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The new low-chromium steel grades 23 and 24 are candidate materials for components of the new power and petrochemical plants, and for the refurbishment and re-powering of older plants. The mechanical and creep properties of both grades are significantly better than the parent grade 22, but long-term creep performance, microstructural evolu-tion, welding characteristics and other properties were not fully defined and assessed. It was also important to improve knowledge of microstructural evolution in order to verify the mechanical behaviour after long-term service.

The consortium has produced trial components by industrial process routes for both grades, but the activities have been focused mainly on grade 23, for commercial rea-sons, and on grade 24 for comparison. New consumables for welding have been devel-oped and tested. Creep test programmes for base material and welded joints, including long-term tests, have been carried out, and some tests will continue beyond the end of the project. The data acquired will also be incorporated in the creep database of the European Creep Collaborative Committee and will be used in the coming assessments for EN standards.

The parallel aim of the project was piping integrity assessment under realistic loading conditions by combined thermal and hydraulic system analysis and stress analysis using the data generated during the project. This work has shown that a P23 pipework system will be more durable than an equivalent CMV system providing that good operational practice is maintained, thereby minimising the risks of severe operational transients.

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Applications of advanced low-alloy steels for new

high-temperature components

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European Commission

Research Fund for Coal and SteelApplications of advanced low-alloy steels for

new high-temperature components

A. Di Gianfrancesco, D. Venditti (1), D. J. Allen, A. Morris (2), S. Caminada (3), S. Pillot (4), M. M. Rodriguez (5), V. Friedman, P. von Hartrott, D. Siegele (6),

S. Holmström, J. Rantala, J. Salonen, P. Nevasmaa, K. Calonius, P. Junninen (7)

(1) CSM — Via di Castel Romano, 100, I-00128 Rome(2) E.ON UK — Ratcliffe-on-Soar, Nottingham NG11 0EE, United Kingdom

(3) Dalmine — Piazza Caduti 6 Luglio 1944, 1, I-24044 Dalmine(4) Industeel — 56, rue Clémenceau, BP 19, F-71202 Le Creusot

(5) ISQ — Av. Prof. Dr Cavaco Silva, 33, Taguspark, PT-2780-994 Porto Salvo(6) Fraunhofer-Institut für Werkstoffmechanik (IWM) — Wöhlerstraße 11, D-79108 Freiburg

(7) VTT — PO Box 1000, FI-02044 VTT

Contract No RFSR-CT-2003-00037 1 September 2003 to 28 February 2007

Final report

Directorate-General for Research

2009 EUR 23598 EN

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A great deal of additional information on the European Union is available on the Internet. It can be accessed through the Europa server (http://europa.eu). Cataloguing data can be found at the end of this publication. Luxembourg: Office for Official Publications of the European Communities, 2009 ISBN 978-92-79-10006-2 ISSN 1018-5593 © European Communities, 2009 Reproduction is authorised provided the source is acknowledged. Printed in Luxembourg PRINTED ON WHITE CHLORINE-FREE PAPER

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Table of Contents

Summary 4 Scientific and technical description of the results 14 Objectives of the project 14 Comparison of initially planned activities and work accomplished 15 Description of activities and discussion 17

o WP1: Characterisation of the ‘as-received’ steels 17 o WP2: Development of weldments 29 o WP3: Mechanical data assessment 48 o WP4: High temperature design and assessment 76 o WP5: High temperature design and assessment 90 o Conclusions 110 o Exploitation and impact of the research results 112

List of figures and tables 113 List of References 116 Appendix 1: Welding procedure specification (WPS) 118 Appendix 2: LM and SEM microstructural investigation of aged specimens 131 Technical Annex 134

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Summary

Introduction The new advanced low Chromium Grades 23 and 24 steels are and will be strong candidates for new power and petrochemical plant construction and for the eventual large-scale replacement of steam pipework on existing power plant. These new grades are improved version of the grade 22 (2¼Cr1Mo) developed in the years ’60 and applied in worldwide the power and petrochemical plants for large amount of components: tubes, pipes, cast, forged. In the years ’80-90 several effort have been devoted to increase the plants performances and therefore materials with enhanced performances were requested. The steelmaker metallurgists and the material researchers started to develop new chemical compositions on the base of existing steels with the addition of elements able to give strengthening by carbides precipitation (V, W, Nb, Ti). The main differences are described in the following table based on ASTM A213 Standard.

For both the grade V was added in the chemical composition and the other elements were respectively: - for Grade 23: W and Nb with lower amount of Mo, with a similar approach used for high alloyed chromium grade 92 respect to grade 91. - for Grade 24: more Cr than grade 22, with the addition of B and Ti. These addition are able to guarantee an increase of creep behaviour with a possible service temperature up to 560-570°C that means 30-50°C more than grade 22. Otherwise the increase of carbides former elements increase the problems in the welded joints. The following figures shown an overview of main components of power and petrochemical plant:

- small diameter and thin wall thickness tubes (a) for waterwalls of boiler (b) and heat exchangers,

- larger diameter and thick wall pipes for headers (c) and steam lines (d) - very large diameter and heavy wall forged components (e) for pressure vessels of

petrochemical power plant reactors (f).

a b c

Grade C Mn P S Si Cr Mo W Nb V B Other min 0.05 0.30 - - - 1.90 0.87 - - - - 22 max 0.15 0.60 0.025 0.025 0.50 2.60 1.13 - - - - -

min 0.04 0.10 - - - 1.90 0.05 1.45 0.02 0.20 0.0005 23 max 0.10 0.60 0.030 0.010 0.50 2.60 0.30 1.75 0.08 0.30 0.0060 N: 0.03 max

min 0.05 0.30 - - 0.15 2.20 0.70 - - 0.20 0.0015 24 max 0.10 0.70 0.020 0.010 0.45 2.60 1.10 - - 0.30 0.0070 N: 0.012 max Ti 0,06-0,10

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d e f The steam lines are one of the main critical component of a power plant due to:

- high temperature steam and high pressure, from headers to the turbine (blue area); - low cycle fatigue as consequence of the stress variations during cyclic operating

condition of the power plant; - the thermal expansion as consequence of the temperature variations during cyclic

operating condition of the power plant generate an additional thermal fatigues stress.

Steam Turbines

Boiler

Steam LineSuperHeater

The additional alloying elements give one increase of the cost of the row materials, but the cost of the final components are not detrimental for the use of these new materials, because it has to be take in account the increase of material performances and consequently the reduction of the wall thickness, as well as, the improvement of the plant efficiency. The main aims of the project were the following:

- Mechanical and microstructural assessment of base materials, - Development of weld material and welding procedure to avoid the “bore cracking”

phenomenon, - Mechanical and microstructural assessment of similar welding, - High temperature design and assessment of welded components including welded

pressure vessels and pipework,

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- Microstructural modelling to predict changes as a function of time and temperature in service operation,

- Piping integrity assessment under realistic loading conditions by combined thermal hydraulic system analysis and stress analysis tools.

Major development needs were:

- to establish knowledge of the base material in term of mechanical and microstructural evolution during the service life,

- welding consumables with proven high temperature properties needed to be developed and validated by creep rupture testing, microstructural and mechanical properties assessment,

- to show that new steels and welds have improved resistance to the in-service bore cracking phenomenon recently identified in existing CrMoV steel steam pipework.

These required comparison between the material properties after the manufacturing heat treatment and after additional simulated service aging, the behaviour of the welded joint and the comparison of low cycle fatigue crack initiation and growth testing, together with creep strain rate, creep rupture and creep crack growth testing, on new and original pipework steels and weldments. These tests have been carried out on experimental materials produced in full size dimensions by industrial process routes. It has also been necessary to model the effects of plant operational conditions and service loadings on the integrity of welded pipework geometries and configurations using combined thermal hydraulic system analysis and structural analysis. This modelling approach has been successfully defined and verified by the laboratory full scale simulation tests. In parallel, microstructural modelling and characterisation techniques have been employed to predict microstructural changes during long term operation of high temperature plant. A large number of tests have been carried out, and an extensive database of information for steelmakers (tubes, pipes, plates, welding consumables) and end users has been generated. This approach and the results obtained will be necessary for component design, manufacture, plant construction, and maintenance and inspection of power and petrochemical plant. The Consortium believes that a major increase of the knowledge of the new material has been acquired, but that not all aspects have been completely clarified, such as the development of a weld metal with sufficient creep ductility. The partners have gained much knowledge and experience in the usability and life management aspects of these materials. This knowledge base can successfully be applied in other projects to come. The acquired understanding of the effect of the alloying elements in the consumables could be used for further consumable development and in predicting the long term microstructural evolution. The developed material models will also be the base for high temperature component simulation. The results obtained in this project will make a small but relevant contribution to increasing plant safety and efficiency, and thus towards achieving the Kyoto Protocol target for Europe of an 8% reduction of CO2 emission by 2010. The main project deliverables for the Consortium partners are: - A database on the materials properties of grade 23 and 24 steels and welded joints for high

temperature component design and structural integrity assessment. This includes strength,

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toughness, fatigue crack initiation and growth data, and minimum creep rate, creep rupture, and creep crack growth rate data;

- New welding consumables for P23 steels designed to avoid the risk of premature high

temperature failure due to low creep ductility; - Weld creep data assessment to improve capability for long term design and operation of

advanced high temperature plant using high and low alloy ferritic steels; - Comparative high temperature fatigue data to enable optimum steel design and selection for

the prevention of in-service cracking in power plant which is required to operate flexibly and is hence subject to transient loading;

- Modelling data on the effects of cyclic plant operation on component structural integrity and

the consequential requirements for design, materials selection and integrity assessment; - Microstructural assessment of materials, weld metals and heat affected zone structures to

study the evolution of service aged and/or tested material and predict the microstructural changes that will occur in high temperature steels. In conjunction with the materials properties database, component modelling and assessment, this will provide a sound basis for plant design, operation and maintenance to prevent the in-service failure of welded components.

