliminary studies on the HRSG without knowing about its size or geometry or its design. They may do these studies on their own without approaching an HRSG supplier, which can
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save a lot of time. All they should have is some knowledge of pinch and approach points. Using simulation, one can: .
.
.
. .
. .
.
Obtain design gas/steam temperature profiles and duty of each surface in the design case. They can play with the HRSG configuration to optimize the temperature profiles or see if multiple pressure steam generation is required or not. Evaluate the steam generation at different steam pressures so that an appropriate steam turbine may be selected. For example, if steam turbines operating at, say, steam pressures, of 100 bar and 60 bar are available, one can quickly do a simulation run and evaluate the steam generation of the HRSG based on given exhaust gas conditions at these pressure levels and then estimate the resulting power generation without approaching an HRSG supplier for this simple information. Obtain an off-design HRSG performance (unfired/fired) at different offdesign gas turbine operations. One may also check if the HRSG needs to be a fired one based on the plant’s needs. Keep in mind that for the gas turbine exhaust gas flow, temperature varies with ambient temperature, load, altitude, whether inlet air chilling is used, etc., and hence performance under all these cases is valuable information for the plant engineer simulating the HRSG performance. Evaluate different gas turbines during the initial project planning stages and see how the steam side performance is different even though the electrical power output of the gas turbine may be nearly the same. Maximize energy recovery by modifying HRSG configuration/rearranging surfaces. If the stack gas temperature is high, they can see if adding a deaerator or low pressure steam generation section or a condensate heater helps. Check for economizer steaming at low loads. This is important if the gas turbine is likely to operate at low loads for a substantial length of time. Evaluate field performance and relate it with HRSG performance guarantees. We can simulate the field data in the design mode and then see the performance based on what the proposal guaranteed. Substantial deviations may be investigated for improper design or inadequate surfacing. Write better HRSG specifications by knowing the HRSG’s capabilities. If the consultant or the plant engineer becomes aware of the capabilities of the HRSG in various modes he is in a position to evaluate the actual HRSG designs better and suggest changes.
If one understands the fact that in the early stages of a cogeneration project all one has is information on gas turbine exhaust gas flow, temperature and analysis and there are several data to be generated (without approaching an HRSG supplier), then simulation is the only tool.
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It should also be noted that the HRSG supplier usually does not perform these analyses as he has limited time and is only interested in coming up with an HRSG configuration and its cost. Examples will be provided below to illustrate the above. However, a simple, manual example will illustrate the power of simulation. The offdesign performance of an HRSG evaporator will be evaluated by knowing the field data or design data at a given condition. The energy transferred to the evaporator is given by: Q ¼ Wg Cpg ðT1 T2 Þ ¼ USDT ¼ USðT1 T2 Þ=ln½ðT1 ts Þ=ðT2 ts Þ
½16:7
Simplifying, In½ðT1 ts Þ=ðT2 ts Þ ¼ US=Wg Cpg
½16:8
In a fire tube boiler, U ¥ Wg0.8. For a water tube boiler, U ¥ Wg0.6, neglecting the effects of temperature. Then W0:2 g In½ðT1 ts Þ=ðT2 ts Þ ¼ K1 ts
for a fire tube boiler
W0:4 g In½ðT1 ts Þ=ðT2 ts Þ ¼ K2
for a water tube boiler ½16:9
and
Example 4: A water tube boiler (evaporator only – no economizer) is designed to generate steam at 1.7 MPa with 45 370 kg/h of flue gas at 538 8C. The exit gas temperature is 260 8C. What is the exit gas temperature and steam generation when 40 834 kg/h of flue gas enters the boiler at 5218C and steam pressure is 1.37 MPa? Solution: First compute K2 using design or known conditions: 45 3700.4ln [(538–208)/(260–208)] = 134.4 = K2. In the off-design case, 40 8340.4ln [(521–198)/(T2198)] = 134.4 or T2 = 245 8C. Evaporator duty and steam generation may be now computed. The approximate steam generation is: 208 and 198 8C are saturation temperatures corresponding to 1.7 and 1.37 MPa respectively.
16.4.1 Examples of simulation analysis Example 5: Exhaust gas flow from a gas turbine is 100 000 kg/h at 5008C. Gas analysis is: % volume CO2 = 3, H2O = 7, N2 = 75, O2 = 15. (a) Determine the steam generation and gas/steam temperature profiles if steam at 5.88 MPa is required at 4508C. The feedwater temperature is
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16.10 HRSG gas/temperature profile – design case.
110 8C and blowdown and heat loss are each 1 %. Suggest a pinch and approach point of 7 8C. Solution: Using the HRSG simulation program, results shown in Fig. 16.10 are obtained. It is seen that the exit gas temperature is 208 8C and steam generation is 11 410 kg/h. (b) What happens to the above HRSG if exhaust gas flow is 90 000 kg/h at 375 8C at the same pressure and feed inlet conditions?
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16.11
Off-design part load HRSG performance.
Solution: By running the off-design performance we note that the HRSG generates about 4900 kg/h and also makes steam in the economizer to the extent of about 2 %. The exit gas temperature is also higher. If this is a long term operation, one has to avoid steaming or minimize its impact, which can be done by several methods (see Fig. 16.11). 1.
2.
The feed control station can be located between the economizer and evaporator so that the economizer operates at a higher pressure and hence the saturation temperature is not reached. Exhaust gases can be bypassed at the economizer so that a much smaller gas flows through it and hence does not transfer as much energy to it.
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4.
627
One may check with the HRSG supplier if the steaming occurs in the vertical leg of the economizer so that the bubbles can flow upwards in the region where steaming occurs. If the operating period is not long and the amount of steaming is low, the blowdown valve may be opened wide so that more water flows through the economizer, thereby avoiding steaming. This may appear to be a waste of water, but if there are hammering or vibration problems in the economizer, this solution can be thought of.
(c) If 17 000 kg/h of superheated steam is desired, how much supplementary firing is required and what is the firing temperature? Gas inlet conditions are the same as in Example 1. Assume also that 3000 kg/h of saturated steam is taken off the drum for deaeration/process purposes. Natural gas with 90 % methane and 5 % each of methane and ethane is used as fuel. Solution: The simulation run in Fig. 16.12 shows that the firing temperature is about 673 8C and fuel input is 5.9 MW. The exit gas temperature is 190 8C, lower than in the design case. This is due to the larger heat sink available at the economizer. The ASME HRSG efficiency is also much higher, namely 72 % compared to 61 % in the design case. If one evaluates the additional fuel input versus the additional steam generation duty, one can see that the additional fuel of about 6 MW generates about 6.3 MW of additional steam duty. This means that the supplementary firing is more than 100 % efficient. One can also justify the higher efficiency by comparing what happens in steam generators. When the excess air is increased in a boiler, the efficiency decreases as more colder air is raised to the exit gas temperature by the fuel. In the HRSG, we are not increasing the excess air but decreasing it by adding the fuel because only the oxygen in the exhaust gas flow is utilized for combustion. Hence the improvement in efficiency is due to a lower exit gas temperature as well as the reduction in excess oxygen. If a cogeneration or power plant wants additional steam, it is better to generate it in the HRSG rather than in a steam generator. A steam generator has a typical efficiency of 93 % on a lower heating value basis, while the HRSG generates additional steam at more than 100 % efficiency. Such results may be obtained using the simulation tool and one does not need the services of a HRSG supplier to obtain these facts. The author may be contacted for the HRSG simulation program or readers can visit his web site (http://vganapathy.tripod.com/boilers.html) for more information.
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16.12
16.5
Off-design fired case.
Improving HRSG efficiency
HRSG efficiency is defined by ASME as the energy absorbed by steam/ water/fluids/(enthalpy of exhaust gas entering the HRSG6times flow + fuel firing in the HRSG on an LHV basis). Hence it is seen that by adding multiple pressure steam levels or condensate heaters, deaeration coil, etc.,
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one can improve the efficiency but the payback must be evaluated. The efficiency improves simply because the exit gas temperature from the HRSG comes down. One may also rearrange heating surfaces (if the HRSG is a multi-pressure unit with high pressure and low pressure steam). Since there are more modules, we can arrange them suitably so that the exit gas temperature is reduced. An example illustrates this. Example 6: Assume that the exhaust gas flow is 200 000 kg/h at 500 8C; the analysis is the same as in Example 5. We need HP steam at 61 kg/cm2a at 450 8C and LP steam at 10 kg/cm2a saturated. The feedwater is at 120 8C. Arrive at the HRSG configuration. Use a pinch/approach of 8 8C for the HP section and, say, 7/5 for the LP section. Solution: A simple approach will be to use an HP section followed by the LP as shown in Fig 16.13a, Here the feedwater enters the HP and LP at 120 8C and we generate 10 100 kg/h HP steam and 3536 kg/h LP steam. How can this be improved if one desires more LP steam? Looking at Fig. 16.13b, we see that, if we used a common economizer concept feeding both the HP and LP sections, it is possible to generate more LP steam and cool the exhaust gases further from 196 8C to 183 8C. The LP steam is now 4240 kg/h. What we have done is to increase the heat sink capacity behind the LP evaporator. By increasing the total flow through the common economizer, we can cool the gases from 200 8C to a much lower temperature. Ideas such as this should be employed to see how the exhaust gases can be cooled further by finding a larger heat sink. The HRSG efficiency also increases from 69 % to 72.5 %. The HRSG cost may increase slightly but the payback period will be short. One has to look at these options, particularly if the gas turbine is large and energy recovered can be more. One can also see this in Fig. 16.13, where multiple modules are used to optimize the gas/steam temperature profiles. The simulation program helps the designer to arrive at the best configuration. Once this has been achieved the HRSG can be physically designed.
16.5.1 Recent trends With rules for limits in emissions of NOx and CO in vogue in many countries (in the range of 3 to 10 ppmv), the HRSG design has to be modified by incorporating catalysts, or SCR units as they are called, (selective catalytic reduction units). The catalysts that lower NOx and CO operate efficiently within a gas temperature range, typically 400 to 500 8C. Hence one has to evaluate the HRSG performance at different loads, unfired and fired if applicable, and ensure that the catalyst will operate satisfactorily within this range of temperatures at all loads. This is done by splitting up the evaporators and connecting them to the common steam
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16.13a Multiple pressure HRSG option 1.
drum system by external downcomers and riser piping. The HRSG cost as well as the gas pressure drop is increased by 75 to 120 mm if catalysts are used and hence the increased backpressure affects the gas turbine power output. Another trend is the quick start and design for cycling, which means more thermal stresses in thick walled components of the HRSG, such as superheater headers and drums, particularly if the steam pressures are high, say around 100–150 bar. In these cases efforts are made to minimize the header and nozzle sizes by having multiple inlet/exit nozzles instead of
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Multiple pressure HRSG option 2.
one: using full penetration welds, using drum heating coils to keep the HRSG warm when the gas turbine is off using smaller diameter tubes to lower the tube wall temperatures and using good steam velocities inside tubes to keep the tubes cooler are some of the methods adopted by large HRSG suppliers. Some HRSG manufacturers unknowingly use large fin densities or external surface areas for superheaters, which increases the heat flux inside the tubes and hence the tube wall temperatures. Using more than 3 fins/in in the superheater is not a good choice. By using appropriate fin
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geometry, the heat flux inside superheater tubes can be kept low and thermal stresses can be minimized [2]. Some other features that are added in these units are: economizer recirculation to minimize the heating time, use of stack dampers to ensure that the HRSG is kept warmer during shutdown, minimizing the infiltration of cold air from outside and hence reducing the startup time, automated generously sized vents and drains for all HRSG sections and multiple attemperators to reduce the superheater tube wall temperatures, particularly in very high pressure units.
16.5.2 Feedwater chemistry Flow accelerated corrosion (FAC) is a recently recognized problem with HRSGs. FAC is a process whereby a normally protective magnetite (Fe3O4) layer on carbon or low alloy steel dissolves into a stream of flowing water or water/steam mixture. Both the PH and temperature and level of dissolved oxygen in the stream influence the stability and solubility of the magnetite oxide layer. Oxidation-reduction potential (ORP) monitoring techniques show good promise as a method to analyze and control feedwater chemistry to minimize FAC. Another promising technique for FAC control in future HRSGs is to fabricate sharp radius elbows out of low chromium alloys, such as the 21/4 Cr–1 Mo variety. This prevents FAC without the need for complex chemistry control methods. Past industry water chemistry practices believe that all of the dissolved oxygen must be eliminated from feedwater to control corrosion. To deoxygenate the feedwater, oxygen was mechanically removed by the condenser/deaerator with supplemental additions of an oxygen scavenger. (e.g. hydrazine) being applied to maintain a 40–100 ppb hydrazine residual. Maintaining an oxygen scavenger residual causes the feedwater to reduce more and more and has for all ferrous systems produced the opposite desired effect of producing a protective oxide film to one where erosion– corrosion of iron based materials is increased. This mechanism is active in condenser shells, feedwater and wet steam piping, feedwater heaters and economizers. FAC occurs in many materials but more in carbon steel piping in the 100 to 250 8C range and hence economizers and low pressure evaporators are affected by this process. Sharp bends where the fluid velocity can erode the protective layer should be avoided. Some European standards require that the maximum O2 in feedwater be increased from 0.02 to 0.1 mg/kg. Flow velocity also has to be lower to reduce erosion of the protective magnetite layer.
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Heat recovery steam generators
16.6
633
Conclusion
HRSG designs have come a long way, particularly for large gas turbines. With multiple pressure level steam generation and multiple modules, the need to optimize the gas/steam temperature profile is obvious. Simulation helps to a great extent. Also advances in gas turbine technology, with higher exhaust gas temperature and large mass flows, warrants the use of reheaters, steam generation at very high pressure and temperature, and the use of condensate preheaters to improve the efficiency. Small cogeneration plants may look at the option of heating thermal fluids such as glycol or therminol in the HRSG or generate low pressure deaeration steam to improve the HRSG efficiency as well as increase the steam turbine power output. Saturated steam from other boilers may also be superheated in the HRSG. Emission regulations have had a great impact on the HRSG design and cost, both operating and installed. Design for cycling has become commonplace.