The project was organized in the following work-packages: WP1: Characterisation of the ‘as-received’ steels (WP Leader: CSM) WP2: Development of weldments (WP Leader: ISQ) WP3: Mechanical data assessment (WP Leader: VTT) WP4: Microstructural assessment (WP Leader: CSM) WP5: High temperature design and assessment (WP Leader: E.On (ex Powergen)) WP6: Project management (WP Leader: CSM) Each Work Package (WP) leader has successfully achieved the specific aim of their WP, and a good network has developed not only between the representatives of each organisation but also other researchers and technicians involved in the specific tasks. This network will be used by the partners for further cooperation in the future. The exchanges of specific information and test results will continue during coming years due to the long term tests still running for a considerable amount of time. Two meetings per year have been held, located on a rotating basis at the specific laboratories and industrial facilities of all the partners, in order to improve the relationships between the different partners involved in the main activity carried out in each semester. The summary of the main activities carried out for each work-package and the relevant results obtained are described below. 3.1 WP1: Characterisation of the “as-received “Steels The main objectives of the WP1 were: • The selection of test materials and plate/tube geometries, • The microstructural characterisation of base material in normalised and tempered condition, • The mechanical properties of base material including strength and toughness.

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The main conclusions of the work package 1 are following summarised. The experimental materials for testing have been produced on industrial scale and successfully qualified by conventional mechanical testing and metallographic characterisation. Grade 23 has been selected for the main test programme. 3.2 WP2: Development of weldments

The main objectives of WP2 were: • The development of welding consumables, • The optimisation of welding procedure (welding technology and parameters), • The qualification and simulation (using Gleeble testing) of welded joints • The determination of mechanical properties of welded joints. The long term performance of similar welded joints in the creep resistant steels is often life-limiting in the design and operation of high temperature power and process plant. For the new low alloy steels, long term creep data on welded joints are not currently available, and the microstructural evolution of welds and base material is not well understood. It was necessary to investigate these factors to develop design data, welding consumables and welding procedures to minimise the risks of plant service failures at welds. Two specific development requirements were addressed in the proposed project. First, problems of poor weld creep ductility have been reported. Welding consumables with proven high temperature properties therefore need to be developed and validated by creep rupture testing, microstructural and mechanical properties assessment. Secondly, the new steels and welds must be shown to have improved resistance to the in-service “bore cracking” phenomenon now identified in existing UK CrMoV steam pipework. This requires comparative low cycle fatigue crack initiation and growth testing, together with creep strain rate, creep rupture and creep crack growth testing, on new and original pipework steels and weldments. The main conclusions of the work package 2 are following summarised:

♦ Welding procedures have been developed and optimised for manufacture of thick section P23 components. An extensive range of weld metal compositions have been investigated and a comprehensive series of full scale test specimens have been produced.

♦ The overmatching, highly alloyed “state-of-the-art” P23 weld metal B323B, selected for its high creep strength and used for the manufacture of long term high temperature test specimens, proved to be a poor choice. This creep-brittle weld metal was found to be liable to reheat cracking during PWHT of full scale thick section butt welds. The occasional presence of pre-existing weld reheat defects in the test specimens therefore led to a pattern of inconsistent and unreliable behaviour.

♦ B323B was deliberately chosen as a test of the viability of the high-strength formulation. The negative result is valuable, in that it clarifies the pitfalls that can occur and indicates a compositional range to be avoided.

♦ Parallel work using Gleeble weld thermal simulation together with the BWI tensile reheat cracking test showed that the P23 weld HAZ is susceptible to reheat cracking, while confirming that B323B weld metal is extremely susceptible.

♦ Later trials showed that alternative compositions, involving slight deviations from the P23 specification, could be developed to produce substantially less creep-brittle weld metals. These successfully survived PWHT without reheat cracking.

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♦ The results indicate that the Monkman-Grant creep ductility parameter M provides a reasonable guide to P23 weld metal performance, and that M values above 1% should generally be adequate.

♦ Improved weld metals, with M values of the order of 2%, can now be formulated. These should be subjected to long term testing in future work.

♦ P23 weld metals nevertheless show intrinsically poorer creep ductility than P23 parent materials. If adequate ductility is to be achieved, some sacrifice in terms of weld metal creep strength appears to be unavoidable. The microstructural reasons for this are not yet clear.

♦ The risks of reheat cracking in P23 weld metals and heat-affected zones can thus be reduced by careful selection of consumables and welding procedures. However, the inherent susceptibility of the P23 composition is a disincentive to its selection for thick section power plant components.

3.3 WP 3: High temperature testing and data assessment The main objectives of WP3 were: • The provision of high temperature mechanical data assessment for base material, • The provision of high temperature mechanical data assessment for welded joint, • The determination of weld stress reduction factors The output of WP3 is towards WP5 High temperature design and assessment. The main objective of AloAS WP3 is to assess the mechanical high temperature behaviour of the new heat resistant low alloy steels T/P23 and T/P24. The material properties for high temperature service are determined mainly by destructive material testing, such as tensile, creep and fracture toughness testing. Special effort is also put into determining the properties of weld metals (consumables) and the cross welds of the project steels. The material data is assessed with state-of-the-art assessment models and tools. The resulting models are then to be applied in WP5 for selected cases. The mechanical testing, modelling and simulation work of the AloAS project has produced a sound basis for life management of P23/24 steels and the development of their weldments. The collation of mechanical property data has been finalised and the final assessments have been performed. However some creep tests are still running and will be continued by the partners in order to increase the database. The results will be reported separately in international conferences and scientific publications. The main observations from the WP3 subtasks are presented in Table 3.3.1. Weld creep data were also produced in subtask WP2. The materials testing programme has been performed both at moderate (550-600°C) and high temperature regimes (above 600°C). The higher temperatures tests were taken into the testing programme due to the high stress levels needed for sensible durations in creep testing at the intermediate temperatures closer to plant service. At high temperatures the stress levels become relevant for plant operation conditions. The realistic plant stresses are estimated from design codes, based on extrapolated creep rupture data with a safety factor. For T23 steel tube, oxidation may limit application to about 575°C, but for P23, higher temperatures could be viable. The recent ASME Code Case 2199-3, 2006 0 interpolated design stresses are 70MPa and 53MPa at 575°C and 600°C respectively. It was concluded that the creep test programme should include test temperatures enabling data to extend down to around 50–70 MPa. Modelling of welded components is also required to analyse the effects of geometry, system support and static and transient plant loading on performance in service. This is necessary to

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determine the design advantages of new materials by properly assessing the interactions between materials properties and plant performance. For example, materials that are stronger in creep can be designed with reduced thickness, thus lowering thermal gradients in service conditions, and thereby improving materials performance under the thermal fatigue loadings that arise in cyclic and flexible operation of power plant. The main conclusions of the work package 3 are following summarised. A comprehensive high temperature materials database has been generated on P23 steels and weldments. These data provide support for high temperature power plant design with P23 steel. Background data on factors such as weld performance and defect development under plant cycling conditions will also support the lifetime management of operational power plant. The results show how important it is to have a good creep ductility for the weld metal of the newer high strength low alloy steels like P23 and P24, in addition to promising creep strength in preliminary qualification. The reasons are related to mechanisms that lead to strain localisation and cracking towards longer term creep exposure, and emphasise the importance of ductility criteria in short term testing and initial selection of candidate compositions. The results have also established an excellent basis for further development of suitable compositions that would perform better, and fundamental understanding on the mechanisms of the creep behaviour of multibead weld metal for high strength low alloy steels like P23 and P24. Further development of the weld metal is clearly needed, and there are new avenues to explore for the purpose. For this class of steels, the project has provided new tools and understanding for faster evaluation and further improvement of high performance welds. 3.4 WP4: Microstructural Assessment The main objectives of the WP4 were: • The evaluation of the microstructural evolution of the base material and welded joint, • The modelling of microstructural evolution. The output is towards WP5 High temperature design and assessment In parallel, microstructural modelling and characterisation techniques have been employed to predict the microstructural changes that will occur in steel components operating in high temperature plant components. This will be used to suggest suitable combinations of time, temperature and possibly stress that could be used to accelerate the ageing process in such a way as to provide material, within the timescales of the project, which could be used to determine the important mechanical properties pertaining to defect development. Because of the requirement for ever more efficient coal fired power plant and petrochemical plant, the demands on the properties of the steels from which the plant is constructed are also increasing. This has resulted in the development of more highly alloyed steels to meet the needs of improved creep and corrosion resistance. A consequence of such developments is steels with less stable microstructures than those currently used. For new plant, the changes in microstructure and properties can be expected to be of little significance. However, as the plant becomes older, service induced microstructural degradation may have significant effects on defect initiation and propagation rates and critical defect sizes. The main objective of WP4 is the evaluation of the microstructural evolution of the base material and welded joint for usability and life management planning for high temperature applications. The intended target for the characterisation is power plant components such as steam lines, water

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walls, etc. The knowledge acquired from aging, material property testing and the microstructural investigation is then to be utilised for modelling. The main conclusions of the work package 4 are following summarised. The assessment of the microstructural evolution of both grades 23 and 24 has been performed by aging tests carried out in the range of temperatures 550-675°C up to 10.000 hours. For both grades, the results obtained in term of microstructural observation seem quite well in agreement with the phases predicted at equilibrium by the JMatPro thermodynamic tool. However due to differences in precipitation and coarsening kinetics also metastable phases may develop during ageing. In grade 23, all the metastable types of precipitates like MC, M23C6 and M7C3 show a tendency to change their size distributions and to form the M6C carbide, which is the more stable phase at equilibrium. In grade 24, the situation is a little bit different because at equilibrium the M23C6 carbide is also a stable phase with a faster growth then M6C. In addition to changes in carbide structure increasing thermal exposure causes gradual coarsening in the lath structure. Obviously, during the aging times performed, the true equilibrium stage has not been reached. However, the aging tests at 675°C for 10.000 hours may approximately be equated, on a time-temperature parameter basis, to an exposure time of 200 years at 575°C. This is the maximum proposed service temperature for these low Cr grades, due to their oxidation rate. Regarding the variation of the mechanical properties during the aging at 550°C up to 10000 hours, the tensile properties at RT (YS and UTS) are practically not affected by exposure. However, the tensile properties measured at 550°C and 600°C show a reduction of about 200MPa in each case. Different behaviour is observed for the impact values. Grade 23 appears to be much better then Grade 24 after heat treatment and also after aging: the Grade 24 material reaches values lower then 50 Joule with brittle fracture. This could be related to the larger dimensions of the precipitates that are present in the Grade 24 steels after aging. The availability of more aged specimens from crept samples which are still running in this project, and in other EU programs such as COST 536, will be very useful to obtain more information on the microstructural evolution of both Grade 23 and 24. 3.5 Work Package 5: High Temperature Design and Assessment

This work package provides an assessment of the durability of proposed new welding consumables for P23 and P24 steels. Where possible this assessment takes into account pragmatic experiences on UK plant with CMV material specified systems. The scope of this assessment is shaped by recent UK experiences with thermal ‘bore cracking’ on CMV pipework systems, which in 2001 was originally confirmed as a threat to pipework integrity in the UK. During early June 2001, a new type of main steam pipe weld cracking was found on a UK station. The cracking initiated from the weld root or adjacent machined recess corner and grew radially outwards, in the worst cases, to a roughly uniform depth in excess of 20mm, fully circumferentially around the bore. The cracking has been found in CMV main steam pipework, which is of nominal 240mm bore and 60mm wall thickness. Subsequent metallurgical investigations suggested that the crack growth mechanism occurs initially due to thermal fatigue,

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with cracks initiated relatively early in the plant life. The pipe girth welds had been inspected during previous statutory outages, however this particular type of weld defect was undetected, despite its size. This was attributed to the nature of the cracks, which meant that they were difficult to find using normal CMV inspection procedures, hence enhanced procedures for detection and sizing were developed. On the initial detection of these defects an extended inspection campaign was initiated across all UK power plants. This revealed many instances of bore cracking, which necessitated the removal of affected welds or in some instances continued operation underpinned by a safety case. The instances of weld bore cracking tended to mainly occur in clusters either at the top of the main steam line, adjacent to the main boiler stop valves, or further downstream adjacent to the steam chests. In some stations the cracking was localised, in others it occurred at several welds along the main steam line. Crack growth rates of 2mm per year due to thermal fatigue are not uncommon. Of more significant concern are cracks that have already propagated a significant way through the pipe wall, which are the subjected to a high risk of accelerated crack propagation to creep crack growth.

At the time of writing, main steam bore cracking still occurs in the UK, but is now considered a ‘managed’ issue, with strategies in place to manage the integrity of affected plant with defects of various sizes. However, the experience with main steam bore cracking on UK main steam CMV pipework has again emphasised the following important issues:

• Successful outage inspection and return to service requires knowledge of expected defects, hence an understanding of the risk of defect initiation and propagation is a prerequisite.

• Generating utilities will not necessarily operate plant in accordance with design intent, particularly if the plant is operated in a commercial environment.

• Safety is paramount in the eyes of the plant owner. • There is no substitute for real service experience, and it is essential that such

experiences are used to shape the requirements for the development of future materials, and construction/repair methods.

• Even considering the extensive service experience gained with CMV pipework and weldments; there are many aspects associated with ‘integrity’ management that are still being addressed as emerging issues.

Hence, the above experience has emphasised the need within the scope of this project to undertake some assessment of the likely durability of the proposed weldments and to compare this against plant experiences with CMV pipework installations. It is within this framework that the tasks in this work package have been undertaken; and they hence comprise a mixture of analytical and pragmatic assessments utilising service experience.

The work package has been split into three specific areas, described in the following, which are intended to provide an assessment of the durability of the low alloy steels under realistic plant operational conditions. Where possible, comparisons against CMV durability have been made.

The main conclusions of the work package 5 are following summarised. Subjected to modest plant thermal transients, associated with a station that is being operated in accordance with good practice, a P23 pipework girth weld with a circumferential crack has a higher factor of safety than an equivalent CMV pipework girth weld. When subject to severe plant thermal transients, associated with poor operational practice, there is no significant benefit associated with a P23 pipework system in preference to a CMV pipework system.

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P23 pipework systems are at least as tolerant of mal-adjustment and poor maintenance practices associated with pipework hanger systems as CMV systems. Thermal transients due to poor operation have been shown to persist along the whole length of a typical pipework system. This influences the pipework weld inspection strategy, since the assumption that weld locations adjacent to the Boiler Stop Valve are most at risk is not true for every station or pipework system. This finding supports inspection data found on some UK stations. The material model developed gives a good prediction for lifetime, within the range of specimens tested. This work has shown that a P23 pipework system will be more durable than an equivalent CMV system providing that:

• Good operational practice is maintained, thereby minimising the risk of severe operational transients;

• Pipe support maintenance and adjustment is undertaken at regular intervals; • Pipe weld integrity will be enhanced if the above two points are adhered to; • Inspection regimes, for the detection and management of thermal bore cracks, should

consider all weld locations along the pipeline.

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Scientific and technical description of the results 1) Objectives of the project

Introduction The new advanced low Chromium Grades 23 and 24 steels are and will be strong candidates for new power and petrochemical plant construction and for the eventual large-scale replacement of steam pipework on existing power plant. The main aims of the project were the following:

- Mechanical and microstructural assessment of base materials, - Development of weld material and welding procedure to avoid the “bore cracking”

phenomenon, - Mechanical and microstructural assessment of similar welding, - High temperature design and assessment of welded components including welded

pressure vessels and pipework, - Microstructural modelling to predict changes as a function of time and temperature in

service operation, - Piping integrity assessment under realistic loading conditions by combined thermal

hydraulic system analysis and stress analysis tools. Some major development hurdles remained to be overcome:

- the basic knowledge of the base material in term of mechanical and microstructural evolution during the service life,

- welding consumables with proven high temperature properties needed to be developed and validated by creep rupture testing, microstructural and mechanical properties assessment,

- new steels and welds must be shown to have improved resistance to the in-service bore cracking phenomenon recently identified in existing CrMoV steel steam pipework.

These required comparison between the material properties after the manufacturing heat treatment and after additional simulated service aging, the behaviour of the welded joint and the comparison of low cycle fatigue crack initiation and growth testing, together with creep strain rate, creep rupture and creep crack growth testing, on new and original pipework steels and weldments. These tests have been carried out on experimental materials produced in full size dimensions by industrial process routes. It has also been necessary to model the effects of plant operational conditions and service loadings on the integrity of welded pipework geometries and configurations using combined thermal hydraulic system analysis and structural analysis. This modelling approach has been successfully defined and verified by the laboratory full scale simulation tests. In parallel, microstructural modelling and characterisation techniques have been employed to predict microstructural changes during long term operation of high temperature plant. This approach and the results obtained will be necessary for component design, manufacture, plant construction, and maintenance and inspection of power and petrochemical plant. A large number of tests have been carried out, and an extensive database of information for steelmakers (tubes, pipes, plates, consumables) and end users has been generated. This will be very useful for components construction and to identify the correct service operational conditions, inspections and maintenance.

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The Consortium believe that not all aspects are yet completely clarified and well defined, but it agreed that all the partners are now much more confident with these grades of materials, their applications and the evolution of their properties. Further work on the basic understanding of the effect of the alloying elements in the consumables, very long term microstructural evolution and modelling of the component behaviours could also be generated. The results obtained in this project will make a small but relevant contribution to increasing plant safety and efficiency, and thus towards achieving the Kyoto Protocol target for Europe of an 8% reduction of CO2 emission by 2010. 2) Comparison of initially planned activities and work accomplished

- The delay in the project conclusion is mainly due to the problems encountered in developing

consumables, and consequently the need for 6 additional months of work to reach at least 10000 hours for creep tests on welded joints has been required and approved,

- The main activities in the project were focused on Grade 23 because, with respect to the

original proposal, Grade 24 has recently seen much more limited requests from the market than Grade 23. Also, although notwithstanding it is included in the ASTM A213 standard for tubes, grade 24 is patented in EU by Vallourec & Mannesmann, and is not included in the ASTM A355 for pipes. Consequently, the world market tends to prefer to use Grade 23, and to avoid dissimilar welding grade 23 to grade 24, i.e. for header manufacture. However Grade 24 material was available as an industrial trial product, and tests were made to compare the mechanical properties and creep behaviours with Grade 23, as well as for some welding test simulations, aging tests, and mechanical tests after aging. The microstructural evolution of Grade 24 has also been evaluated. This modification in the project target was showed and approved in the mid term presentation.