16.7
Further reading
J. Daiber, ‘Fluid dynamics of the HRSG gas side’, Power, March 2006. HRSG users group, Bozeman, Montana, HRSG Users Handbook, 2009. D. Taylor and A. Pasha, ‘Economic operation of fast starting HRSG’, Power, June 2010.
16.8 1. 2.
References
V. Ganapathy, ‘Simplify HRSG evaluation’, Hydrocarbon Processing, March 1990. V. Ganapathy, ‘Industrial Boilers and HRSGs’, Taylor and Francis, 2003.
16.9
Appendix: nomenclature
Cpg gas specific heat (kcal/kg 8C) hso, hs, hw1, hw2 enthalpy of steam at the HRSG exit, saturated steam, feedwater in and leaving the economizer (kcal/kg) K, K1, K2 constants Q duty (MW), where the subscript 12 refers to energy absorbed by both the superheater and evaporator Q1, Q2, Q3 energy absorbed by the superheater, evaporator and economizer (MW) Q12 energy absorbed by the superheater and evaporator (MW) S surface area (m2) S1, S2, S3 surface areas of the superheater, evaporator and economizer (m2)
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tg1, tg2, tg3, tg4 tw1, tw2 T1, T2 ts tso ΔT U (US)c Wg, Ws
gas temperatures entering the superheater, evaporator, economizer and exit (8C) feedwater temperature entering and leaving the economizer (8C) evaporator inlet/exit gas temperatures (8C) saturation steam temperature (8C) superheated steam temperature (8C) log-mean-temperature difference (8C) overall heat transfer coefficient (kW/m2 8C) product of U and S corrected gas and steam flow (kg/h)
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17 Power plant welds and joints: materials management and performance improvement D . J . A B S O N , TWI, UK
Abstract: Gains in thermal efficiency over many years have been achieved by the development of steels with improved creep properties. However, creep strength is degraded by welding, with the HAZ being the creepweak region. Compared with parent steel development, less effort has, until comparatively recently, been expended in determining cross-weld creep strength and in understanding the reasons for the lower creep strength of the HAZ. This review considers some of the issues arising in low alloy Cr–Mo steels, particularly the 9–12%Cr steels, in creep service in power plants, including consideration of HAZ behaviour. Key words: creep, HAZ, welding, repair, coatings, inspection, failure, type IV.
17.1
Introduction
The drive towards achieving higher thermal efficiency in fossil-fuel power stations is translated into a need for an increase in operating temperatures and pressures. This challenge is being met progressively, at least as far as parent steel properties are concerned (Mayer and Masuyama, 2008). However, the need to accommodate carbon dioxide capture and storage will offset some of the gains in thermal efficiency that have been made in recent years. The welding and service issues relating to the low alloy Cr–Mo and Cr–Mo–V steels, which have been in service for many years, are well understood. Hence, this review will concentrate primarily on the 9–12%Cr steels, where lessons are still being learned, and where code changes are slow to be implemented. Extensive research has been carried out on the newer, 9–12%Cr, parent steels, particularly with regard to their creep rupture strength (Masuyama, 1999; Scarlin et al., 2004; Cipolla and Gabrel, 2005; Viswanathan et al.,
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17.1 Influence of a brief hold at different austenitisation temperatures, plus subsequent PWHT at 570 or 600 8C, on the creep rupture life for grade 91 steel (Middleton and Metcalfe, 1990).
2005), with weld metal development largely having kept pace. However, until recently, there has generally been much less research into cross-weld performance, particularly long-term performance, even though type IV (HAZ) cracking in low alloy power plant steels had been a problem for around 20 years prior to the adoption of grade 91 (9%Cr–1%Mo) steel. The low creep rupture strength of simulated HAZ regions in grade 91 steel, shown in Fig. 17.1 (Middleton and Metcalfe, 1990), should have sounded a warning note, but the early warnings of the consequences of low cross-weld creep strength given by Middleton (1990) and also by Bru¨hl et al. (1990a, 1990b), as discussed below, were not heeded. Grade 91 steel has been widely used, with the ‘discovery’ of its poor cross-weld creep performance arising after a few years in service, when failures occurred after service lives as short as 20 000 h (Brett and Allen, 1999; Brett et al. 1999; Brett, 2001a).
17.2
Materials selection and development
In the UK, 0.5Cr–0.5Mo–0.25V steel found widespread use in coal-fired plant operating at 565 8C and in oil-fired plant operating at <540 8C (Middleton, 1990). In other European countries, X20CrMoV121 (X20) was widely used, generally operating at <540 8C. Grade 91 was developed as a cheaper and more easily fabricated replacement for X20. Steel development continues, with longer term creep testing, modelling studies and transmission electron microscopy all contributing to an increased understanding of the factors that influence long-term parent steel creep strength. This understanding has allowed the development of alloys with progressively higher creep strength. This improvement is illustrated in Fig. 17.2 (von Hagen and Bendick, 2001). A listing of the nominal compositions of boiler steels that are currently being employed is given in Table 17.1 (Cerjak and
© Woodhead Publishing Limited, 2011
© Woodhead Publishing Limited, 2011
0.17 0.23 0.07 0.14
0.50
0.50
0.50
0.50 0.10 0.50
0.25 1.00
0.50 0.15 0.45
0.50
0.35
SI
0.70
1.00
0.30 0.60 0.30 0.60 0.30 0.60 0.30 0.60
0.40 1.00 0.40 0.80 0.10 0.60 0.30 0.70
Mn
10.00 12.50 10.00 12.50
8.00 10.00 8.00 9.50 8.50 9.50 8.50 9.50
0.70 1.15 2.00 2.50 1.90 2.60 2.20 2.60
Cr
0.80 1.20 0.26 0.60
0.90 1.10 0.85 1.05 0.90 1.10 0.30 0.60
0.40 0.60 0.90 1.10 0.05 0.30 0.90 1.10
Mo
0.30 1.70
0.30
0.3
0.30
0.30
Cu
1.50 2.50
0.90 1.10 1.50 2.00
1.45 1.75
W
0.50
0.30 0.80
0.40
0.30 0.10 0.40
Ni
0.25 0.35 0.15 0.30
0.18 0.25 0.13 0.25 0.15 0.25
0.20 0.30 0.20 0.30
V
0.04 0.10
0.06 0.10 0.06 0.10 0.04 0.09
0.02 0.08
Nb
0.0005 0.0050
0.0005 0.0050 0.0010 0.0060
0.0005 0.0060 0.0015 0.0070
B
0.040 0.100
0.030 0.070 0.050 0.090 0.030 0.070
0.010
0.030
0.012
0.012
N
0.05 0.10
Ti
600
103
2
49
113
600
600
98
94
92
147
130
68
49
600
600
550
550
550
550
550
Tw(8C)1 R2 m6106/Iw (MPa)
Service temperature. 100 000 h creep rupture strength at service temperature. Source: Software Stahlschlu¨ssel 2004, Verlag Stahlschlu¨ssel Wegst GmbH, Version 4.00.0005; www.stahlschluessel.de. The T23/T24 Book, New Grades for Waterwalls and Superheaters, Vallourec & Mannesmann Tubes, 2000. ECCC Data Shoots, 2005, compiled and published by D G Robertson and S R Holdsworth, ETD Ltd. METI (Ministry of Economy, Trade and Industry), Thermal Power Standard Code, Japan (single phase with no delta ferrite).
1
T/P122 HCM12A
min max min max
min max min max min max min max
9%Cr steels x11CrMo9-1 (T9) X10CrMoVNb9-1 (T/P91) X11CrMoWV Nb9-1-1 (E911) X10CrWMoVNb9-2 (T/P92)
12%Cr steels X20CrMoV12-1
0.08 0.18 0.08 0.14 0.04 0.10 0.05 0.10
min max min max min max min max
13CrMo4-4 (T/P12) 10CrMo9-10 IT/P22I 7CrWVMoNb9-6 T/P23 7CrMoVTiB1010 (T/P24)
0.08 0.15 0.08 0.12 0.09 0.13 0.07 0.13
C
Low Cr steels
Table 17.1 Chemical composition (wt%) and creep rupture strength of widely used creep-resistant ferritic steels for application in thermal power plants
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17.2 Requirements for creep rupture strength with increasing pressure and temperature (von Hagen and Bendick, 2001).
Mayr, 2008). Austenitic steels are used in the higher temperature sections, such as the final stages of superheaters and reheaters, where increased creep strength and resistance to oxidation are required (Laha et al., 2005). The development of weld metals for creep service has largely kept pace with parent steel development (Heuser and Jochum, 2001; Zhang et al., 2001), and therefore the weld metal itself is seldom the location of failure. However, 9%Cr steel weld metals contain higher levels of Mn and Ni than that of the parent steels, which imposes limits on the PWHT temperatures. For grade 91 steel, de Smet and van Wortel (2006) gave an equation for the A1 temperature: A1 ¼ 848 42½%Ni þ %Mn C They pointed out that this gives an A1 temperature of approximately 785 8C for [Ni + Mn] = 1.5 %, and that this is close to the upper temperature limit for PWHT. Since, with steels that are currently available, the HAZ is the creepweak region, as discussed below, there is little incentive for further weld metal development. However, if type IV cracking problems were to be resolved in new steels, the focus for further development would shift to weld metal.
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17.3
639
Weld/joint degradation
17.3.1 Failure locations The nomenclature for the different types of cracking in creep-resistant steels, due to Schu¨ller et al. (1974), is given in Fig. 17.3, taken from Brear and Fleming (2004). A succinct explanation of the terminology, given by Francis et al. (2004), is reproduced below. (A discourse on a range of cracking mechanisms observed in 0.5%CrMoV steam pipework systems has been given by Brett (1998, 2003)): The cracking of welded joints is usually classified according to the position of the crack; Type I and Type II modes occur within the weld metal, the former confined to the weld metal whereas the latter may grow out of the weld into the plate; Type III cracking occurs in the coarse grained region of the heat-affected zone. Type IV cracking is a feature of welded joints in creep-resistant steels. It is associated with an enhanced
17.3 Classification of cracking in weldments (Brear and Fleming, 2004, according to Schu¨ller et al., 1974).
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17.4 Failure locations for cross-weld samples of grade 91 steel for different rupture lives and temperatures describing the temperature dependence of type IV cracking (Middleton et al., 2001)
rate of creep void formation in the fine grained and intercritically annealed heat-affected zones of the weld, leading to premature failure when compared with creep tests on the unwelded steel. These zones contain coarse carbide particles, leading to a reduction in creep strength; the particles also help nucleate voids. In a cross-weld test, the weakened Type IV region is sandwiched between the stronger base plate and coarsegrained heat affected zone. The resulting accumulation of creep damage in the Type IV region causes the premature failure. In cross-weld creep tests carried out at high stress in the laboratory, failure commonly occurs in the parent steel. However, the failure location shifts to the HAZ as the stress is lowered towards the levels experienced in long-term service. For grade 91 steel, the test duration at which this shift occurs has been neatly illustrated by Middleton et al. (2001) (see Fig. 17.4). While the HAZ region in which type IV failure occurs at comparatively high stress levels is the fine-grained HAZ, it shifts to the intercritical region as the stress is lowered (Abson and Rothwell, 2010). These different regions have been illustrated, and related to the phase diagram, by Cerjak and Mayr (2008) (see Fig. 17.5).
17.3.2 Increasing awareness of low cross-weld creep strength There is now an awareness of a major shortcoming of the original composition specification for grade 91 steel. By examining data from grade 91 steel power station components that had suffered cracking after
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17.5 Schematic representation of subzones of the HAZ corresponding to the calculated phase diagram of X10CrMoVNb9-1 (grade 91 type) steel (Cerjak and Mayr, 2008).
comparatively short service lives, Brett and co-workers (Brett et al., 1999; Brett, 2001a) identified a problem with the Al level of 0.04 % that was originally allowed. By combining with nitrogen, such a high Al level leaves insufficient nitrogen to form the fine-scale MX nitrides that contribute to the creep strength of this alloy. The ASME specification has now been revised, giving an upper limit of 0.02 %Al. However, at the time of writing, the European code EN 10216-2: 2002 allows <0.04%Al for most of the steels for elevated temperature service, including the 9–12%Cr steels. An increasing body of test data is being assembled that shows not only that the HAZ in 9%Cr steels is creep weak, but that the ratio of cross-weld creep strength to parent steel creep strength decreases as the stress is lowered. Some of the early work by Abe and co-workers (Abe, 2005) (Fig. 17.6) was reflected in publications by others, including Abson et al. (2007a), Schubert et al. (2005) (Fig. 17.7), Abson and Rothwell (2010) (Fig. 17.8) and Holmstro¨m and Auerkari (2006) (Fig 17.9). In their analysis of weldment creep data for grade 91 steel, Amato et al. (2010) presented available weld strength factor data as a function of temperature, and concluded that, for long-term high temperature exposure, it tended towards a value of 0.60. Wilshire and Bache (2009) represent creep rupture data by dividing the creep rupture strength by the tensile strength at the test temperature and by incorporating an exponential term containing the activation energy for selfdiffusion into the abscissa. In their elegant analysis of available data, Holstro¨m and Auekari (2009) used the Wilshire approach and ECCC cross-
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17.6 Creep rupture data of Abe (2005) for the parent steel and crossweld specimens of 92 grade and boron-containing steel tested at 650 8C.