- The consortium agreed that it is better to ensure good and consistent results on Grade 23, to

be compared with reference to grade 22 for creep and with grade 24 (comparison tests have been carried out for creep and tensile properties, aging response, and welding), instead of to scatter resources more widely and include materials with lower industrial interest. The possibilities of additional improvement in consumables for grade 23 were considered by the partners more important and promising, from the industrial point of view, than carrying out trials on dissimilar welding (e.g. welding grade 23 to grade 91). The latter task was therefore cancelled.

- In the originally submitted plan it was intended to undertake a design review of Seam welded

pressure vessels and pipework (Task 5.1.2). During the start-up project meeting in Rome on the January 29th, 2004 it was agreed that it was of less importance to devote time to the assessment of seam welded installations, due to the predominance of girth welded designs in the European Union. Hence task 3.5.1.2 was not pursued and it was declared in the second progress report.

- The decision to concentrate on the girth weld geometry has been emphasised by findings on

the summer 2004 outages on UK plant. Since bore cracking was discovered in 2001 E.ON UK instigated a comprehensive inspection/assessment programme to minimise the risk. This also involved advising stations regarding how they can minimise the risk of initiating circumferential cracks at girth welds. As a result of this work all the major defects were removed and any in-situ cracks were being monitored as part of a normal inspection regime.

15

However during the summer 2004 outages deep bore cracks have been found in one UK station, in this case at locations spread along the whole length of the steam line. Usually E.ON have found that the cracks are focussed either at the top, near to the BSV’s or at the bottom of the steam line, near to the turbine loop pipes. Metallurgical samples of these cracks indicated that the growth rates are approximately 2mm/year and are driven by thermal fatigue. In addition E.ON has also discovered bore cracks on some thinner sections of pipework, which is unexpected.

- The creep tests still running in the CSM, Dalmine and Industeel laboratories will continue

until rupture for later addition to the creep database. The results will be disseminated to the other partners of the Consortium annually in the form of a news letter.

These results will also be shared with the ECCC (European Creep Collaborative Committee). ECCC is a major EU consortium of Universities, R&D Centres, Steelmakers, Boilermakers, Turbine manufacturers, plant constructors and end users, and is currently working on the generation and assessment of creep data to define new strength values to be introduced into the EN standards. For instance the EN for the Grade 23 has been not yet issued.

16

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3) Description of activities and discussion 3.1 WP1: Characterisation of the “as-received “Steels Objectives : • Selection of test materials and plate/tube geometries • Microstructural characterisation of base material in normalised and tempered condition • Mechanical properties of base material including strength and toughness, Task 3.1.1: material production 3.1.1.1 Production of material Grade 23 and 24 steel by industrial routes (UsI, Dal) or on the laboratory scale The partners agreed to work on real component types, as defined in terms of the main market requirements for tube and pipe diameters and thicknesses. 3.1.1.1a: Grade 23 tubes and pipes A heat of grade 23 steel was ordered from ABS (Acciaierie Bertoli Safau S.p.A., Pozzuolo del Friuli, UD, Italy) by Dalmine. The steel was cast in ingots. Bars were rolled from ingots and fully annealed. The quantities were as follows: - 25 tonnes of bars of OD 280mm, used to produce pipes with OD 219mm (B) - 55 tonnes of bars of OD 360mm, used to produce pipes with OD 355mm (A) A heat from Ascometal (France) was used to produce tubes with diameter below 90mm (H), while a third heat produced in Dalmine was used to produce a second batch of pipes with OD 219mm (G). 3.1.1.1b: Grade 24 tubes and pipes A heat of grade 24 steel was cast in bars of OD 145 mm for the production of small diameter tubes for experimental use only. 3.1.1.1c: Plate product Grades 23 and 24 are not currently supplied industrially as plate. Only grade 23 is standardized in ASTM A1017. Initially, Industeel aimed to develop contacts with material producers with the aim of buying a small heat, e.g. an ingot of about 5 tonnes, of each of the grades 23 and 24 as required for the project. Industeel would like to produce thick plates with its own rolling mill equipment, but the minimum capacity of its melting shop, 60 tonnes, is really too large for a prototype plate. In the course of this search, it was noted that Industeel’s customers are not currently interested in the second grade, P/T 24. Marketing and commercial departments are often asked for information on new CrMoV steel grades, but only for P/T23 and P/T92 (W-enhanced CrMoV grades). Hence, Industeel chose to produce only P/T23 plate, in a “thin” but representative thickness, and make a fuller assessment of the properties of this new promising steel grade. In the event, the only viable option proved to be to buy a cylindrical ingot designed for pipe production from Vallourec & Mannesmann, and then produce prototype plates via a forging route. The initial ingot geometry was a cylinder of 300mm diameter by 7000mm long (about 4 tons). The chemical analyses of the Grade 23 and 24 heats used in the program are reported in the table 3.1.1.

18

Grade

C Mn Si P S Cu Ni Cr Mo Ti V Nb W B

0.04 0.10 - - - - - 1.90 0.05 - 0.20 0.02 1.45 0.0005 ASTM

Grade 23

0.10 0.60 0.5 0.010 0.030 - - 2.60 0.30 - 0.30 0.08 1.75 0.0060

23 (heat 274284) 0.06 0.47 0.24 0.008 0.001 0.12 0.09 2.15 0.13 0.03 0.24 0.06 1.57 0.0026

23 (heat 945203 0.070 0.450 0.290 0.013 0.002 0.10 0.09 2.330 0.210 0.030 0.220 0.043 1.516 0.003

23 (heat 109194) 0.067 0.460 0.260 0.019 0.005 0.14 0.13 2.102 0.112 0.027 0.214 0.045 1.579 0.003

Plate 23 0.051 0.43 0.29 0.01 0.015 0.14 0.09 1.90 0.07 0.05 0.21 0.041 1.58 0.0015

0.05 0.30 0.15 - - - - 2.20 0.70 0.20 - - 0.0015 ASTM

Grade 24

0.10 0.70 0.45 0.020 0.010 - - 2.60 1.10 0.30 - - 0.0070

24 (heat 931224) 0.06 0.44 0.25 0.012 0.003 0.16 0.13 2.32 0.97 0.06 0.22 0.004 0.0024

Table 3.1.1: ASTM Standards and chemical analysis of Grade 23 and 24 heats Table 3.1.2 summarizes the heat treatment requirements of the ASTM standards. Industrial products have been produced for both grades. Table 1.3 shows the dimension of the tubes and pipes manufactured by Dalmine and plates by Industeel.

ASTM Standard

Normalising (N) or Austenitising (A) temperature Cooling Tempering temperature

A213 N at not less than 1040°C Not defined Not less than 730°C

A335 N at not less than 1040°C Air cooling Or Accelerated cooling Not less than 730°C

A1017 N at 1040 to 1095°C Or A at 1040 to 1095°C Air Blasting Or Liquid Quenching Not less than 730°C

Table 3.1.2: ASTM Standards for heat treatment

The following route has been followed by Dalmine to produce the tubes and pipes: - heating of the billets, - piercing and hot rolling in small or medium size mills, - heat treatment in tunnel furnaces - NDT and finishing.

The following route has been followed by Industeel to produce the small prototype plates:

- A four tons cylindrical ingot (φ300 x L7000mm) was bought by Industeel, - Ingot was cut into slices (φ300 x L280mm), - Slices were forged into small plates (1000x400x50mm) in a small forge near Le Creusot.

After forging, the small plates were sent to the plant of Charleroi (Industeel Belgium) to test various heat treatments and choose the most promising. Figure 3.1.1 shows examples of the pipes and plates produced.

Figure 3.1.1.: Grade 23 pipes (OD219 x WT31.75mm) and plate (1000 x 400 x 50mm)

3.1.1.2. Heat treatment of base material for testing and conventional mechanical tests for product qualification

19

The products have been heat treated by austenitising and tempering as summarised in table 3.1.3 Quenching after austenitisation was necessary for products with wall thickness higher then 10mm in order to avoid the formation of ferrite, which reduces the mechanical and creep properties of the final products. To improve the impact properties of large diameter pipes in grade 23, the tempering temperature has been increased to 780°C.

Table 3.1.3: product forms and heat treatments (P= Pipe; T=tube) Task 3.1.2: microstructural characterization of base material 3.1.2.1 Characterisation of microstructures using light (LM) and electron microscopy (SEM, TEM) with EDS analysis and other analytical techniques to obtain background information on the grain size, dislocation density, type, dimension and density of precipitates present in the matrix All the materials have been characterised by light microscopy in order to evaluate the grain size and microstructure and the extent of the decarburisation layer on the external surface. Figures 3.1.2 and 3.1.3 show some examples of grade 23 and 24. Both grades show a fully tempered bainitic microstructure. For the pipe casts G and J, the measured decarburisation layer thickness is of the order of 190 and 130 μm respectively for the external and internal surfaces. Grain size is about 5 μm for both casts. Table 3.1.4 summarises the results obtained.