17.7 Weld strength factor predicted for various power plant steels after 100 000 hours (Schubert et al., 2005).
weld creep rupture data to derive 100 000 h weld strength factors as a function of stress for grade 91 steel over the temperature range 550 8C to 650 8C (see Fig. 17.9). Data points added to the original graph are the weld strength factors given by Schubert et al. (2005), from which relevant parent steel 100 000 h strength values can be estimated, and 2009 ECCC 100 000 h
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17.8 Creep rupture weld strength factors for grade 92 (Abe, 2005), grade 122 (Abe and Tabuchi, 2004) and grade 91 (Laha et al., 2007). Values derived from creep rupture data of cross-weld or simulated finegrain HAZ specimens tested at 650 8C. A linear extrapolation up to 300 kh has been made (Abson and Rothwell, 2010. Reproduced by permission of TWI Limited).
strength values, from which relevant weld strength factor values can be estimated. These added data points reveal clearly the progressive decrease in weld strength factor at 100 000 h with increasing service temperature. From Figs 17.7, 17.8 and 17.9, it can be seen that, for 600 8C/100 000 h service, weld strength factors (the ratio of weldment creep rupture strength to parent steel creep rupture strength) of <~0.75 are common. Some values of weld strength factor for weldments in a range of steels, appropriate to 100 000 h service at the stated temperature, are given in Table 17.2. A note on the so-called ‘weld strength reduction factor’1 was introduced in the 2008 edition of ASME Section I (in Table PG26:) ‘Weld strength reduction factors to be applied when calculating maximum allowable working pressure or minimum required thickness of components fabricated with a longitudinal seam weld’. For quenched and tempered ‘CSEF’ steels, 1
The use of this terminology, which is common in the U.S.A. and Japan, is deprecated. It refers not (as the wording implies) to a quantity that is deducted from the parent steel creep strength but rather to a ratio of weldment creep rupture strength to parent steel creep rupture strength. The ECCC term ‘weld strength factor’, which is in common use in Europe, is preferred.
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17.9 Stress versus WSF for grade 91 steel derived by Holmstro¨m and Auerkari (2006) from 2005 ECCC data, to which has been added the following data points (Abson and Rothwell, 2010. Reproduced by = 100 000 h WSF values from Schubert et permission of TWI Limited). = 2009 ECCC al. (2005), which indicate associated strength values; 100 000 h strength values, which allow WSF values to be estimated.
the factor is 0.5 for service temperatures from 510 8C to 649 8C. This is a clear recognition, finally, that long-term cross-weld creep strength is substantially lower that that of the parent steel. If such steels are normalised and tempered, the factor decreases progressively from 1.00 to 0.77 over the same temperature range. These changes are having a far-reaching effect on the design of high temperature power plant. As noted earlier, considerable research effort has been expended in generating long-term creep data for the parent steel, with comparatively little effort spent on determining HAZ performance. However, the importance of HAZ creep properties has now been recognised. With an increasing volume of cross-weld data, a better prediction of weldment creep life can be made. More precise data will allow better modelling of system behaviour, including finite element modelling. The information generated will aid both the system design and the development of inspection strategies.
© Woodhead Publishing Limited, 2011
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550 550 550 550 550 600 600 600 600 600
T/P12 T/P22 T/P23 T/P24 T9 T/P91 E911 T/P92 X20CrMoVNb9-2 T/P122 (HCM12A)
Source
Service temperature (8C)
Cerjak and Mayr (2008)
49 68 130 147 92 94 98 113 49 103 0.5 0.68
0.75
0.78
0.69
0.66
0.76
0.86
0.73
Kubonˇ and Schubert Laha Kimura Vlasa´k Holstro¨m and Sobotka (2004) et al. (2005) et al. (2007) et al (2008) et al. (2008) Auekari (2009)
0.71
100 000 h creep Weld strength factor strength at service temperature (MPa)
Weld strength factor for weldments in a range of steels, appropriate to 100 000 h at the stated service temperature
Steel identity
Table 17.2
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17.4
Application of degradation-protection technologies
While some protective coatings may be applied as weld cladding, a range of other techniques is used, including high velocity oxy-fuel spraying, slurry coating and, for extreme service, co-extrusion. The application of coatings for corrosion protection is only justifiable if an increase in lifetime is effected, preferably by a factor of three or more (Meadowcroft, 1987). The selection of coating procedures and compositions involves a number of criteria, including component size, component design lifetime, field or shop applicability and cost (Stringer, 1987). Agu¨ero (2006) added component geometry, compatibility of the coating with the substrate and the application temperature (plus any subsequent heat treatment). All the coating materials and surface modification techniques explored for the protection of steam generating plant at elevated temperatures have advantages and disadvantages; hence, a number of issues remain to be resolved before oxidation-resistant coatings can be used in steam turbines (Agu¨ero, 2008). Erosion–corrosion interactions in equipment associated with energy production have been discussed and represented schematically by Wright (1987). Two forms of corrosion in coal-fired boilers have been discussed by Stringer (1987). These are superheater corrosion (which is a molten salt form of attack, caused by the deposition of sodium, potassium and other sulphates on the surfaces of the tubes (Meadowcroft, 1987)) and accelerated water wastage of the water walls in the general vicinity of the burners. The use of co-extruded tubes with an outer layer of 310 stainless steel has proved beneficial in severe cases of superheater corrosion, and also, as an alternative to plasma-sprayed coatings, to combat water-wall damage (Stringer, 1987). A very different problem is encountered in oil-fired boilers, where the principal form of high temperature degradation is V2O5–Na2SO4 molten salt corrosion of superheater tubes, for which correct adjustment of the fuel:air ratio and the addition of fuel additives are beneficial. An increase in the vanadium to sodium ratio is associated with an increase in the corrosion rate (Meadowcroft, 1987). Agu¨ero (2006), who discussed a wide range of coating processes for use in steam plant, noted that diffusion aluminide coatings are not suitable where nitrides contribute to the strengthening mechanism of the substrate as, for example, in grade 91 and 92 steels. Thermal spray coatings based on aluminium and chromium oxides are very resistant to steam oxidation, and do not require a heat treatment. They do, however, clearly require line-ofsight during application. Agu¨ero also pointed out that, while much laboratory data on steam oxidation resistance has been generated,
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laboratory conditions do not reflect accurately those experienced in service, and there is a comparative dearth of service-exposed data. She has, however, reported satisfactory performance of slurry aluminide coatings in a test tube attached to a steam turbine operating at 640 8C, but commented that a number of issues remain to be resolved before oxidation-resistant coatings can be used in steam turbines (Agu¨ero, 2008). Among the coatings to improve oxidation resistance that can be sprayed on to alloy steels in steam service are Ni–Cr powders (for which minor additions of Al and Ti can improve coating adhesion (Higuera Hidalgo et al., 2001)) and Al. Surface preparation is important, with shot-blasting being the usual process. High velocity oxy-fuel (HVOF) gave both the best surface adhesion and coatings that were substantially free from porosity, although plasma spraying and flame spraying were also employed. Coating layer thicknesses were up to approximately 60 μm. In their investigation of the steam oxidation resistance of HVOF and plasma-sprayed 50%Ni–50% Cr coatings applied to 9%Cr–1%Mo steel substrates, with exposure times of up to 1000 h, Sundararajan et al. (2003) found that scale growth at the interface occurred for plasma-sprayed coatings, when tested at 750 8C, but not for the denser HVOF coatings. In a later study (Sundararajan et al., 2004), a 50%Ni–50%Cr coating was observed to perform better than 80% Ni–20%Cr. In more demanding environments, such as superheater tubing in oil-fired or coal-fired boilers, greater layer thicknesses are required, with co-extruded tubing becoming more attractive (Meadowcroft, 1987). Duallayer coating systems consisting of a thin, ductile oxidation-resistant layer of an expensive material (MCrAlY, Stellite or WC-Co) protected from direct exposure to steam by a thicker lower cost layer (electroless Ni, Al, Al-Si or Cr alloys) appear worthy of consideration (Scarlin et al., 2003). Among the issues that require further in-depth study are coating–substrate interdiffusion, protective oxide stability in high pressure steam and the spalling mechanisms of protective scale and overlay coatings (Agu¨ero, 2008).
17.5
Impact on power plant performance/life management
For many years, the UK power industry has been fighting a ‘rearguard action’ to monitor the creep rupture performance of welded joints in 0.5Cr– 0.5Mo–0.25V and 2.25%Cr–1%Mo, in the light of the occurrence of type IV cracking in these steels after long-term service. However, prior to the widespread use of grade 91 steel, its cross-weld properties appear largely not to have been studied and relevant information that was available did not achieve wide acceptance. As noted above, as early as 1990, Middleton warned of the risk of type IV cracking in grade 91 steel: ‘The operation of
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supercritical plant constructed from either P91 or X20 at ~6008C would therefore involve a degree of risk of type IV cracking, the level depending on the medium to long term strength loss found for cross-weld testing, unless a modest overdesign were applied.’ A similar warning note was sounded by Bru¨hl et al. (1990a, 1990b) who measured cross-weld creep strengths that were substantially lower than parent steel values. They noted that, while weld metal, HAZ and parent steel are subjected to the same strain for circumferential welds, the lower HAZ creep strength of P91 is of concern for longitudinal welds, as they are subjected to the same stress. However, these warnings were not heeded, and grade 91 steel is in service in situations where it is essentially under-designed. Hence, the ‘rearguard action’ is also now being fought on grade 91 steel. For example, Brett et al. (1999) and Brett (2001a) describe the catastrophic failure of a superheater header end plate made of forged grade 91 steel. Failure took place at the type IV region of the heat affected zone on the end plate side of the end plate-to-header weld after only 36 000 hours in operation at 586 8C. Strategies for monitoring critical areas are now in place; see, for example, Brett and Allen (1999). The monitoring and replacement or repair of some power station components such as this is now a major issue for the power industry.
17.6
Dissimilar joints
In power station plant, dissimilar joints occur where alloy steels with different compositions (which may have been chosen because they operate in different temperature regimes) meet or where a component of one steel composition is removed from service and replaced by another of a different composition. In power plants, the existing material may already have suffered some creep damage and the HAZ thermal cycle will inevitably shorten the creep life of the regions adjoining the weld. Hence, low heat input weld beads should be used, at least where they impinge on the side of the joint with the lower creep strength. Carbon diffusion from the lower chromium regions during any PWHT and during service, creating a decarburised zone, will cause further lowering of creep strength. The practice of creating dissimilar joints in low alloy ferritic or bainitic steels and the choice of welding consumables is well established. However, experimental work is continuing on combinations that include the newer 9–12% Cr steels. In existing power stations, dissimilar joints involving higher alloy (9–12% Cr) steels can occur, for example, where a low alloy steel header is replaced by a higher alloy steel, typically grade 91. In principle, there are several possible welding consumables that could be used for such a joint. The obvious contenders are consumables that are near-matching for grade 22 or grade 91. However, additional consumables have been tested experimen-
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tally, for example 5%Cr (Roman et al., 1998) and grades 23 and 24 (Heuser et al., 2008). The issue has been addressed by several investigators, examples of some of which are given below. While their findings do not lead to a clear choice of welding consumables, the use of a consumable that is a near-match for the parent steel of lower creep strength, as advocated by Allen (2002) and reviewed below, appears to give the most satisfactory approach. A table giving the recommended weld metal compositions is presented here as Table 17.3 (AWS, 1996). Roman et al. (1998) buttered a grade 91 steel weld preparation with 9%Cr steel weld metal and then gave an intermediate heat treatment at 765 8C. They also buttered a further grade 91 steel weld preparation with 5%Cr steel weld metal and proceeded without any intermediate heat treatment. Butt welds to grade 22 steel were then completed for these and for unbuttered preparations, followed by PWHT at 720 8C; grade 22 weld metal was used in all cases, except that joint completion with 9%Cr steel weld metal followed buttering with 9%Cr weld metal. For the unbuttered joints, carbon migration occurred at both sides of the weld. However, failure in creep tests occurred in the grade 22 parent steel, presumably because the tests were short term. The 5%Cr steel buttering plus grade 22 weld metal fill gave lower cross-weld creep strength than the unbuttered joint welded with grade 22 weld metal, and the joint buttered with 9%Cr steel weld metal gave yet lower cross-weld creep strength. In their investigations, Heuser et al. (2008) gave a PWHT at 740 8C following their welding of dissimilar joints with grades 23 and 24 weld metal. They detected only slight hardness reductions close to the fusion boundaries, but did not report any cross-weld creep data. Brozda et al. (2006) also did not report any cross-weld creep data of their grade 23/grade 91 dissimilar joints, which they subjected to long-term exposure tests (of 3000 h and 10 000 h) at 550 8C and 600 8C. HAZ and weld metal toughness were generally low, with grade 23 weld metal being preferred to grade 91 because of its higher toughness. The ranking of different joints by Voda´rek et al. (2009) is at variance with many other investigations. They also joined grade 23 and grade 91 steels, with near-matching weld metal for either the grade 23 or the grade 91 side. They gave a 750 8C/2 h PWHT to both welds. Their creep testing at 500 8C, 550 8C and 600 8C revealed substantially longer creep lives where the grade 91 weld metal was used, with the failure location depending on the testing parameters and local microstructural characteristics. By contrast, the predominant failure location for welds deposited with grade 23 weld metal was in the partially decarburised zone of the weld metal, close to the fusion boundary with the grade 91 steel. Allen (2002) carried out cross-weld creep testing after depositing butt welds between grade 22 and grade 91 pipes. Buttering the grade 91 pipe with
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Table 17.3 Suggested filler metals for various parent metal combinations (AWS D10.8-96, Table 2, reproduced with permission from the American Welding Society (AWS), Miami, Florida) Base Metals
C-Steel
C-Steel C-Mo 1/2-1-1/4 Cr-Mo 2 & 2-1/4 Cr-Mo 3 & 5 Cr-Mo 7 Cr-Mo 9 Cr-Mo 9 Cr-Mo-V 3XX-SS
C-Mo 1/2-1-1/4 2 & 2-1/4 3 & 5 Cr-Mo Cr-Mo Cr-Mo
1 A A A A A A A I3
A B B B B B B B I3
A B C2 C C C C C I3
A B C D D D D D I3
Letter
Composition
SMAW6
A
C-Steel
E7108
B
C-Mo
E7018-AI
C
1-1/4 Cr-Mo
D
2-1/4 Cr-Mo
E
5 Cr-Mo
F
7 Cr-Mo
G
9 Cr-Mo
H I
9Cr-Mo-V Ni-Cr-Fe
E70l8-B2L or E8018-B2 E80l8-B3L or E9018-B3 E80l8-B6 or B6L B502-158 E8018-B7or B7L E7 Cr-158 E8018-B8 or B8L E505-158 E9018-B9 ENiCrFe-2 or 3
7 9 9 3XX-SS Cr-Mo Cr-Mo Cr-Mo-V
A B C D E E E E I3
A B C D E F F F I3
A B C D E F G C I3
A B C D £ F G H I3
GTAW/CM AW
FCAW
ER70S-2 or 3 ER70S-B2L or ER80S-B2(7) ER70S-B2L or ER80S-B2 ERSDS-B3L or ER90S-B3 ER80S-B6 ER5029 Note 7
E7X-T5M.-T9M, or -TI2M E7XT5-A1M or E8XT1-A1M E8XTX-B2LM or E8XTX B2M E9XTX-B3UM or E9XTX-B3M Note 7
ER80S-B8 ER5059 ER90S-B9 ERNiCr-3
I3 I3 I3 I3 I3 I3 I3 I3 4
Note 7 Note 7 Note 7 Note 7
Notes: 1
See AWS DI0.12, Recommended Practice and Procedures for Welding Plain Carbon Steel Pipe. For 1/2 Cr-Mo base metals. E80I8-B1 and E8xTl-B1M or -B1/LM may be used for SMAW and FCAW respectively. 3 For non-cyclic thermal service up to about 6008F (3158C), type E309 or E309Mo stainless steel filler metal is frequently used. 4 See AWS D10.4. Recommended Practices for Welding Austenitic Chromium-Nickel Stainless Steel Piping and Tubing. 5 These classifications are shown in the following AWS filler metal specifications for the applicable welding processes: 2
Material
SMAW
GTAW/GMAW
FCAW
C-Steel Low-Alloy Steel Nickel Alloys 300 Series SS
A5. I A5.5 A5.11 A5.4
A5.18 A5.28 A5.14 A5.9
A5.20 A5.29 — A5.22
See appropriate AWS filler metal specifications for other welding processes. 6
While electrode with type 18 coverings or coatings are shown on this table, electrodes with type 15 and 16 coatings are equally acceptable. Filler metals with matching chemical compositions have not been classified by AWS and may not be commercially available. 8 Newer E8018-BX are classified in AWS specification A.5.5 . . . . . . . . . . . . . . . Older E50X-XX and E7Cr are classified in AWS specification A.5.4 9 Newer ER80S-BX are classified in AWS specification A.5.28 . . . . . . . . . . . . . Older ER5OX are classified in AWS specification A.5.9. 7
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9%Cr steel weld metal was followed by a 765 8C/3 h intermediate heat treatment, with a 720 8C/3 h PWHT following completion of the weld with the same consumable. The alternative welding procedure was to weld with grade 22 weld metal, followed by a 730 8C/3 h PWHT. A further heat treatment simulated service by 600, 650 or 730 8C ageing. Allen reasoned that the constraint conditions between a uniaxial creep test and plant conditions are similar when 9%Cr weld metal is used, but that the reduced constraint in a uniaxial test when a grade 22 weld metal is used (because of the development of a wider decarburised zone) gave a pessimistic assessment of cross-weld creep strength. From a detailed assessment of his extensive creep rupture data, he concluded that the preferred approach is to use a grade 22 weld metal. It will be noticed that, in Table 17.3, an established nickel-based alloy (ENiCrFe-2 or ENiCrFe-3) is the recommended consumable for joining an alloy steel to an austenitic stainless steel. However, Coleman (2007) recommends that the EPRI P87 consumable (which only contains approximately 9 % chromium) developed by EPRI is used, since ENiCrFe-3 (alloy 625) may embrittle at the PWHT temperatures required for the alloy steels. Coleman noted that buttering a low alloy steel with EPRI P87 weld metal and subjecting the buttered preparation to a PWHT in the shop provides a means of avoiding the need for a PWHT of the field weld. A further means of avoiding PWHT in the field is to weld a short Nibase alloy transition piece to the lower alloy steel and perform a PWHT in the shop, when the Ni-base alloy to stainless steel weld performed in the field does not require a PWHT. The consumable will also tolerate solution treatment of the alloy steel, thereby eliminating the type IV region. In his report, Coleman (2007) also gives a table for the selection of PWHT temperatures for dissimilar metal joints. The consumable developed initially was a manual metal arc electrode; however, a solid wire variant has subsequently been developed (Siefert et al., 2010).
17.7
Inspection and hardness testing
Increasing use is being made of risk-based inspection, giving improvements in safety and in the efficient use of resources; see, for example, Cane (2000). Preliminary inspection is likely to be visual, which may be followed by other techniques, including dimensional checks and ultrasonic inspection and crack detection. The examination of process data, such as thermocouple measurements and checks for hot-spots by other techniques, including thermographic inspection, may also follow. For materials in creep service, other techniques may also be used, including the intermittent monitoring of creep strain (for example by using electron speckle pattern interferometry (Penny and Kohlho¨fer, 2005) or continuous monitoring of creep strain
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(using capacitance strain gauges (Parker, 2007)). The preparation of acetate replicas (Parker, 1982; ASTM, 1996), which are capable of revealing the presence of creep cavitation, and also micro-focus X-ray may also be used. In their review of non-destructive techniques for the detection of creep damage, Sposito et al. (2010), conscious of the shortcomings of the use of replicas (which include their inability to detect buried cracks), claim that ultrasonic and potential drop techniques appear to be the most promising techniques, but that further research is needed. In their review of possible strategies for strain monitoring, Maharaj et al. (2009) included digital image correlation (DIC) and auto-reference creep management and control (ARCMAC), and suggested that the two methods might be combined. The measurement of hardness, using a portable hardness tester, such as the Krautkramer Microdur 20 or Proceq’s Equotip, is a valuable tool. In principle, it could be used to assess the extent of softening throughout creep service, but would need an initial measurement of hardness to permit its use for an estimate of the remaining creep life. Where steels in creep service have received an incorrect heat treatment, as is reported to be the case for some grade 91 steel, or where their composition is not ideal, in situ measurement of hardness can be a valuable tool. An example of its use in this way has been given by Brett (2001b), who effectively established a correlation between room temperature hardness and creep strength. See Figure 17.10, which shows creep life versus hardness for grade 91 steel specimens tested at 600 8C at an applied stress of 155 MPa. According to de Smet and van Wortel (2006), following proper heat treatment of grade 91 steel, ‘The hardness of weld metal and base metal will be between 200 and 270 Vickers hardness.’
17.10 Creep life versus room temperature hardness for a range of grade 91 steel samples tested at 600 8C, at an applied stress of 155 MPa (Brett, 2001b).
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653
Repair
Bhaduri et al. (2001) commented that: ‘Repair welding plays a very important role in the economic management of steam generator components. The underlying philosophy of repair welding is not to repair components that have reached the end of their design lifetime; it is to put back in service those components that have prematurely failed due to design or manufacture-related defects. Further, these welding repairs should, as far as practicable, be within the scope of the prevailing codes and regulations. As most of the repair-welding strategies are closely guarded information with component manufacturers, there is a growing necessity for in-house development of repair welding procedures by power utilities in close liaison with their Welding Technology Group.’ The weld repair of components in service at elevated temperatures was reviewed a few years ago (Issler et al., 2004), with examples of repairs that have been carried out included in the article. More recently, Aloraier et al. (2010) discussed the use of different welding processes for repair welding, including temper bead repairs, and repairs carried out without preheat, using Ni-base consumables. They also gave some consideration to the resulting residual stresses. Storesund and Samuelson (2002) reported a survey of forty-four cases of weld repairs carried out to Swedish and Danish power plant, mainly on low alloy steam piping after 100 000 to 200 000 h of service. They also carried out finite element simulation of several possible repair geometries, from which they were able to make some recommendations for weld repair. As a result of their simulations and their survey, they gave the following recommendations: . . . .
Avoid welding procedures that may cause a strongly creep-soft HAZ. Select design solutions that minimise design stresses. Select material for weld repair somewhat overmatched (creep-hard) in relation to the remaining service exposed material. Wide and medium deep geometry for the excavation is optimal for the lifetime of a weld repair.
The finding that depositing a creep-hard repair weld metal gave an increased life resulted from the greater creep strain in the surrounding material shedding a greater proportion of the load on to the harder weld metal compared with a matching or a creep-soft weld metal. This conclusion presumably relates to circumferential welds where (as noted earlier) the weld metal, HAZ and parent steel are subjected to the same strain. It is, however, contrary to conventional wisdom, which is that a lower creep strength weld metal is preferred when a PWHT is not given, in order to allow more rapid relaxation of residual stresses in early service. Controlled deposition procedures were originally introduced to achieve
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fine-grained HAZs in order to avoid reheat cracking in low alloy steels used in the power industry. Extensive studies on computer modelling and controlled deposition welding trials at the UK Central Electricity Generating Board resulted in a two-layer refinement technique for MMA welding (Alberry and Jones, 1980). With the primary emphasis having shifted to repair without PWHT, joint research programmes between The Welding Institute in the UK and The Edison Welding Institute in the USA resulted in the development of a four-layer controlled deposition procedure for Cr–Mo steels (Friedman, 1996). The MMA consumable used was 2.25% Cr–1%Mo, with the low carbon variant being preferred, as it would allow faster relief of residual stresses in service. The decay of residual stresses during simulated service exposure was investigated by Leggatt and Friedman (1996). Weld repair practices and the resulting residual stresses have been reviewed by Aloraier et al. (2010). Some ASME codes permit temper bead weld repair without subsequent PWHT (ASME, 2010a; ASME, 2010c) also given in API 510 (2006). They specify how the MMA welding is carried out. Provision is also made for controlled deposition welding in ASME IX (2010b). With certain specified exceptions, manual, semi-automatic, TIG and plasma welding are prohibited. In the National Board Inspection Code (2009) in the USA, covering repair after service exposure, repair without PWHT was permitted initially for C–, C–Mn, C–Mn–Si, C–0.5Mo and 0.5Cr–0.5Mo steels (in the 1977 issue). However, in the light of data provided by several research programmes, the 1995 issue extended the list of steels to include ASME P4 and P5 (Cr–Mo) steels. Controlled deposition techniques used on the lower alloy creep-resistant steels to avoid the need for PWHT have so far not been widely used for the newer, higher Cr steels. Maximising the refinement and in particular achieving some tempering of the coarse grain HAZ and weld metal microstructure through the application of controlled deposition procedures may be a suitable way to avoid the use of PWHT in some situations. However, the necessary procedures can be time-consuming and difficult to carry out, and rely heavily on the skill of the welder. The challenge is that the transformed HAZ is always martensite and is very resistant to tempering. Coussement et al. (1990) carried out research on 5 mm thick T91 that included use of the ‘half tempering’ technique. They showed that, compared with conventional quenching and tempering, TIG welding and post-weld heat treatment, a higher cross-weld creep rupture strength was determined for weldments deposited after quenching and half tempering (620–650 8C/1 h), and then subjected to a 750–760 8C/1 h PWHT after joint completion, before being creep tested at 575, 600 and 625 8C. The authors commented that, when the tempering after austenitisation is performed at lower temperatures (around 620–650 8C), the fine V-carbides are not
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precipitated. Hence globularisation of these carbides cannot occur in the intercritical zone (where the peak temperatures are in the range of approximately 850–950 8C). In their studies of controlled deposition simulated repairs in both 2.25% Cr–1%Mo and 9%Cr–1%Mo steels, each with a near-matching filler, Bhaduri et al. (2001) investigated the use of both half bead and temper bead repairs. For 9%Cr–1%Mo steel, they preferred the latter, as it resulted in lower residual stresses in the HAZ. Vekeman and Huysmans (2008) used a carefully constructed, controlled deposition TIG procedure for ‘cold’ weld repair of new T91. The usual pre-heat and minimum interpass temperature for grade 91 steel is 200 8C. However, in order to achieve tempering of earlier passes, it is imperative that the temperature must be allowed to fall to ~100 8C to permit complete transformation to martensite in the weld metal and underlying HAZ between passes (or between layers) and prior to any PWHT.
17.9
Future trends
17.9.1 The risk of failure in existing plants The operators of power plant will need to continue to be vigilant in monitoring materials in steam service. While this should allow any necessary repair and replacement to be carried out during planned outages, catastrophic failures cannot be entirely ruled out, particularly as there are almost certainly some problem heats and heat treatments of grade 91 steel in service. In the light of new knowledge that has now been incorporated into some of the standards, checks can be made on steel compositions to look for heats containing high levels of Al, on piping systems to look for longitudinal seam welds (that may reflect under design) and on hardness (that would reveal likely low creep strength). Hence, in addition to other on-going monitoring, some precautionary measures can be implemented to reduce the risk of catastrophic failures, both in existing fabrications and in new build.