Material code

Material Dimensions (mm) Heat treatment conditions

A P23 355x35 (ODxWT) austenitised at 1070°C, quenched and tempered at 760°C, and tempered at 780°C

B P23 219x31.75 (ODxWT) austenitised at 1070°C, quenched and tempered at 760°C

G P23 219x31.75 (ODxWT) austenitised at 1070°C, quenched and tempered at 780°C

H P23 88.9x17 (ODxWT) austenitised at 1050°C, quenched and tempered at 760°C

P (plate) Plate 23 IND (1000 x 400 x 50mm)

Tests performed : normalised (austenitised at 1070°C for 2Hr + air cooling) + tempered at temperatures between 730 and 770°C for 2Hr

quenched (austenitisation at 1070°C for 2Hr + water quenching) + tempered at temperatures between 730 and 770°C for 2Hr

Final heat treatment for delivered plates is:

quenched (austenitisation at 1070°C for 2Hr + water quenching) + tempered at 770°C for 2Hr

J P23 457x17.5 (ODxWT) DAL preliminary test

material austenitised at 1070°C, quenched and tempered at 780°C

C T24 44.5x 6.3 (ODxWT) normalised at 1000°C and tempered at 750°C

E T24 76x12.5 (ODxWT) austenitised at 1000°C, quenched and tempered at 750°C

(M) (T24)1) 159x20 (ODxWT) (M) from Vallourec & Mannesmann only for aging test at VTT

N T23 76x12,5 (ODxWT) normalised at 1050°C and tempered at 760°C

20

P23 code H (ODxWT 88.9x17mm) T24 Code C (ODxWT 3x6.8 mm)

Figure 3.1.2- Examples of microstructure for grade 23 and 24

Figure 3.1.3: T24 code C, external surface

Code G

OD 219mm x WT 31.75mm Code J

OD 457mm x WT 17.5mm

External L: 202 – 202 - 198 T: 202 – 196 – 200

L: 227 – 228 – 236 T: 240 – 236 – 238

Middle L: 207 – 208 – 209 T: 209 – 206 – 205

L: 228 – 232 – 230 T: 233 – 238 – 232

Hardness (HV10)

Internal L: 212 – 215 – 215 T: 203 – 206 – 202

L: 230 – 230 – 228 T: 228 – 225 – 227

Grain size 5.3 μm 4.8 μm

External: 180 External: 200 Decarburisation (μ m) Internal: 120 Internal: 140

Table 3.1.4: example of the results obtained (L= longitudinal direction; T= transverse direction)

More detailed analysis using SEM and TEM have been carried out only on products selected for the main mechanical characterisation. Figures 3.1.4 show an example of the SEM microstructure of Grade 23 cast G. Both grades have also been characterised by TEM in the final heat treatment condition to investigate the detailed microstructure, precipitate morphology and dislocation distribution. These analyses were necessary to provide the reference condition to compare with the results to be generated after aging. Both grades are characterised by small precipitates along the grain boundary and inside the grains and subgrains. Figure 3.1.5 shows examples of the microstructures of casts G and C in thin foils, while figure 3.1.6 shows precipitates in Grade 23 cast G on a carbon extraction replica. Figure 3.1.7 summarises the bimodal size distribution of precipitates. The smaller size class is mainly due to MX particles and the bigger is due to M23C6 carbides.

21

Figure 3.1.4: P23 cast G (OD 219mm x WT 31.75mm): SEM microstructure

P23 T24

Figure 3.1.5: Quenched and tempered casts G and C (TEM on Thin Foil)

Figure 3.1.6: precipitates (TEM extraction replica) on P23 Cast G quenched and tempered

22

0

5

10

15

20

25

30

0 100 200 300 400 500Particles size (nm)

Freq

uenc

y (%

)

Figure 3.1.7: size distribution of precipitates The plate materials have been characterised by optical microscopy (Figure 3.1.8). All the products produced by different process routes after heat treatment show a very good homogeneity of microstructure.

Skin 1 (X200) Mid thickness (X200) Skin 2 (X200)

¼ thickness 1 (X200) ¼ thickness 2 (X200)

Figure 3.1.8: Microstructure of plate in different positions Task 3.1.3: mechanical characterization of base material Characterisation of as-received materials properties by the following tests: 3.1.3.1.Tensile strength from room temperature up to 650°C a) Tube and Pipes Table 3.1.5 summarises the minimum mechanical properties specified in ASTM Standards.

ASTM Standard YS min (MPa) UTS (MPa) El% min for 2’ of 50mm (%) Maximum hardness

A213 400 510 min 20 220HB Or 230HV Or 97HRB

A335 400 510 min 20 (Basic for 8mm thick) Or El% min = 1,25xth+10

Not defined for grade 23

A1017 400 510 to 690 Not defined 220HB or 97HRB

Table 3.1.5: minimum values for ASTM standards Tensile and hardness tests on both grades and all the products have been carried out, and a summary of the results is presented in table 3.1.6.

23

Pipe size ODxWT Code YS (MPa) UTS (MPa) Elongation (%) Hardness HV10 219.1x31.75 mm G 581 675 26 <215

457x17.5 mm J 602 687 39 <245 88.9x17.5 mm H 612 694 36 223 76x12.5 mm N 590 686 33 225

44.5mm x 6.3 mm C 599 700 24 205 Table 3.1.6: Examples of mechanical properties obtained on different products and grades

Hot tensile properties have been also performed on selected grade 23 pipes. Results for casts G and J as a function of temperature up to 650°C are reported in figure 3.1.9.

0

100

200

300

400

500

600

300 350 400 450 500 550 600 650 700

Temperature [°C]

MPa

YieldRupturePoli. (Rupture)Poli. (Yield)

0

100

200

300

400

500

600

700

300 350 400 450 500 550 600 650 700

Temperature [°C]

MPa

YieldRupturePoli. (Rupture)Poli. (Yield)

Figure 3.1.9 - Hot tensile data on P23 pipe casts G and J

b) Plates Two sets of plates have been tested, normalized at 1070°C for 2 hours, and then cooled to room temperature by water quenching for the first set and by air quenching for the second set. Each set was then tempered for 2 hours at three temperatures (730, 750 and 770°C) to give final mechanical properties. Tensile properties are good with these heat treatments and the requirements are fulfilled. The table 3.1.7 summarizes the tensile properties obtained in each heat treatment configuration at the plate mid-thickness position.

1070°C-2h + Water quenching 1070°C-2h + Air quenching

Orientation Transversal Longitudinal Transversal Longitudinal

Tempering YS (MPa)

UTS (MPa)

El (%)

YS (MPa)

UTS (MPa)

El (%)

YS (MPa)

UTS (MPa)

El (%)

YS (MPa)

UTS (MPa)

El (%)

730°C 2h 681 733 36 696 747 41 654 729 37 654 720 43 750°C 2h 664 723 41 654 716 44 651 721 42 652 722 40

770°C 2h 573 642 40 579 641 39 565 641 40 559 638 42

Table 3.1.7: mechanical test results as function of cooling rate and tempering conditions

The delivered products have been heat treated using the following conditions: - Water Quenching after austenitisation at 1070°C for 2Hrs - Tempering at 770°C for 2Hrs

Figure 3.1.10 shows a comparison of the tensile test results between the pipe A and the plate P (delivered conditions). 3.1.3.2 Hardness The hardness values of the different products have been measured. The results are shown in tables 3.1.4 and 3.1.6: they fulfil the standard requirements for the grades 23 and 24.

YS

UTS

YS

UTS

YS

UTS

YS

UTS

24

Material A and P - Tensile Tests

0

100

200

300

400

500

600

700

800

-100 0 100 200 300 400 500 600 700 800Temperature (°C)

YS /

UTS

(MPa

)

YS Longitudinal (Tube A) UTS Longitudinal (Tube A)

YS Circumferential (Tube A) UTS Circumferential (Tube A)

YS Longitudinal (Plate P) UTS Longitudinal (Plate P)

YS Transverse (Plate P) UTS Transversel (Plate P)

Figure 3.1.10: comparison of the mechanical properties of the Pipe A and the Plate P

3.1.3.3. Impact and FATT curve The impact properties and fracture appearance transition temperature (FATT) curves have been measured on Cast G in the circumferential orientation in the as received condition. Figure 3.1.11 shows the results. The orientation is C-L (normal to notch = circumferential; orientation of propagation = longitudinal). Location is quarter thickness (inside and outside diameter). The same orientation is used for fracture mechanics.

ALOAS A213gr23 - DALMINE PIPEAs received - Circumferential orientation

0

50

100

150

200

250

-60 -40 -20 0 20 40 60 80 100 120 140

Temperature (°C)

KV

(J)

ALOAS A213gr23 - DALMINE PIPEAs received - Circumferential orientation

0

10

20

30

40

50

60

70

80

90

100

-60 -40 -20 0 20 40 60 80 100 120 140

Temperature (°C)

Frac

ture

app

eara

nce

cris

talli

nity

(%)

Figure 3.1.11: FATT curve for pipe Code G

The following table 3.1.8 summarizes the Charpy V-notch (CVN) properties obtained at mid-thickness and in the sub-skin position in transverse orientation at 0°C for the plate P.

1070°C-2h + Water quenching 1070°C-2h + Air quenching

Orientation Transverse Sub-Skin

Transverse Mid-Thickness

Transverse Sub-Skin

Transverse Mid-Thickness

Tempering CVN1 (J)

CVN2 (J)

CVN3 (J)

CVN1 (J)

CVN2 (J)

CVN3 (J)

CVN1 (J)

CVN2 (J)

CVN3 (J)

CVN1 (J)

CVN2 (J)

CVN3 (J)

730°C 2h 18 17 14 178 54 220 6 6 6 6 6 6 750°C 2h 13 14 14 38 35 43 6 5 5 5 5 5

770°C 2h 19 27 38 61 230 231 8 8 8 9 8 8

Table 3.1.8: Charpy V impact results on plate P with different cooling rate and tempering conditions

25

Air quenching is clearly not sufficient to get the correct microstructure in the plates. Water quenching is the only way to get the required fully bainitic microstructure in 50mm thick product. The tempering heat treatment chosen by Industeel for delivery is 770°C for 2 hours. This is the best compromise to get both correct tensile properties (to satisfy A1017 maximum UTS criteria) and good Charpy V properties. The sub-skin CVN results are not as good as those for mid-thickness specimens. This is due to the forging process which involves low deformation of the sub-skin region near the dies. This phenomenon involves poor recrystallisation of the material, leading to larger grain size and lower CVN properties. Impact toughness has been measured on material A and P. CVN transition curves (figure 3.1.12) have been plotted in longitudinal and circumferential orientation for material A, in as delivered and after accelerated ageing by step cooling, and in longitudinal and transverse orientation for plates P. Table 3.1.9 summarises the results obtained.