17.9.2 Parent steel developments It is likely that some of the heat treatments that are currently employed do not realise the full potential of newer creep-resistant steels. This has been amply demonstrated by the work of Morris et al. (2010), who subjected grade 92 steel to a normalising and double tempering heat treatment schedule (normalising at 1150 8C and double tempering at 660 8C for 3 h + 3 h), thereby increasing the creep life more than three-fold. A further substantial gain in creep life was obtained by reducing the C:N ratio. Cross-
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weld creep data will clearly be of interest, to see whether the cross-weld creep strength and the weld strength factor have improved. An appropriate boron addition improves the creep strength of the parent steel. With such an addition, the transient creep region of the parent steel continued for longer times than for a boron-free steel, which resulted in a lower minimum creep rate and a longer rupture time (Abe et al., 2001). This behaviour was attributed to the stabilisation of the lath martensite microstructure for longer times by the stabilisation of M23C6 carbides enriched in boron and to the stabilisation of grain and martensite lath boundaries by the segregation of boron (Abe, 2007). However, among the interesting developments of recent years has been the production of boroncontaining steels that do not develop a fine-grained HAZ and do not appear to suffer type IV cracking. Such a steel was pioneered in Japan, by Abe and collaborators (for example Kondo et al., 2006, Abe, 2005). They observed this behaviour in a 9%Cr–0.5%Mo–1.8%W–VNb–0.0139%B–0.0015%N steel. Metallographic observation of the HAZ revealed the absence of a finegrained HAZ. Further work on this type of steel is being continued in Europe (Mayr and Cerjak, 2007). The potential of these steels has not yet been realised in power plant. If their early promise is realised, providing for their use will present an interesting challenge for the writers of pressure vessel codes. It has long been recognised that Si has a beneficial influence on steam oxidation resistance in 9 to 12%Cr steels. Therefore, we can expect to see more research effort expended in determining the upper limit of Si that can be added safely to this class of alloys. Abe (2005) reported that, with its 0.7%Si, his MARB2 steel (9%Cr–3%Co–3%W–0.7%Si–VNb0.02%B) has excellent steam oxidation resistance, which was enhanced considerably by a pre-oxidation treatment in argon at 700 8C. Chromium is, of course, also beneficial for steam oxidation resistance. However, it is widely recognised that the creep strength of 12%Cr steels is lower than that of 9%Cr steels. Hald (2009) commented: ‘If the oxidation protection is to be achieved through alloy additions instead of surface coatings, it is necessary to increase the chromium content in the steels from 9 % to 12 %. However, so far all such attempts to make stronger 12%Cr steels have been unsuccessful because the high chromium content introduced severe microstructure instabilities in the tested steels, which led to breakdowns in the long-term creep strength.’ Moreover, Abe (2009) noted considerable variation in the creep strength of different heats of 12%Cr steels, which he attributed to differences in the concentration of the nitrogen that was not combined as AlN or TiN. If we are to avoid the need to apply protective coatings to high chromium steels in order to allow their use at yet higher service temperatures, further attention is likely to be given to developing 12%Cr steels, probably with higher levels of silicon, although it
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will clearly be necessary to solve the problem of microstructural instability. This is likely to involve achieving more precise control over the level of nitrogen. One novel approach is to arrange for the formation of a fine dispersion of Z-phase, as claimed by Danielsen and Hald (2008) in their European patent on 9–12% Cr steels. Their invention relates to steel alloys having a martensitic or martensitic–ferritic structure and containing Zphase [(Cr,X)N] particles, where X is one or more of the elements V, Nb, Ta, and where the Z-phase particles have an average size of less than 400 nm. A further interesting development, reported by Toda et al. (2003) and by Kimura et al. (2009), is a boron-containing steel with a nominal composition of Fe–15%Cr–3%Co–6%W, which is reported to have almost the same creep rupture life at 650 8C as NF616. However, as the authors observe, because of its higher Cr content, this steel is expected to show better oxidation resistance than that of conventional creep-resistant steels. If its promise is realised, it may be employed in future fossil-fired plant.
17.9.3 Weld metals of higher creep strength As noted earlier, weld metal development has broadly kept pace with the development of parent steels, and greater creep strength in weld metals is not currently required, since the creep-weak region lies in the HAZ. This was shown, for example, by Mayr et al. (2007). However, if, or more likely, when steels that are resistant to type IV cracking are developed, further weld metal development will be required. Weld metal compositions differ from those of parent steels. Among other differences, in order to improve toughness, the levels of Ni are higher and Nb lower (Bergquist, 1999). In future weld metal developments, we are likely to see higher levels of Co and W, and the incorporation of boron, as in the experimental weld metal of Mayr et al. (2009). Also, as reported by Abson et al. (2007b), Ti has the potential to give a useful increase in creep strength, although it is difficult to control the recovery of Ti. If 15%Cr steels are developed, we may see weld metal with similar levels of Cr, together with slightly elevated levels of Si, since both Cr and Si give enhanced resistance to steam oxidation.
17.9.4 Repair procedures Some repair procedures, including controlled deposition repair procedures, have been discussed above. We will undoubtedly see further work in this area for the repair of 9–12%Cr steels without a subsequent PWHT. One of the major utilities developed a controlled deposition repair procedure for grade 91 steel, with a view to deploying it without a subsequent PWHT, but had not used it in the repair of plant (Brett, 2007). This is an area of study at TWI, with some novel approaches being investigated.
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17.10 Sources of further information and advice Steels for steam service in power plant are currently receiving considerable attention worldwide, with articles on many different aspects of their behaviour, including creep/fatigue, repair and computer modelling appearing in major journals and in conferences. Further information can be found in the reviews cited in this chapter and in papers that are presented at the major creep conferences, which occur either annually or every two or three years. These include the following: EPRI, Baltica, Liege, ECCC, Charles Parsons, International Conference on Creep and Fracture of Engineering Materials and Structures, IOM Communications and International Thermal Spray Conference.
17.11 Acknowledgements The author is grateful to colleagues at TWI for assistance in the preparation of this chapter, and to the reviewer for some constructive comments.
17.12 References Abe F (2005), ‘High performance creep resistant steels for 21st century power plants’, in Proceedings of 1st Conference on Super High Strength Steels, Rome, Italy, 2–5 November 2005, Associazione Italiana di Metallurgia, Fig. 29. Abe F (2007), ‘Effect of boron on creep deformation behaviour of 9Cr steel for USC boilers at 650 8C’, in Proceedings of Seventh International Charles Parsons Turbine Conference on Power Generation in an Era of Climate Change, University of Strathclyde, Glasgow, Scotland, 11–13 September 2007, pp. 477–488. Abe F (2009), ‘Heat-to-heat variation long-tem creep strength of some ferritic steels’, in Proceedings of Conference on Creep and Fracture in High Temperature Components – Design and Life Assessment Issues, Zu¨rich, Switzerland, 21–23 April 2009, Eds: I A Shibli and S R Houldsworth, European Creep Collaborative Committee, 2009, pp. 5–18. Abe F and Tabuchi M (2004) ‘Microstructure and creep strength of welds in advanced ferritic power plant steels’, Science and Technology of Welding and Joining, 9(1), 22–30. Abe F, Igarashi M, Wanikawa, S, Tabuchi M, Itagaki T, Kimura K and Yamaguchi K (2001), ‘New ferritic heat resisting steels for 650 8C USC boilers’, in Proceedings of 3rd [EPRI] Conference on Advances in Materials Technology for Fossil Power Plants, Swansea, Wales, 5–6 April 2001, Ed: R Viswanathan, W T Bakker and J D Parker. Institute of Materials, London, UK, pp. 79–88. Abson D J and Rothwell J S (2010), ‘Review of type IV cracking in 9–12%Cr creep strength enhanced ferritic steels’, TWI Confidential Research Report for Industrial Members, June 2010. Accepted for Publication in International Materials Reviews, Fig. 16. Abson D J, Rothwell J S and Cane B J (2007a), ‘Advances in welded creep resistant
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9–12%Cr steels’, in Proceedings of 5th International EPRI Conference on Advances in Materials Technology for Fossil Power Plants, Miami, Florida, 2–5 October 2007. Abson D J, Rothwell J S and Woollin P (2007b), ‘The influence of Ti, Al and Nb on the toughness and creep rupture strength of grade 92 steel weld metal’, in Proceedings of Conference on Integrity of High Temperature Welds, 24–26 April 2007, IOM3, London. Agu¨ero A (2006), ‘Coatings for protection of high temperature new generation steam plant components: a review’, in Proceedings of 8th Liege International Conference on Materials for Advanced Power Engineering, Liege, Belgium, 18– 20 September 2006, Eds: J Lecomte-Beckers, M Carton, F Schubert and P J Ennis. Julich, Germany, Forschungszentrum Julich GmbH; also Advanced Steam Turbine Materials, Welding Technology, 2006. 53, Part II, pp. 949–963. Agu¨ero A (2008), ‘Progress in the development of coatings for protection of new generation steam plant components’, Energy Materials, 3(1), 35–44. Alberry P J and Jones K E (1980), ‘Two-layer refinement techniques for pipe welding’, in Proceedings of 2nd International Conference on Pipewelding, London, 20–22 November 1979. The Welding Institute, Abington, Cambridge, UK, 1980, Vol. 1, Session 4, Paper 9, pp. 185–199; discussion Vol. 2, 1981, pp. 383–387. Allen D (2002), ‘High temperature performance of dissimilar P91 to P22 [Cr–Mo] steel weldments’, in Proceedings of Fifth International EPRI (Electric Power Research Institute) – RRAC (Repair and Replacement Applications Center) Conference, Welding and Repair Technology for Power Plants. Point Clear, Alabama, 26–28 June 2002, EPRI (Electric Power Research Institute), P91 Session, Paper P2. Aloraier A, Al-Mazrouee A, Price J W H and Shehata T (2010), ‘Weld repair practices without post weld heat treatment and their consequences on residual stress: a review’, Int. J. Pressure Vessels and Piping, 87, 127–133. Amato A, Moino G and Shingledecker J (2010), ‘Development of weld strength reduction factors and weld joint influence factors for service in the creep regime and application to ASME codes: Task 1b/3 Report’, November 2010, EPRI/ ASME Standards and Technology, LLC project/3052. API 510 (2006), Pressure Vessel Inspection Code: Maintenance Inspection, Rating, Repair and Alteration, ninth edition, June 2006. ASME (2008), Boiler and Pressure Vessel Code, Section I, Rules for Construction of Power Boilers, American Society of Mechanical Engineers, New York. ASME (2010a), Boiler and Pressure Vessel Code, Section III, Rules for Construction of Nuclear Facility Components, Paragraph NB-4622.9, Temper bead weld repair, American Society of Mechanical Engineers, New York. ASME (2010b), Boiler and Pressure Vessel Code, Section IX, Welding and Brazing Qualification, Paragraph QW-290, Temper bead welding, American Society of Mechanical Engineers, New York. ASME (2010c), Boiler and Pressure Vessel Code, Section XI, Rules for Service Inspection of Nuclear Power Plant Components, Paragraph IWA-4623.1, Temper bead welding of similar materials – shielded metal arc welding, American Society of Mechanical Engineers, New York.
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© Woodhead Publishing Limited, 2011
Index
abrasion, 8 acceptance test, 554–5 acid gas removal, 114–15 availability issues, 125–7 COS hydrolysis, 125–6 MDEA, 126–7 rectisol, 126 selexol, 126 acoustic emissions, 56–7 acoustic leak monitoring, 44 active combustion pattern factor controller, 100–1 advanced gas reactor, 549 advanced ultra-supercritical, 478 air-borne noise, 44–5 air separation unit, 111–13, 579 availability issues, 116–19 liquid oxygen storage, 119 reported failures, 116–19 alignment keys and packers, 565 inner casing packers, 565 all-volatile treatment, 477 Alstom, 381–2 alumina, 436 aluminide coatings, 369 aluminium oxide (Al2O3), 361 American EPRI-1403, 224 Andrade’s equation, 342 ANSYS software, 491, 507 approach point, 616 HRSG size and steam generation, 618–21 Armitt type approach, 474 Arrhenius constant, 464 Arrhenius kinetics, 214 ash, 6 ash analysis, 7 ash fusion, 9–10
ASME Boiler and Pressure Vessel Code Section III Subsection NH, 287 ASME Code Case N-47, 233 ASME PTC 19.1-2005, 555 ASTM 4-6, 258 ASTM 8-9, 258 ASTM International, 5 austenitic steels, 460 scale morphology and spallation, 470–2 cross-section of scale formed on Super304H after 4000 hours in steam, 472 typical oxide morphology formed on 300 series austenitic steels in steam, 472 auto-reference creep management and control, 652 biochemical composition, 8 boiler components component operating environments, 147–52 fuel effects, 150–2 schematic diagram of circulating fluidised bed biomass fired unit, 149 schematic diagram of pulverised fuel power plant, 148 schematic diagram of waste-fired grate unit, 149 schematic flow diagram of power plant steam/water system, 147 schematic representation of fuel combustion, 151 solid fuel composition examples, 150
666 © Woodhead Publishing Limited, 2011
Index superheaters/reheaters/waterwalls, etc., 147–50 environmental degradation, 145–75 damage quantification and protective measures, 169–71 degradation mechanisms and modelling, 152–69 future trends, 172–5 boiler materials creep, mechanisms, measurement and modelling, 180–216 creep deformation and damage mechanisms, 181–91 information source and advice, 216 measurement methods, 191–7 operating environment effect, 197– 211 predictive modelling, 211–15 boiler steels, 222–66 characteristics of microstructure and property degradation processes, 248–60 changes in the dislocation density in P92 steel after creep, 257 degradation of microstructure and austenitic steel properties, 258–60 degradation of microstructure and mild alloy steel properties, 250–3 degradation of microstructure and properties of high-alloy martensitic steels of 9–12% Cr group, 253–8 influence of carbide type in the ferrite on the 1Cr-0.5Mo steel durability, 253 influence of carbide type in the ferrite on the martensitic steel durability, 257 microstructure of austenitic steels with isolated primary MC carbides, 259 phase composition of carbides, temperature and operating time on creep resistance of 1Cr-0.5Mo type, 252 VM12 microstructure in the initial state and after service with a strong effect of degradation, 255–6 material assessment and residual durability determination, 230–48 assessment method for the material