Material A - longitudinal orientation - As Received

0

50

100

150

200

250

300

-40 -20 0 20 40 60 80 100

Temperature (°C)

KV

(J)

0

20

40

60

80

100

120

Shea

r App

eara

nce

(%)

Material A - circumferential orientation - As Received

0

50

100

150

200

250

300

-40 -20 0 20 40 60 80 100

Temperature (°C)

KV

(J)

0

20

40

60

80

100

120

Shea

r App

eara

nce

(%)

Material A - circumferential orientation - After Step Cooling

0

50

100

150

200

250

-40 -20 0 20 40 60 80 100

Temperature (°C)

KV

(J)

0

10

20

30

40

50

60

70

80

90

100

Shea

r App

eara

nce

(%)

As delivered - Transverse orientation

0

50

100

150

200

250

300

-40 -20 0 20 40 60 80 100

Temperature (°C)

KV

(J)

0

10

20

30

40

50

60

70

80

90

100

Shea

r App

eara

nce

(%)

EnergyShear Appearance

As delivered - Longitudinal orientation

0

50

100

150

200

250

300

-40 -20 0 20 40 60 80 100

Temperature (°C)

KV

(J)

0

10

20

30

40

50

60

70

80

90

100

Shea

r App

eara

nce

(%)

EnergyShear Appearance

Figure 3.1.12: Charpy transition curves - transverse (a) and longitudinal orientation (b) of plates.

Material Conditions TK28J (°C) TK41J (°C) TK54J (°C) FATT (°C)

A – Circumferential As received 24 32 38 54 A – Circumferential Step Cooling 10 18 24 40

A - Longitudinal As received 23 30 34 44 P - Longitudinal As received 11 15 18 20 P - Transversal As received 0 3 7 15

Table 3.1.9: Summary of the impact toughness results 3.1.3.4 Short term creep tests The creep test programme includes short, medium and long term tests at 550, 575, and 600°C, with further tests at 620-625°C and 650-660°C to compare with data on welded joints. Results are detailed in the report on Work Package 3.

26

3.1.3.5 Selection of the most promising material The materials selected for the main characterisation were: - Grade 23 cast G for the full testing program, including welding trials, - Grade 23 cast H for the simulation tests at IWM, - Grade 24 cast E for comparison of mechanical, creep, and aging tests, - Grade 23 Plate P for a large testing program, including welding trials. 3.1.3.6 Fracture toughness The test method used here is ASTM E 1921-02 (last revision), also known as “Mastercurve determination” and is based on the multi-temperature option. The material has been tested in as-received condition (pipe G from Dalmine) and after accelerated ageing (step cooling) to measure the effect of thermal ageing on the toughness transition. The values of To (temperature for which median quasi static fracture toughness KJc is equal to 100MPa.m-1) obtained with the multi-temperature option of the mastercurve (figure 3.1.13) approach are given in the table 3.1.10.

To As Received multi 20°C To Step Cooling multi 27°C Table 3.1.10: To values

Both mono-temperature (one set of specimens is broken at one unique temperature and then To is calculated with the values of the tests) and multi-temperature (the set of specimens is broken at different temperature, allowing a better estimation of transition behaviour) option of the calculus of mastercurve give very similar results. The dispersion of the results is about 1°C between mono and multi-temperature. It could have been sufficient to do tests at temperature predicted by Charpy-V tests. It proves that A213 grade 23 toughness transition is well predicted by the mastercurve approach. Only tests done at “high temperature” are ductile and are not in the tolerance bounds. All tests were carried out in compliance with the ASTM E 1921-02 standard. The validity of these tests performed to predict the Mastercurve parameter has been proved.

Mastercurve - Multitemperature option - A213 gr. 23 - Tube Dalmine

0

50

100

150

200

250

300

350

400

-60 -40 -20 0 20 40 60

T-To (°C)

KJc

(1') (

MPa

.m1/

2 )

Rupture 1%

Rupture 5%

Rupture 50%

Rupture 95%

Rupture 99%

ALOAS A213gr23 As Received Tube Dalmine

ALOAS A213gr23 Step Cooling Tube Dalmine

Fully ductile tests

Figure 3.1.13: mastercurve

Figure 3.1.14 shows the macrographs of all the specimens. The decision to carry out all the calculations using both the mono and the multi-temperature option seems to be a good choice. The mono-temperature option gives a good prediction of To

but not the shape of the transition curve. On the other hand, the multi-temperature option determines the evolution of toughness as a function of temperature.

27

7 8 9 10 11 12

1 2 3 4 5 6

Figure 3.1.14: Examples of broken specimens

It appears that A213 grade 23 is a very ductile steel even at temperature just above To. It also appears that thermal ageing (via step cooling) has no major effects on the toughness transition. Selection of Test Material Grade Considering the programme in relation to the original proposal: During the first year of activities, it was noted that there was much less demand from the market for Grade 24 than for Grade 23. It was also discovered that although it is included in ASTM A213 for tubes, Grade 24 is patented in the EU by Vallourec & Mannesmann, and is not included in the ASTM A355 for pipes. Consequently, the world market is more prone to use Grade 23 for tube and pipe products. Selection of Grade 23 throughout also avoids dissimilar welding Grades 23 to 24 in header manufacture. Today, Grades 23 and 24 are not supplied industrially as plate. Only Grade 23 is standardized in ASTM A1017 for the plate product. The partners therefore agreed to do more activities on Grade 23, and carry out only limited testing of Grade 24 for comparison. Grade 24 was used to compare creep behaviour with Grade 23 and to determine mechanical properties and microstructural evolution after aging. This modification to the project target was shown, explained and approved in the presentation of the mid term progress report. Conclusions of Work Package 1 The materials for testing have been produced in industrial plant and successfully qualified by conventional mechanical testing and metallographic characterisation. Grade 23 has been selected for the main test programme.

28

3.2 WP2: Development of weldments

Objectives: • development of welding consumables • optimisation of welding procedure (welding technology and parameters) • qualification and simulation (using Gleeble testing) of welded joints • determination of mechanical properties of welded joints Task 3.2.1: Welding trials and modelling: Welding development has involved two main activities. Development of an optimised welding consumable type by means of short term creep testing trials of candidate compositions was undertaken throughout the ALoAS project. At month 12 of the project, the best consumable composition available at that stage, B323B, was selected for the main programme of weld characterisation and testing. B323B was used by project partner ISQ to produce a series of test blocks for long term high temperature creep, LCF, crack growth, and pressurised tube testing. The data produced on B323B provide baseline information on P23 weld metals. The development work that has continued since B323B was selected now shows that further improvements to P23 consumable formulation are achievable. 3.2.1.1 - Development of welding consumables for P23 and P24 steel to avoid premature low creep ductility failure Metrode, acting as subcontractors to ISQ, produced several series of experimental P24 and P23 weld metals to explore various options in compositional formulation within and outside the base material specifications, table 2.1. These consumables were used to make small AWS type weld deposit samples, figure 2.1, for assessment of ambient temperature tensile and impact properties (by MPL) and short term all-weld-metal creep testing (by E.ON UK). Selected Metrode consumables and commercial alternatives were also supplied to ISQ for deposition of weld pads, section 2.1.3.

10-12mm

20°

12.5mm

CMn

Buttering

The initial trial consumables and those obtained from commercial suppliers were: - METRODE Chromet 23L - The Metrode electrode Chromet 23L, developed for tube welding without PWHT, has a low carbon level around 0.05% and a deliberate nickel addition to optimize as-welded toughness. Another characteristic is the relatively moderate levels of creep strengthening elements such as W, Nb and V, and of the ambient temperature strengthening elements Si and Mn. - METRODE Chromet 23H (B322 and B323)) - The original Metrode Chromet 23H development has a composition very close to the P23 base material with no nickel addition. With this starting point, Metrode manufactured the first two trial SMAW batches B322 and B323. The first batch, B322, aimed to replicate Chromet 23H, while the second, B323, aimed to maximize the carbide formers (creep strengtheners, especially V, Nb, B) whilst nominally remaining within the ASTM A213 T23 analysis limits (B323). Neither batch had any nickel addition. - METRODE B323B – The B323 formulation, based on a small batch, was replicated by manufacture of a much larger second batch, designated B323B (sizes 3.2mm and 4.0mm diameter), for test material production welding by ISQ.