667
state of an element composed of the steel 10CrMo910, 235 assessment of pressure component elements made from ferriticmatrix steel, 230–45 assessment of pressure component elements made of austenitic steel, 245–8 creep damage classification in 0.5Cr-0.5Mo-0.25V steel, 245 creep damage classification in 12Cr1Mo-V steel, 242 internal defects during long-term operation, 233 microstructural patterns for 12Cr1Mo-V steel after operation under creep conditions, 243–4 microstructural patterns including defects for 0.5Cr-0.5Mo-0.25V steel, 246–7 microstructural patterns of change in microstructure image of 0.5Cr0.5Mo-0.25V steel, 240–1 microstructural patterns of change in microstructure image of 12Cr1Mo-V steel, 237–8 microstructure change of 0.5Cr0.5Mo-0, +925V steel during long-term services, 239 microstructure change of 12Cr1Mo-V steel during long-term services, 236 ways to approach the evaluation of the state of material or component, 232 microstructural degradation, 222–66 modelling degradation processes and their use, 261–6 austenitic steels, 265 ferritic-pearlitic/bainitic steels, 265 martensitic steels, 265 preparation of classification system for material after operation, 260–1 assessment of the state of component process for 12Cr1Mo-V steel, 263 assessment of the state of component processes for 0.5Cr0.5Mo-0.25V steel, 262 classification method involved in assigning the exhaustion degree with the main structural class, 261
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Index
steel development for power engineering, 223–30 allowable stress at creep strength of martensitic and austenitic steels, 228 comparison of creep strength of bainitic steels with martensitic steel 9Cr–Mo–VNNb type, 226 creep rupture strength for martensitic, austenitic steel and nickel superalloys, 228 nominal chemical compositions of austenitic steels used in power boilers construction, 227 oxidation resistance comparison of (9–12)% Cr steels and austenitic TP347 steel in pure steam, 229 selected bainitic and martensitic steels for thick wall components and coils, 225 Boltzmann constant, 185 bond coat, 84 borescope ports, 561 Brayton cycle, 577 British Standards Institute (BSI), 5 brittle, 339 Brownian diffusion, 156 bulk density, 8 Burgers vector, 182 carbon capture and storage, 34, 114 carbon capture systems, 173 carbon tax, 579–80 casing baffles, 566 illustration, 566 charge-coupled device, 58 chemical scrubbers, 115 chemical vapour deposition external gas generation, 438 internal gas generation, 438 system schematic representation, 437 ChlorOut, 32 chromia, 436 chromide diffusion coatings, 369–70 circulation ratio, 613 co-firing with biomass, 479–80 coal gasification systems, 581–2 effects of temperature and pressure on syngas compositions, 581 coal petrology, 10 coating chemistry, 444–5 metallic systems, 444 thermal barrier systems, 444–5
Coble creep, 185 Coffin–Manson equation, 408 Coffin–Manson relationship, 407 combined cycle, 116 analysed IGCC plants, 117 availability issues, 127–9 gas treatment and sulphur recovery in gasification plants, 118 IGCC ASU summary, 118 combined water treatment, 478 combustion monitoring system, 76 combustion systems, 33 combustion turbines, 392 condition-based maintenance, 411 conditioning monitoring method, 78 convective air heater, 599 convective syngas coolers, 123 coordinate measurement machines, 558 corrosion fatigue, 353 corrosion monitoring, 47–9 conventional and advanced corrosion-coupon probes, 47–8 electrical resistance corrosion probes, 48–9 boiler areas requiring corrosion mapping and results in terms of corrosion rate and overall thinning, 49 COST 501 programme, 524 crack growth rate, 319 cracking, 639–40 creep, 341–7 boiler materials, mechanisms, measurement and modelling, 180– 216 predictive modelling, 211–15 creep deformation and damage mechanisms, 181–91 creep damage mechanisms and fracture, 187–91 creep deformation mechanisms, 181–87 gas turbine superalloys, 308–14 development of the temperature capability of an Ni alloy for gas turbine blades, 310 typical shape and trends of the creep curve, 309 measurement methods, 191–7 creep fracture growth measurements, 194 creep testing under non-steady conditions, 194
© Woodhead Publishing Limited, 2011
Index residual creep life assessment, 195– 6 standard creep testing methods, 192–3 stress concentrations and multiaxial stress testing, 193–4 stress relation testing, 193 synergistics effects measurement, 196–7 very low creep strains measurements, 195 operating environment effect, 197– 211 advanced boiler steels after longterm isothermal ageing, 197–202 creep behaviour intermittent heating effect, 202–11 power plant boilers, 283–5 creep cavities on grain boundaries, 283 creep–fatigue design curves based on linear damage summation method, 287 creep–fatigue interaction, 286–8 deformation process evolution from cavity formation to macrocracks and creep rupture failure, 284 creep behaviour advanced boiler steels after long-term isothermal ageing, 197–202 calculated equilibrium mole fractions, 204 stress dependences, times to fracture for P91, P92 and P23 steels, 198 TEM micrographs, P23 steel microstructure, 203 TEM micrographs, P91 steel microstructure after ageing at 650 8C, 200 TEM micrographs, P92 steel microstructure after ageing at 650 8C, 201 intermittent heating effect, 202–11 creep curves at 150 MPa with different loading histories, P92 and E911 steels, 207 P91 steel creep curves, different loading histories at 125 MPa, 205 scheme of intermittent loading and heating during creep exposure, 205
669
TEM micrographs, E911 steel subjected to creep at 150 MPa, 210 TEM micrographs, P92 steel subjected to creep at 150 MPa, 209 TEM micrographs, P91 steel subjected to monotonic creep at 125 MPa, 208 time to fracture for standard creep tests, cyclic creep, and corrected time cyclic tests, 207 creep damage mechanisms and fracture, 187–91 fracture values of areal fraction of cavities, 190 intergranular creep damage development scheme, 187 creep deformation, 186–7 time to fracture dependence on stress, P91 steel compared to extrapolation, 186 creep–fatigue interaction power plant boilers, 286–8 creep rate, 341 creep strains measurement, 507–10 critical pressure components influence of start-up and shut-down operations on thermal and total stress loads, 491–7 division into finite elements, 492 equivalent thermal and total stress histories during cooling operation, 497 equivalent thermal and total stress histories during heating operation, 496 first method of boiler start-up and shut-down operations, 493 13HMF steel, 492 second method of boiler start-up and shut-down operations, 494 CrMoV, 520–1 CrMoVNbN, 523 cross-weld creep tests, 640, 649, 651 damage assessment-SFM, 297 damage summation model, 405–6 degradation-protection technologies, 646–7 densitometry, 188 deposition, 152–7
© Woodhead Publishing Limited, 2011
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Index
deposit compositions, 156–7 structure deposit observed on superheater tube in coal-fired boiler, 157 direct inertial impaction, 154–5 other deposition processes, 155–6 schematic representation of interaction between superheater tube and local environment, 152 vapour condensation, 153–4 calculation of saturated vapour pressures of alkali sulphates and chlorides, 154 sulphur and chlorine effect and alkali metal availability on deposit compositions, 154 DICTRA, 593 diffusion coatings, 369–70, 436–8 CVD: external gas generation, 438 CVD: internal gas generation, 438 modified systems, 438–9 precious metal aluminising, 439 spray coatings, 439 pack cementation, 437 diffusional creep, 185–6 mechanism, 185 digital control system, 61 digital image correlation, 652 digital speckle photogrammetry, 58 directional solidification, 400 dislocation movement, 182–3 dissimilar joints, 648–51 filler metals for various parent metal combinations, 650 distributed control system, 41 Drakelow C, 281 Drax, 282 dry filtration, 124 dry low emission, 76 ductility, 339 dynamic slag application furnace, 584, 589 eddy current scanning, 69 eddy diffusion, 156 Eddystone B, 281 Electric Power Research Institute, 579 electrical sensitivity, 188 electro-co-deposition, 443 electrochemical noise, 170 electron beam physical vapour deposition, 91, 403, 443 electron beam welding, 431–2
electron microscopy, 169–70 emission monitoring system, 76 EN 10216-2, 641 EN 12952-3, 491, 500 EN-12952-4/2000, 233 ε–N method, 353 energy content, 9 energy dispersive X-ray analysis, 169–70 enhanced oil recovery, 115 environmental barrier coating, 401 environmental degradation boiler components, 145–75 component operating environments, 147–52 damage quantification and protective measures, 169–71 component and materials monitoring methods, 169–71 protective coatings, 171 degradation mechanisms and modelling, 152–69 deposition, 152–7 erosion/wear, 168–9 fireside corrosion, 159–68 oxidation, 157–9 future trends, 172–5 gas compositions produced in oxyfired and air-fired pulverised fuel systems, 174 Equotip, 652 erosion, 288 Esshete 1250, 281 European Committee for Standardisation, 5 European COST 501, 224 European COST 502, 224 European COST 522, 224 European COST 536, 224 fail-safe concept, 350 failure mode effect and critical analysis, 132 fatigue gas turbine superalloys, 314–19 constant amplitude fatigue performance of alloys IN738LC and FSX-414, 316 fatigue crack growth rate of IN738LC, 318 HCF fatigue strength for IN738LC and IN939, 317 temperature dependence of LCF
© Woodhead Publishing Limited, 2011
Index performance for alloy IN738LC, 317 power plant boilers, 285–6 fatigue life, 349 ferritic steels, 460 scale morphology and spallation, 467–70, 471 12Cr ferritics, 470 cross-section of scale formed on MARB-2 after 4000 hours in steam, 471 cross-sections of scales formed on T91 after 4000 hours in steam, 469–70 ferritic steels containing up to 9Cr, 468 oxide morphology formed on T22 in steam, 468 T91, 469 fine grain, 471 fire side, 453 fire tube boilers, 609 fireside corrosion, 159–68 superheater/reheater corrosion, 161–8 bell-shaped curve characteristic, fireside corrosion damage mechanism, 165 metal temperature effect on fireside corrosion, 168 metal temperatures effect on corrosion rates in conventional pulverised-fuel-fired power system, 165 phase diagram for alkali-iron trisulphates, 162 phase diagram for alkali pyrosulphates, 162 phase diagram for alkali– sulphates–chlorides, 163 phase diagram for zinc chloridespotassium chlorides, 164 physical state of complex alkali sulphates and temperature vs austenitic alloys corrosion behaviour, 166 potential corrosion mechanism example, 167 tube lives, steam temperatures and degradation rates in coal, biomass and waste systems, 168 waterwall corrosion, 160–1 deposit/scale structure, corroding waterwall surface, 160
671
fixed carbon, 6–7 flame temperature sensor, 100–1 flow accelerated corrosion, 529–30, 632 inlaid rings for erosion protection, 530 flue gas desulphurisation, 480, 537 fracture mechanics, 353 frequency scanning eddy current method, 87 fuel chemistry, 5–30 biomass fuels, 21–6 biomass analysis standards examples, 23 minor and trace elements present in different biomass range, 25 typical biomass analyses, 24 coal fuels, 10–21 coal analysis standards examples, 12–13 coal classification, 15 coal constituents breakdown, 10 coal rank, 16 coal types and coal carbon and moisture contents relationship, 14 common minerals found in coal, 19 H/C and O/C ratios relationship for coals and biomass, 16 sulphur and chlorine contents, selected biomass and coals, 18 trace elements in coal, 20 typical coal analyses, 17 volatility classification, trace elements in combustion system, 21 specifications and analysis methods, 5–10 chemical analyses, 6–8 other general fuel properties, 9–10 physical analyses, 8 proximate analysis parameters and analysis basis relationship, 7 ultimate analysis parameters and analysis basis relationship, 7 waste derived fuels, 26–30 European waste analysis standards, 29–30 MSW composition data, fuel origins, 27 MSW detailed breakdown from UK survey, 28 waste fuel analysis, 30 fuel injection systems, 384–5
© Woodhead Publishing Limited, 2011
672
Index
fuel preparation systems, 33 fuel transport/handling/storage systems, 32–3 full-arc admission, 542–4 ‘full-scale’ tests, 554 FutureGen plants, 581 gas turbine advanced materials and coatings, 399–403 evolution of high-temperature capability of superalloys, 400 coating and finishing technology, 436–45 chemical vapour deposition system, 437 coating chemistry, 444–5 diffusion coatings, 436–8 high velocity oxyfuel coating system, 442 low pressure plasma system, 441 MCrAlY coating applied by LPPS, 441 modified diffusion systems, 438–9 overlay coatings, 439–43 robotic application of the coating applied to a turbine component, 442 thermal barrier ceramic coating applied over a metallic corrosion resistant coating, 444 competitive arena for power generation, 331 factors affecting maintenance intervals, 423–5 fuel types, 424 starts and stops, 424 trips from load, 425 water/steam injection, 425 field service overhaul and maintenance, 421–6 condition based inspections, 425–6 heavy industrial gas turbine rotor being disassembled, 423 photograph detailing information that can be identified remotely from the control system, 426 final repair operations, 445 in situ assessment, by non-destructive techniques, 84–95 advanced eddy current method, coating thickness and depletion evaluation, 87–90
electrical impedance vs frequency plots, 88 frequency-scanned eddy current (FSECT) measurements, 90 methods comparison, 97 photoluminescence piezospectroscopy, 90–5 PLPS result on blade section after service exposure in the plant, 96 PLPS spectra comparison, 94 PLPS spectra of TGO, 93 PLPS technique and EB-PVD TBC columnar structure, 92 pulsed thermography configuration and blade thermal image, 86 TBC coating example, 85 thermography, check blade cooling and coating adhesion, 85–7 life management and diagnostic, 403– 9 damage summation model, 405–6 strain-range partitioning model, 407–9 maintenance, refurbishment and repair, 420–49 future trends, 448–9 materials applications, 377–99 benefits of single crystal, 394 casing, 399 chemical composition of some common turbine disc alloys, 398 combustion section, 384–7 combustion turbine alloy and coatings for hot section blading, 391 combustion turbine materials, 386 composition of cast DS superalloy for turbine blade application, 393 composition of single crystal superalloys, 395 compressor blade alloys, 384 compressor section, 383–4 evolution of large land-based gas turbine design features, 378 nominal composition of selected nickel-base superalloys used for turbines blades, 390 nominal compositions of selected alloys used for industrial gas turbine vanes, 388 rotor, 395–9 seals, 399 superalloy sheet materials used for