Figure 3.2.1. Metrode weld deposit sample geometry

29

- METRODE 2CrWV - In addition to the SMAW electrodes Metrode also obtained P23 solid wire for submerged arc welding. This consumable contains about 0.1% nickel and has weld metal properties intermediate between 23L and 23H (B323). - METRODE P24 consumable - Metrode developed a Chromet 24 variant B87 with higher carbon. - HITEHR consumables - Experimental consumables develop by METRODE and ISQ during an earlier European project HITEHR (MH) were also characterized. This composition is similar to P24 but included cobalt and boron. - SUMITOMO P23 consumables - ISQ purchased Sumitomo electrodes for evaluation. Welds were deposited with 3.2mm diameter electrodes at ~120A and ~24V, 175°C preheat and maximum interpass temperature, and PWHT of 715°C/1 hour. - Bohler-Thyssen Thermanit P24 consumables – These have a hybrid formulation based on P24 but with Nb substituted for Ti (supplied via COST 536 project). Following stress rupture testing, three further trial series of consumables were also manufactured, as summarised in table 3.2.1 and discussed in section 2.1.3. B575B and B576 were formulated to clarify the separate influences of creep strengthening and surface-active elements, B702 and B703 explored the effects of nickel additions, and the final series A71 - A72 - A84 - A85 explored further element combinations with emphasis on optimising the niobium content.

Batch C Mn Si S P Cr Mo Ni B* N Ti V Al Cu Nb W ASTM A213 Grade T23

0.04 0.10

0.10 0.60

0.50 Max

0.010 Max

0.030 Max

1.9 2.6

0.05 0.30 NS 5 -

60 0.030 Max NS 0.20

0.30 0.030 Max NS 0.02

0.08 1.45 1.75

Chromet 23L 0.05 0.5 0.2 0.01 0.01 2.2 0.1 0.80 10 <0.02 - 0.21 0.005 - 0.03 1.5

B322 0.056 0.55 0.29 0.011 0.012 2.38 0.12 0.03 12 0.022 0.004 0.24 <0.001 0.05 0.05 1.74

B323 0.078 0.59 0.39 0.008 0.010 2.36 0.12 0.03 48 0.028 0.025 0.30 <0.001 0.04 0.09 1.75 Metrode SAW deposit

0.058 0.51 0.36 0.002 0.015 2.09 0.07 0.12 8 0.012 0.009 0.19 0.013 0.11 0.05 1.49

HITHER (MH1) 0.08 0.52 0.27 0.008 0.010 2.33 0.97 0.05 0.004 0.019 0.057 0.23 - 0.07 0.007 -

HITHER (MH2) 0.08 0.48 0.23 0.007 0.008 2.12 0.91 0.04 0.003 0.014 0.051 0.19 - 0.03 0.006 -

Sumitomo P23 HCM2S

0.07 0.71 0.43 0.002 0.009 2.01 0.09 0.99 -- -- -- 0.28 -- -- 0.04 1.47

Bohler Thyssen P24

0.10 0.55 0.24 0.007 0.010 2.51 0.93 0.15 0.003 -- 0.0002 0.21 -- 0.04 0.049 --

B323B(3.2) WO22919 3.2mm

0.075 0.51 0.36 0.010 0.012 2.11 0.11 0.06 41 0.032 0.017 0.29 <0.001 0.06 0.08 1.87

B323B(4.0) WO22920 4.0mm

0.078 0.55 0.37 0.009 0.011 1.98 0.10 0.05 46 0.022 0.024 0.27 <0.001 0.06 0.09 1.70

B575B 0.065 0.63 0.36 0.008 0.009 2.23 0.11 0.04 47 0.021 0.025 0.21 <0.001 0.03 0.05 1.64

B576 0.078 0.53 0.26 0.009 0.009 2.15 0.12 0.03 13 0.016 0.004 0.29 <0.001 0.03 0.08 1.70

B702 0.057 0.48 0.24 0.009 0.009 2.09 0.14 0.75 7 0.028 0.004 0.22 0.002 0.06 0.04 1.78

B703 0.084 0.55 0.36 0.008 0.009 2.10 0.13 0.74 39 0.029 0.021 0.32 <0.001 0.05 0.08 1.99

A71 0.076 0.52 0.33 0.010 0.009 2.08 0.11 0.74 11 0.014 0.005 0.28 <0.001 0.05 0.096 1.43

A72 0.065 0.71 0.36 0.011 0.009 2.13 0.12 0.77 9 0.035 0.004 0.26 <0.001 0.05 0.061 1.54

A84 0.063 0.70 0.36 0.010 0.010 2.15 0.11 0.78 9 0.029 0.005 0.28 <0.001 0.05 0.030 1.64

A85 0.075 0.51 0.37 0.010 0.007 2.23 0.92 0.77 7 0.024 0.005 0.22 0.001 0.05 0.061 0.04

Table 3.2.1: All weld metal analysis and aim compositions, wt %. Boron quoted as ppm. NS = not specified

30

The results of the mechanical tests on these weld metals are given in table 3.2.2. More details on WPS applied are included in APPENDIX 1

+20ºc

0ºc

-20ºc

Rm MPa

Rp 0.2% MPa

Elongation

%

Z %

Hardness, Hv(10Kg)

Cap Mid thickness

PWHT ºC/h

CVN J

CVN J

CVN J

4D 5D Average Max Average Max

23L AW 22 - - - 870 19 16 50 290 - 350 - AW - - - - - - - - 312 342 322 336

715/2 50(32) - - 721 639 20.5 17 61 273 281 229 245 B322 740/2 123(111) - - 666 581 22.5 20 66 238 243 241 249 AW - - - - - - - - 372 397 350 376

715/2 20(18) - - 779 711 10 8.5 15 313 325 281 283 B323 740/2 65(31) - - 723 657 9 7.5 15 267 276 264 258 AW - - - - - - - - 287 322 287 322

715/2 35(20) - - 695 630 16 12 38 250 270 260 281 SAW 740/2 145(48) - - 645 572 22 17.5 53 233 245 244 254

MH1 705/10 165 151 145 672 585 18 16 32 240 343 226 235 MH2 705/10 188 163 138 - - - - - 232 240 218 227

Sumitomo 705/10 72 - - 803 750 21 - - - - - - B-Thyssen 705/10 120 - - 620 520 - 18 - - - - -

740/2 70(60) - - 745 668 8.5 8 20 - - 263 272 B323B(3.2) WO22919 (3.2) 760/2 99(70) - - 725 637 15 15 30 - - 241 249

740/2 61(51) - - 688 640 9.5 9 17 - - 255 270 B323B(4.0) WO22920 (4.0) 760/2 91(45) - - 686 621 10 9 16 - - 240 251

B575B 746/1 128(115) - 31(21) 706 629 18 15 36 256 260 236 243 B576 746/1 66(63) - 17(10) 719 627 22 19 60 262 274 247 254

740/2 111(97 - 46(32) 682 599 24 20 64 239 245 230 233 B702 760/2 146(138 - 90(76) 637 549 23.5 21 66 216 221 212 215

740/2 60(52) - 23(20) 797 722 13.5 10.5 21 271 274 263 279 B703 760/2 82(63) - 23(15) 764 685 13 10 19 258 258 258 266 A71 740/2 94(80) - 32(26) 725 640 18 16.5 55 250 254 248 254 A72 740/2 127(112) - 66(57) 732 625 20.5 18 62 246 251 247 251 A84 740/2 136(130) - 95(74) 704 620 20.5 18 52 139 242 240 242 A85 740/2 125(119) - 67(53) 748 672 19.5 17.5 66 251 258 249 256

Table 3.2.2: Ambient temperature all weld metal tensile, hardness and impact properties 3.2.1.2 – Welding simulation and modelling, using Gleeble testing A methodology was developed [1,2] for the evaluation of material weldability in terms of reheat cracking susceptibility and toughness and ductility. It is based on the application of single- and multi-cycle thermal (Gleeble) simulation of base material in order to produce representative HAZ microstructures over a range of peak temperatures Tp and weld t8/5 cooling times, as well as, investigation of actual weldments both in the as-welded, thermally re-heated and PWHT conditions. The weldment studied here was a circumferential multipass weld on a thick-wall P23 steam-pipe steel made at ISQ using Metrode B323B (experimental) filler metal with broadly matching chemical composition. Simulated HAZs and actual weldments were subjected to the Belgian Welding Institute (BWI) slow strain rate tensile reheat cracking test [3], standard Charpy impact toughness testing, microstructural investigation using optical microscopy, parent steel & weld chemical analyses, and Vickers hardness measurements. Reheat cracking susceptibility was expressed in terms of the attained value of reduction of area (RA) in a specimen, whereas toughness was assessed by determining the individual Charpy absorbed energies (J) and the 28 J impact toughness transition temperature, T28J. Characterisation of the various HAZ and weld metal microstructures makes it possible to link cracking sensitivity and toughness and ductility behaviour to the material chemical composition and corresponding microstructural features. Overall, the results demonstrated that the B323B weld metal is far more critical than the P23 steel HAZ, both in terms of reheat cracking sensitivity and ductility and toughness. In the as-welded condition, the weld metal exhibited an excessive hardness of ≈380 HV and extremely

31

low Charpy toughness of ≈7 J at +20 °C. The highest and the lowest hardness were both associated with re-heated microstructures close to bead boundaries between successive weld layers. Formation of local, hardened and softened microstructures within re-heated multipass weld metal is presumably a result of complex precipitation reactions causing re-precipitation, over-ageing and tempering. PWHT (760 °C / 2h) greatly improved the toughness of the weldment. However, it also inevitably raises the risk of reheat cracking, particularly in the weld metal which, at worst, showed reduction of area (RA) no more than 2–3 % in the BWI cracking test. A startling observation was that the recorded RA values for the thermally re-heated weld metal were significantly lower than those for the simulated HAZ, even though the applied Gleeble cycles were identical, see Table 3.2.3. Fractography of reheat cracked specimens revealed pronounced localisation of damage, see Figures. 3.2.2 and 3.2.3. According to optical microscopy, the bainitic-martensitic microstructures appeared more or less similar between the thermally re-heated weld metal and the thermally simulated HAZ associated with identical Gleeble cycles. Thus, underlying causes of deficient RA and hence extremely high susceptibility to reheat cracking in the CG-(R)-WM, in particular, cannot be attributed to the original weld solidification structure. Differences in reheat cracking sensitivity between the weld metal and the HAZ must originate from sub-structural features, e.g., carbide precipitation structure and/or martensite lath morphology. The high reheat cracking sensitivity of the weld metal might be explained by its high B, Nb (and W) and hence pronounced strengthening of grain interiors because of W in solution and/or precipitates and precipitation of fine, coherent Nb-rich carbides/ carbonitrides. Complete characterisation of the sub-structure would, however, require TEM microscopy, which was outside the scope of the present work.