© Woodhead Publishing Limited, 2011
Index industrial gas turbine combustors and transition pieces, 385 turbine section, 387–95 use of iron-base alloys in industrial gas turbines, 397 materials selection, life management and performance improvement, 330–419 future trends, 409–11 part life extension and optimisation, 446–8 stress analysis model of typical turbine blade, 447 typical metallographic sectioning plan for life assessment, 447 parts refurbishment, 426–9 parts repair, 429–36 electron beam weld detailing the typical weld morphology, 433 electron beam welding process, 433 fully repaired crack after a transient phase restoration process, 435 high temperature vacuum brazing, 432–6 hydrogen fluoride cleaning system, 434 laser powder fusion welding system, 431 laser weld on tip section of turbine blade, 432 tungsten inert gas welding system, 430 welding, 429–32 protective coatings, 368–77 commercial overlay coatings, 371 commercial diffusion coatings, 370 diffusion coatings, 369–70 failure mechanism in thermal barrier coating, 377 failure modes, 373–7 multilayer thermal barrier coating system, 373 overlay coatings, 370–2 thermal barrier coatings, 372–3 quality control and first article inspection, 445–6 superalloys, 332–67 gas turbine heat recovery steam generator, 611–18 forced circulation HRSG, 613 furnace fired gas turbine HRSG, 615 HRSG performance, 616–18 gas/steam temperature profile, 616
673
multiple pressure fired HRSG for a large gas turbine, 612 two pass HRSG for a small gas turbine, 612 unfired vs fired HRSG, 615 gas turbine superalloys combined creep and fatigue, 319–22 effect of frequency on the crack growth rate of nickel alloy API, 320 linear life fraction mapping, 321 utilising the life fraction principle for maintenance of gas turbines, 321 creep, 308–14 development of the temperature capability of an Ni alloy for gas turbine blades, 310 typical shape and trends of the creep curve, 309 creep, fatigue and microstructural degradation, 307–28 creep rupture strength cast alloy IN738LC, 312 cast alloy Udimet 500, 313 cast blade alloy GTD-111, 314 wrought alloy Udimet 720, 313 fatigue, 314–19 constant amplitude fatigue performance of alloys IN738LC and FSX-414, 316 fatigue crack growth rate of IN738LC, 318 HCF fatigue strength for IN738LC and IN939, 317 temperature dependence of LCF performance for alloy IN738LC, 317 future trends, 327–8 microstructural degradation, 322–7 agglomeration and rafting of the γγ ´ microstructure, 325 cast blade (GTD-111) microstructure, 323 cross-section of turbine blade with server irreversible, 327 ex-service blade (GTD-111) with coarse partially fused γ ´ particles, 324 ex-service cast turbine blade microstructure, 323 plate-like TCP phase precipitates below the coating, 326
© Woodhead Publishing Limited, 2011
674
Index
single crystal blade microstructure, 324 whitish zone on crack surfaces, 326 gas turbines condition monitoring, 74– 83 combustion instabilities, 76–7 matrix assisting in identifying causes of combustion instabilities, 77 electrically charged particulate at turbine exhaust, 77–81 EDMS result over a six months monitoring period, 80 EDMS sensor and cable, 79 EDMS sensor installed in gas turbine exhaust duct, 80 gas temperature and hot parts integrity, 81–3 pyrometer installation into the GT, 82 pyrometer installed in GT, 83 schematic of gas turbine, 75 gasification unit, 113–14 availability issues, 120–5 feed system, 120 gasifier, 120–2 raw gas cooling, 123–4 raw gas pre-cleaning, 124–5 slag removal, 122–3 gasifier, 113–14, 120–2 burner, 122 corrosion, 122 feedstock, 120–1 reactor wall, 121–2 GE H System™ gas turbine, 379 grain boundary sliding, 184–5, 583 contribution dependence on grain boundary sliding to total creep strain, 184 greenhouse gases, 172 grindability, 8 GT26/GT24 gas turbine, 381 GTD-111, 392 GTD-450, 384 hardness measurements, 188 hardness testing, 651–2 creep life vs temperature hardness, 652 Hastelloy X, 385, 387 Haynes 188, 385 hazard and operability, 132 heat affected zone, 429
heat balance diagrams, 561 heat recovery steam generator, 421, 547–8 improving efficiency, 628–32 feedwater chemistry, 632 multiple pressure HRSG option 1, 630 multiple pressure HRSG option 2, 631 recent trends, 629–32 performance management and improvement, 606–34 effect of steam pressure on tube thickness, 609 fire tube boiler, 611 gas turbine, 611–18 typical gas data, 608 waste heat boiler classification, 607 waste heat boiler for fume incinerator, 610 pinch and approach points affecting HRSG size and steam generation, 618–21 K values, 619 logic for off-design HRSG performance calculations, 622 simulation, 621–8 examples of simulation analysis, 624–8 gas/temperature profile, 625 off-design fired case, 628 off-design part load, 626 hematite, 462 Herring–Nabarro diffusional creep, 583 high temperature cylinders, 519–24 design criteria, 519 general concerns, 519 high temperature alloys and their ASTM equivalents, 520 innovations, 522–4 9%CrMoVNbB moving blades, 524 increased creep strength and advanced materials, 523 modified 9%CrMo and 10% CrMoVNbN diaphragms, 522–3 prototype diaphragm in 10% CrMoVNbN alloy, 523 material selection, 520–2 bolting, 522 casing materials, 521 moving blades, 521 rotor materials, 520–1
© Woodhead Publishing Limited, 2011
Index stationary blades, 522 high-temperature heat exchanger construction, 582–97 flowing slag corrosion testing, 584– 9 ODS alloys, 589–97 UTRC tubes-in-a box HTHX design, 583 materials management and performance improvement of IFCC systems, 575–603 pilot-scale testing, 597–602 design of the EERC slagging furnace system, 597–9 HTHX testing during oxy-firing, 599–602 high-temperature hot corrosion, 362–3 high-temperature oxidation, 374–5 high-temperature vacuum brazing, 432– 6 blending, 435 cleaning, 434 heat treatment, 435 inspection, 436 lay up, 434–5 high velocity oxyfuel, 441–3, 647 high velocity oxygen fuel thermal spraying, 171 higher heating values, 9 horizontal silo burners, 128 hot corrosion, 361–3, 375 HR3C, 226 HS-188, 387 Huyghens principle, 69 impression creep, 196 in situ non-destructive techniques assessment, 58–61 boiler and HRSG components, 59 boiler inspection, 58–9 inner oxide thickness assessment methods in waterwall tubes, 59–60 material condition assessment methods, serviced steam headers and steam lines, 61 SH and RH tubes inspection, 60 steam headers and steam lines inspection, 60 indirectly fired combined cycle systems, 577–81 high-temperature heat exchanger, 575–603 coal gasification systems, 581–2
675
construction, 582–97 pilot-scale HTHX testing, 597–602 ultrasupercritical systems, 576 LCOE for coal-fired power systems employing CO2 recovery, 580 LCOE for standard pc-fired power systems, 580 UTRC HiPPS concept, 578 infrared, 51 inlet connection, 565 sleeve and piston ring, 564 Institute of Nuclear Researches, 498 integrated gasification combined cycle (IGCC) power plants availability analysis, 110–38 acid gas removal (AGR) and sulphur recovery availability issues, 125–7 ASU availability issues, 116–19 combined cycle availability issues, 127–9 gasification units availability issues, 120–5 existing plants summary, 129–32 different main units contribution to total downtimes, 130 IGCC availability developments, 131 main units downtime at Puertollano, 130 future trends, 135–8 availability experiences of Polk Power Station vs several studies, 137 integrated gasification structure, 111– 16 acid gas removal and sulphur recovery, 114–15 air separation unit, 111–13 combined cycle, 116 gasification unit, 113–14 IGCC basic structure, 112 RAM modelling forecast, 132–5 RAM studies results, 133 inter alia DICTRA (diffusion-controlled transformation), 262 International Energy Agency Clean Coal Centre, 20 International Organisation for Standardisation (ISO), 5 Irsching 4, 379 Ising model, 263
© Woodhead Publishing Limited, 2011
676
Index
‘J-series’ gas turbine, 381 Japanese R&D, 224 jetting damage, 288 Johnson–Mehl–Avrami equation, modified, 264 KA26-2 ICS Multi Shaft, 382 KA26-2 Multi Shaft 2 on 1, 382 KA26-1 Single Shaft 1 on 1, 382 Kachanov–Rabotnov, 189 Kawasaki Combined Cycle Thermal Power Station, 381 Larson–Miller parameter, 215, 346 laser cladding, 430 laser powder fusion welding, 430–2 laser vibrometry, 76 last stage blade, 550–1 exhaust loss vs annulus velocity, 550 Laves phase, 281 leak detection, 42–7 acoustic leak detection and monitoring, 44–5 boiler equipped with leak detection system, 46 leak detection and monitoring, microvibrations, 45–6 leak detection motivations and boilers monitoring, 42–3 damage causes, tubes in boilers and HRSGs, 50 NDT methods used on boiler and HRSG components, 59 leak-induced sound and vibration methods, 44 air-borne leak sensor, 45 mass flow balance methods, 43–4 requirements and monitoring methods, 43 signal processing and diagnostic criteria, 46–7 leak detection by acoustic monitoring, 47 levelised cost of electricity, 579 Lifshitz sliding, 186 ‘limit loading’, 550 linear cumulative damage, 405 linear damage summation, 286, 406 linear life fraction, 405 linear life fraction rule, 312, 319 linear polarisation resistance, 170 logistic creep strain prediction, 213 long term overheating, 284
long term service agreements, 333 low-cycle fatigue, 356, 367 low pressure plasma spray, 440–1 low temperature cylinders, 524–33 design criteria, 524 general concerns, 524–31 avoidance of flow accelerated corrosion/erosion–corrosion, 529–30 avoidance of stress corrosion cracking, 524–9 minimisation of water droplet erosion, 530–1 material selection, 531–3 bolting, 533 casing materials, 532 LP cylinder alloys and their ASTM equivalents, 532 moving blades, 532 rotor materials, 532 stationary blades, 533 low-temperature hot corrosion, 362 lower heating values, 9 MatCalc, 263 MCrAlY, 371 mechanical distress, 375–6 MesoPLasma, 101 metal coatings, 373 metallography, 188 M701G2 gas turbine, 380–1 Microdur 20, 652 microstructural degradation boiler steels, 222–66 characteristics of microstructure and property degradation processes, 248–60 material assessment and residual durability determination, 230–48 modelling degradation processes and their use, 261–6 preparation of classification system for material after operation, 260–1 steel development for power engineering, 223–30 gas turbine superalloys, 322–7 agglomeration and rafting of the γ– γ ´ microstructure, 325 cast blade (GTD-111) microstructure, 323 cross-section of turbine blade with severe irreversible, 327
© Woodhead Publishing Limited, 2011
Index ex-service blade (GTD-111) with coarse partially fused γ ´ particles, 324 ex-service cast turbine blade microstructure, 323 plate-like TCP phase precipitates below the coating, 326 single crystal blade microstructure, 324 whitish zone on crack surfaces, 326 Miner rule, 287, 319, 405 minimum creep rate, 308, 342 minor and trace metal analysis, 7–8 moisture, 6 molecular dynamics, 185 Monte Carlo method, 263 MS9001 H gas turbine, 380 multiaxial stress testing, 193–4 municipal solid wastes, 26 Nabarro–Herring creep, 185 neutral water treatment, 478 Ni-base superalloys, 460y¨ scale morphology and spallation, 473 Norton creep law, 343 Norton equation, 213 ‘nozzle guide vanes’, 388 ODS alloys, 589–97 bolted MA956, 591–2 tubing-elbow joint, 591 brazed MA754, 590–1 one-half of the cross-section MA754 elbow, 590 evaporative metal bonding, 595–7 backscattered electron image and X-ray map of EM joint, 596 creep rupture specimen, 595 transient liquid-phase bonding of MA956, 592–4 close-up of surface exposed from tensile break, 594 MA 956 pipe sections showing globules of bonding alloy, 592 MA 956 pipe sections using the heating schedule from DICTRA calculations, 593 MA956 pipe sections showing tensile break, 594 optical fibres thermometers, 54–5 optical microscopy, 169–70 optical techniques, 76, 100 Orowan equation, 182
677
overheating and heat exchange monitoring, 49–55 advanced fibre optic methods, temperature monitoring inside boiler, 54–5 infrared methods, temperature monitoring, 51 monitoring temperature motivations and boiler tubes heat exchange, 49–51 monitoring temperature profile, 52 permanently inserted thermocouples usage for SH and RH tubes temperature and heat exchange monitoring, 54 permanently inserted thermocouples usage for temperature and heat exchange monitoring, 51–3 overlay coatings, 371–2, 439–43 electro-co-deposition, 443 electron beam physical vapour deposition, 443 high velocity oxyfuel, 441–3 low pressure plasma spray, 440–1 vacuum plasma, 440 oxidation, 157–9 growth rate constants, 159 heat exchanger tube materials nominal compositions, 158 oxidation partial pressure, 460–3 chromia evaporation, 463 expected oxide phases, 461–2 oxidation-reduction potential monitoring techniques, 632 oxide dispersion strengthened alloys, 401–2 oxide spallation, 454 oxy-firing HTHX testing, 599–602 heat recovery data, 600 oxyfuel combustion, 480 oxygenated treatment, 477–8 Palmgren–Miner, 498, 510 Paris equation, 355 Parsons Engineering, 579–80 partial-arc admission, 542–4 particle size distribution, 8 percolation theory, 190 performance monitoring, 555–6 HP cylinder efficiency deterioration, 556
© Woodhead Publishing Limited, 2011
678
Index
phased-array ultrasonic (PAU) techniques, 69 photoluminescence piezospectroscopy, 91 physical scrubbers, 115 pinch point, 616 HRSG size and steam generation, 618–21 Poisson–Voronoi tesselation, 189 Polish Standards PN-75/H-84024 (1975), 230 power-law creep, 345 power-law relationship, 355 power plant future trends, 655–7 parent steel developments, 655–7 repair procedure, 657 risk of failure in existing plants, 655 weld metals of higher creep strength, 657 life management and performance improvement of steam turbine upgrade, 535–72 drivers, 537–9 installation, 563–9 installed generating capacity vs age, 536 mechanical design, 557–63 performance improvement, 547–57 product selection and specification, 539–46 welds and joints materials management and performance improvement, 635–57 applications of degradationprotection technologies, 646–7 chemical composition and creep rupture strength of ferritic steels, 637 dissimilar joints, 648–51 impact on power plant performance/life management, 647–8 influence of a brief hold at different austenitisation temperatures, 636 inspection and hardness testing, 651–2 materials selection and development, 636–8 repair, 653–5 requirements for creep rupture strength, 638 weld/joint degradation, 639–45
power plant boilers damage inspection and monitoring and integrity/life assessment issues in high Cr martensitic steels, 291– 301 atomic/scanning force microscopy for on-site cavitation damage assessment, 295 cracking and exfoliation of internal oxide scale, 292 detection of creep damage by ultrasonic velocity change method, 298 hardness monitoring as an integrity assessment, 301 methods for identification and quantitative description of microstructural parameters, 294 microstructure based integrity assessment, 293–5 optical and scanning electron microscopy for cavitation measurement, 295 portable SFM being used on a creep damaged P91 welded pipe, 295 potential drop as a damage monitoring technique, 298 potential drop ratio vs creep life fraction relationship, 301 SFM images acquired before and after chemical etching, 296 tube specimens A, B, and C, 299 ‘Type IV’ cracking in the finer region of P91 HAZ, 292 ultrasonic noise method for damage monitoring, 297 use of ASME P91 and P92 type high Cr martensitic steels, 291–3 UT velocity change results, 300 damage mechanisms, inspection and life assessment, 272–301 damage mechanisms and component failure, 283–90 creep, 283–5 creep–fatigue interaction, 286–8 effect of ageing on ferritic steels microstructure, 289 erosion, 288 erosive wear in power generation, 288 fatigue, 285–6 microstructural degradation,
© Woodhead Publishing Limited, 2011
Index particle coarsening and spheroidisation, 289–90 steam oxidation, 290 materials, metallurgy and microstructure, 274–82 basic characteristics of minor phases in 9–12%Cr steels, 279 chemical compositions of 9–12%Cr steels, 278 low alloy steels used for tubes and pipes, 276 materials for boiler tubing and headers, 274 materials for steam pipework, 274– 5 metallurgy of 9–12%Cr martensitic steels, 276–81 metallurgy of austenitic stainless steels, 281–2 metallurgy of low alloy ferritic steels, 275–6 P91 tempered martensite microstructure, 276–7, 280–1 power plant components condition monitoring and assessment, 38–102 gas turbines condition monitoring, 74–83 in situ assessment, gas turbine hot parts by non-destructive techniques, 84–95 information sources and advice, 101–2 plant condition monitoring levels, 40 remote monitoring solutions, 95–8 steam turbines and generators, 61– 74 future trends, 98–101 condition assessment and monitoring policies, 98–9 technology innovations examples, 99–101 monitoring boiler and heat recovery steam generators, 41–61 corrosion monitoring, 47–9 in situ non-destructive techniques assessment, 58–61 leak detection, 42–7 overheating and heat exchange monitoring, 49–55 plant condition monitoring levels
679
with typical sensors and delivered information, 42 structural integrity monitoring, headers and steam lines, 55–8 power plant fuel flexibility alternative fuel usage in combustion power plants and technology application to improve fuel flexibility, 30–3 co-firing fuels, 32–3 fuel substitution, 31–2 solid fuel composition, 3–34 alternative fuel usage in combustion power plants and technology application to improve fuel flexibility, 30–3 fuel chemistry and composition, 5– 30 future trends, 33–4 predictive modelling, 211–15 creep models summary, 215–16 creep processes, 211–12 phenomenological models based on creep curve self-similarity, 213–14 physical description, 212 simplified physically based models, 212–13 time–temperature parameters, 214–15 pressure swing adsorption, 115 Propulsion Instrumentation Working Group (PIWG), 81 protective coatings, 368–77 diffusion coatings, 369–70 overlay coatings, 370–2 thermal barrier coatings, 372–3 pulsed thermography, 85–6 ‘pure H2O’, 461 RA-333, 387 radiant syngas cooler, 123 ‘rainflow method’, 499 Raman spectroscopy, 91 Rankine cycle, 457, 577 raw gas cooling, 113–14 raw gas pre-cleaning, 124–5 dry filtration, 124 wet scrubbing, 125 raw gas pre-treatment, 114 refused derived fuels, 28 reheater, 41 relaxation, 314 reliability factor, 134 remnant lifetime monitoring, 498–514
© Woodhead Publishing Limited, 2011
680
Index
analysis of obligatory boiler standards and regulations, 500–2 creep, 499–500 creep life prediction for cylindrical element, 499 creep behaviour and creep strains measurement, 507–10 circumferential strains, 508 circumferential stressed caused by creep at points P1 and P51, 509 equivalent total stress history at point P1, 509 low cyclic fatigue, 498–9 fatigue investigation chart, 498 measurement of creep strains, 510–14 analyzed pipelines data, 513 creep curves of pipelines, 513–14 general pipeline arrangement, 511 pins for diametral measurements, 512 schematic creep curve, 514 temperature and stress distributions, 502–7 axial thermal stress transients, 506 calculated allowable stresses and heating and cooling rates, 503 circumferential thermal stress transients, 506 division of the half cross-section of the horizontal cylindrical element, 504 equivalent stress transients on outlet head’s inner surface, 507 measured pressure and temperature transients, 505 monitored pressure elements of a power boiler, 503 temperature transients on the outlet header’s inner surface, 505 repair welding, 653–5 residual life RCL, 231 Robinsons rule, 287, 312 Rolls-Royce, 381 root mean square, 46 S–N method, 353 safe-life approach, 349–50 safety integrity level, 132 Sanicro 25, 226 SCR units, 629 secondary creep, 308, 342, 343 selective catalytic reduction, 116 Sherby–Dorn parameter, 214
short term overheating, 284 Siemens Power Generation, 379 silicon, 370 silicon carbide high-temperature sensor, 100–1 single-crystal directionally solidified (SCDS) superalloy, 366 single-crystal (SC) castings, 400 slag corrosion testing, 584–9 aluminium content in MA956, 589 chromium content in MA754, 588 DSAF system, 585 MA754 and MA956 tube sections, 586 SEM image of MA754 tube cross section, 587 SEM image of MA956 tube cross section, 587 slag removal, 114 slagging/fouling/corrosion systems, 33 slagging furnace system, 597–9 bare-tube HTHX, 599 schematics of system used in the testing, 598 small angle neutron scattering, 188 small punch creep tests, 195–6 small punch (SP) method, 61 sodium sulphate, 362–3 solid fuel composition power plant fuel flexibility, 3–34 alternative fuel usage, combustion power plants and technology application, fuel flexibility improvement, 30–3 fuel chemistry and composition, 5– 30 future trends, 33–4 solid recovered fuels, 28 solid-state diffusion, 376 solution precursor plasma spray (SPPS) process, 403 spallation, 376–7 standard creep testing methods, 192–3 STANJAN computer code, 466 steady state creep, 182, 342, 344 steam admission control, 542–4 FAA nine stage impulse inner module retrofit, 544 FAA vs PAA heat rate, 545 PAA eight stage impulse inner module retrofit, 544 steam boiler
© Woodhead Publishing Limited, 2011
Index component loading, monitoring and life assessment, 490–515 remnant lifetime of pressure components monitoring, 498–514 thermal and strength analysis for start-up and shut-down operations, 491–7 steam cycles, 547 expansion line fossil-fired and nuclear cycles, 548 steam extractions, 566 steam oxidation boiler and turbine environments, 453– 89 future trends, 478–80 advanced ultra-supercritical, 478 co-firing with biomass, 479–80 major international research and development efforts, 479 oxyfuel combustion, 480 management, 476–8 materials requirements, 459–60 bolts, 460 high pressure and intermediate pressure rotors, 459 rotating blades, 460 steam piping and headers, 459 superheater and reheater tubing, 459 waterwall tubing, 459 operations, 477–8 steam cycle chemistry, 477–8 outage management, 476 ultrasonic testing, 476 visual examination, 476 oxidation thermodynamics and kinetics, 460–7 kinetics, 463–7 oxidation data in steam from mass change measurements, 466 oxidation data in steam from scale thickness measurements, 466 oxidation partial pressure, 460–3 phase stability of iron as function of temperature and oxygen partial pressure, 462 steam oxidation kinetics based on mass change measurements, 465 steam oxidation kinetics based on scale thickness measurements, 465 thermodynamics, 460
681
scale morphology and spallation, 467–75 Armitt diagram, 474 austenitic steels, 470–2 coefficients of thermal expansion for relevant oxides and alloys, 475 cross-section of scale formed on Inconel 740 after 4000 hours in steam, 473 ferritic steels, 467–70 Ni-base superalloys, 473 spallation models, 474 spallation superimposed on to the Armitt et al diagram, 475 steam boiler and turbine environments, 454–60 boiler alloys, 460 Rankine cycle for subcritical boiler, 457 Rankine cycle for supercritical boiler, 458 simplified water path in large supercritical utility boiler, 456 steam cycle, 455–9 water properties, 455 steam side, 453 steam turbine and generators, 61–74 damage locations and causes, steam turbine blades, 67 damage locations and causes, steam turbine diaphragms and nozzle boxes, 68 damage locations and causes, steam turbine rotors, 66 damage locations and causes, steam turbine valves and steam chests, 68 damage mechanisms, 65–8 main steam turbine components, structural integrity assessment, 70–1 miniaturised tests, rotor material assessment, 71–4 non-destructive examination, 68–70 on-line vibration and dynamic condition monitoring, 61–5 procedure flowchart small punch test result and other nondestructive measurements, 74 rotor vibration amplitude versus rotation speed, 63
© Woodhead Publishing Limited, 2011
682
Index
sample extraction from serviced steam turbine rotor for small punch testing, rotor aspect and actual size of extracted material, 73 two couples of velocity and acceleration sensors, 63 vibration orbits along the shaft line, 64 drivers, 537–9 availability, reliability and risk management, 538 extended power uprates, 539 life extension, 538 net income vs time for power plants before and after retrofit, 538 performance improvement, 539 installation, 563–9 alignment, 566–7 design for installation, 563–6 planning and management, 567–9 potential durations, 569 technical field advice, 567 materials selection, life management and performance improvement, 518–33 high temperature cylinders, 519–24 low temperature cylinders, 524–33 mechanical design, 557–63 concept design process, 561–3 design life, 560 flow diagram concept design process, 562 maintenance requirement, 560–1 measurement process, 557–60 performance improvement, 547–57 capacity change, 549–50 control philosophy, 551–2 cylinder efficiency, 557 effect of LP blade exhaust area, 550–1 efficiency improvement, 549 efficiency vs load for FAA and PAA control, 552 guarantee alternatives for fossilfired cycles, 563 guarantee alternatives for nuclear cycles, 563 guarantees, 552–3 monitoring, 555–6 plant conversions, 547–8 steam cycles, 547 turbine architecture, 548
typical benefits, 556–7 verification performance improvement, 554–5 product selection and specification, 539–46 cold end considerations, 546 design conditions and life, 543–6 hot end considerations, 546 improvement potential, 539–40 steam admission control, 542–4 types, 540–2 HP inner module impulse turbine retrofit with disc and diaphragm construction, 541 HP inner module impulse turbine retrofit with drum rotor construction, 541 upgrades for power plant life management and performance improvement, 535–72 ‘Steel B’, 524 strain life, 353 strain-range partitioning (SRP) model, 407–9 stress concentrations, 193–4 stress corrosion cracking, 524–9 effect of yield strength on initiation of SCC, 526 flexibility of welded rotor, 529 threshold stress approach, 527 threshold stress approach for SCC in a normal plant stream environment, 528 stress distributions, 502–7 stress intensity factor range, 354–5 stress relaxation testing, 193 structural integrity monitoring headers and steam lines, 55–8 acoustic emission methods, 56–7 creep damage monitoring, 57–8 structural monitoring, steam header by acoustic emission, 57 structure-borne noise, 45–6 sulphur recovery, 114–15 availability issues, 127 superalloys creep–fatigue interaction, 359–60 partitioning of the strain range into different component strains, 360 gas turbine, 332–68 firing temperature trend with blade material improvement, 334 nominal composition of IGT cast
© Woodhead Publishing Limited, 2011
Index Co-base and Ni-base superalloys, 337 role of alloying elements, 336 metallurgical behaviour, 338–59 allowance for factors that affect fatigue, 352 creep-rupture schematic curve, 341 deformation mechanism map, 345 effects of mean stress on alternating stress amplitude, 352 fatigue crack initiation behaviour of various steels, 355 Larsen–Miller parameter for various types of blades, 347 room-temperature mechanical properties for various materials, 340 S–N diagram trends, 351 stress and temperature dependence of the creep rate, 345 time-dependent deformation, 342 time to rupture as function of steady-state creep rate, 344 variation of fatigue-crack-growth rate with alternating stress intensity, 356 oxidation and hot corrosion, 361–3 corrosion rate variation with temperature, 362 properties, 363–7 physical properties of cast nickelbase and cobalt base alloys, 364 physical properties of superalloy base elements, 364 Super304H, 226 superheater, 41 suspension plasma spray (SPS), 403 T-pipe thermal and strength analysis, 491–7 division into finite elements, 492 equivalent thermal and total stress histories during cooling operation, 497 equivalent thermal and total stress histories during heating operation, 496 first method of boiler start-up and shut-down operations, 493 13HMF steel, 492 second method of boiler start-up and shut-down operations, 494 technical field advisors, 567
683
temperature distributions, 502–7 tertiary creep, 342 thermal barrier coatings (TBC), 368, 372–3, 402–3 thermal fatigue, 367 thermally grown oxide, 84 Thermo-Calc system, 262–3 thermocouples, 51–2 thermomechanical fatigue, 356–7 thermomechanical fatigue cracking, 377 thermophoresis, 155–6 three-dimensional laser scanner, 558 illustration, 559 time–temperature parameters, 214 toughness, 347 TP347HFG, 226 transient liquid-phase bonding, 592–4 TRD-508, 233 trend monitoring systems, 75 TTT-type (time, temperature, transformation) diagrams, 264 tungsten inert gas (TIG) welding, 429– 30 turbine inlet, 81 ultimate analysis, 7 ultrasonics, 188 ultrasupercritical systems, 576 UTRC HiPPS project, 582–4 vacuum casting, 400 vacuum plasma, 440 van Krevelen diagram, 15 verification test see acceptance test volatile matter, 6 water droplet erosion, 530–1 hardened zone at leading edges of moving blades, 531 water tube boilers, 609 weld/joint degradation, 639–45 failure locations, 639–40 cracking classification, 639 failure locations for cross-weld samples, 640 schematic representation of HAZ subzones, 641 increasing awareness of low-weld creep strength, 640–5 creep rupture data, 642 grade 2 creep rupture weld strength factors, 643 predicted weld strength factor, 642
© Woodhead Publishing Limited, 2011
684
Index
stress vs WSF for grade 91 steel, 644 weld strength factor for weldments in a range of steels, 645 low cross-weld creep strength, 640–5 ‘weld strength reduction factor’, 643 wet scrubbing, 125
wu¨stite (FeO), 461 X-ray diffraction, 169–70 X20CrMoV12.1 steel, 224, 234 yttria-stabilised zirconia, 88
© Woodhead Publishing Limited, 2011