Figure 3.2.2 - Reheat cracking in P23 weld Figure 3.2.3 – Reheat cracking in P23 weld metal subsequently reheated into the coarse- metal discovered from a CT specimen grained-temperature region (Tp = 1340 °C): extracted from a laboratory-scale welded CG(R)-WM – filler metal B323B. sample (PWHT) – B323B filler metal Degradation of the toughness in the CGHAZ in the as-welded condition was attributed to coarse prior-austenite grain size and the presence of un-tempered martensite with very high hardness (Table 2.3). Reasonably high toughness occurring in the (i) FGHAZ and (ii) the CGHAZ in the PWHT condition, can in turn be ascribed, respectively, to (i) undissolved carbides, low-C martensite and small grain size, and (ii) tempered low-C martensite with reduced hardness. In the case of multipass weldments, low toughness can occur due to a microstructure composed of initial coarse martensite in conjunction with new small-grain untempered martensite; a structure that results from a peak temperature high enough to cause partial re-austenisation that overrides prior martensite tempering, but inadequately high to result in complete grain

32

refinement/normalising. Otherwise, multi-cycle welds are likely to exhibit recovery in toughness, in relation to single-cycle welds. In the present case, PWHT seems mandatory to guarantee weldment hardness not exceeding 350 HV, as well as to ensure adequate toughness for the multipass weld metal made using matching B323B filler metal. Welding techniques accentuating the normalising effect of multipass welding should be beneficial both in terms of reheat cracking and toughness. For thick-section welds, filler metals with mis-matching chemical compositions can be expected to reduce weld metal reheat cracking sensitivity.

Specimen / Case Specimen orientation

Initial diameter (mm)

Reduction of area (%)

P23 Gleeble-simulated HAZ CGHAZ_10 Transverse 7 15.1 CGHAZ_10 Transverse 7 16.4 CGHAZ_15 Transverse 7 19.0 CGHAZ_15 Transverse 7 16.4 CGHAZ_25 Transverse 7 16.4 CGHAZ_25 Transverse 7 16.4 FG(R)-CGHAZ_10 Transverse 7 58.7 FG(R)-CGHAZ_10 Transverse 7 58.7 FG(R)-CGHAZ_25 Transverse 7 73.5 FG(R)-CGHAZ_25 Transverse 7 56.8 FG(R)-CGHAZ_25 Transverse 7 69.7

P22 Gleeble-simulated HAZ (reference material) P22: CGHAZ_10 Transverse 7 26.5 P22: FG(R)-CGHAZ_10 Transverse 7 59.0

P23 thermally (Gleeble) re-heated multipass weld metal – B323(experimental) 7 3.0 Transverse 7 3.0 7 1.7

CG(R)-WM_10

Longitudinal 7 1.4 7 15 Transverse 7 15 7 16

FG(R)-CG(R)-WM_10

Longitudinal 7 17

Note: CGHAZ_10, CGHAZ_15 and CGHAZ 25: single-cycled CGHAZ (Tp = 1340°C) simulated using the weld t8/5 cooling time of 10, 15 and 25 sec., respectively. Note: FG-(R)-CGHAZ_10 and FG-(R)-CGHAZ_25: double-cycled FG(R)-CGHAZ, i.e., the initial CGHAZ (Tp1 = 1340°C) subsequently reheated to the Tp2 = 1020°C corresponding to the FGHAZ, simulated using the weld t8/5 cooling time of 10 and 25 sec., respectively. Note: CG(R)-WM: Tp1 = 1340 °C ; FG(R)-CG(R)-WM: Tp1 = 1340 °C & Tp2 = 1020 °C ; t8/5 = 10 sec.; Belgian Welding Institute (BWI) reheat-cracking test: strain to fracture, 0.5 mm/min; PWHT = 760 °C

Table 3.2.3: Reheat cracking test results for the simulated HAZs of P23 steel G, P22 reference steel, and the thermally Gleeble re-heated microstructures of the B323B multipass weld metal.

It was found that the hardness measurements of Gleeble simulated and actual weldment HAZ microstructures showed consistent results. This confirms that thermal simulation had resulted in microstructures that resemble corresponding HAZs of the original P23 multipass pipe weldment studied here. Furthermore, also thermally (Gleeble) re-heated and original welds revealed more or less similar hardness, which proves that the Gleeble re-heated weld metal microstructures correspond realistically to those of the actual multipass weld metal reheated due to successive weld thermal cycles. The unexceptionally high weld metal maximum hardness of ≈380 HV can be explained by the fact that the present P23 weldment was investigated in the as-welded condition, i.e., without PWHT. Besides, the weld had rather high contents of B and Nb that both contribute to hardness elevation. That hardness in the cases of thermally re-heated and actual weld metal were practically equivalent, gives a good reason to believe that the BWI reheat cracking test results associated with the Gleeble simulated/re-heated microstructures can describe realistically reheat cracking susceptibility of actual weldments, as far as microstructural features are concerned.

33

Further evidence demonstrating high reheat cracking sensitivity of the present B323B weld metal was obtained when investigating another P23 weldment otherwise similar to the present one, except that it was delivered in the PWHT condition. Microstructural examination of some CT specimens extracted from the weld metal revealed significant number of small reheat cracks close to the weld surface in the farthermost location from the pre-fatigued notch, see Fig 3.2.3. The fact that severe reheat cracking had occurred in the weld metal in the case of a relatively small laboratory-scale sample, can be considered as undisputable evidence of the high cracking sensitivity of the weld made using closely matching B323B filler metal. This suggests that the high susceptibility of the present B323B weld metal to reheat cracking experienced in the BWI tensile tests on thermally re-heated weld metals, is likely not to be an overestimation of cracking susceptibility associated with the Gleeble simulation & BWI test itself, but instead, a realistic demonstration of high inherent sensitivity of the original B323B weld metal to reheat cracking. Summary The present methodology comprising thermal (Gleeble) re-heating of the actual weld metal in conjunction with the BWI tensile reheat cracking test was successfully applied in characterising the material sensitivity to reheat cracking [1-5]. This is in harmony also with the findings reported elsewhere [3–5]. The following conclusions were drawn: (1) The P23 steel CGHAZ was found ‘susceptible’ and the coarse-grained reheated B323B

multipass weld metal ‘extremely susceptible’, to reheat cracking. Introduction of the second thermal cycle producing grain refinement resulted in lesser cracking sensitivity.

(2) In the as-welded condition, the B323B weld metal exhibited considerably lower toughness than the CGHAZ of the P23 steel. All HAZ microstructures fulfilled the PED toughness requirement of 27 J at +20 °C even in the as-welded condition, whilst the weld metal fulfilled the corresponding requirement only in the PWHT condition.

(3) PWHT (760 °C / 2h) seems mandatory to guarantee weldment hardness not exceeding 350 HV, as well as to ensure adequate toughness for thick-section multipass welds made using B323B filler metal with broadly matching chemical composition to P23 steel.

(4) Welding techniques (e.g., bead placement, groove geometry) accentuating the normalising effect of multipass welding, in conjunction with intermediate heat inputs yielding the weld t8/5 of ≈10–20 sec, should be beneficial both in terms of reheat cracking and toughness.

3.2.1.3 – Welding trials As noted above, B323B was selected as a “state of the art” consumable – the best option available at the time of this selection – which was used for long term testing and characterization by the project partnership. ISQ optimised and produced all the welding samples for long term testing using the selected B323B consumables, as well as weld pads made with commercial consumables and alternative variants used for comparison testing and characterization. The Manual Metal Arc Welding and Submerged Arc Welding processes were selected as appropriate for the dimensions of tube and plates required for testing. These are the most common industrially used processes and their optimisation and qualification in the project represent major results in industrial terms. The optimisation of the welds has been carried out following a sequence of trials for each consumable, depending on the problems encountered in each case. The choice of welding procedure has been supported by technical experience, visual analysis, X-ray tests and in specific cases more extensive tests. A PWHT condition of 760ºC for 2 hours was selected as standard for this weld metal, in view of the ambient temperature ductility issues. Figure 3.2.4 shows the weld geometry and Figure 3.2.5 the Pipe G welded joint after PWHT and cutting for testing. Table 3.2.4 summarises all the welds produced by ISQ.

34

Item Nº Type Dimensions Welding

direction Consumable PWHT

Item 1a Deposit on flat plate

160 (welding direction)*110*35

SMAW B232 2h / 760ºC 1 Specimen

Item 1b Deposit on tube (219*32)

200 (along the tube axis)*100*25

Along the tube axis

SMAW B232 2h / 760ºC 1 specimen

Item 1c Deposit on tube (219*32)

200 (around the tube)*100*25

Around the tube axis

SMAW B232 2h / 760ºC 1 specimen

Item 2 Butt welds on tube (89*17)

-- -- SMAW B323 2h / 760ºC 4 specimen (3 real test + 1 characterization)

Butt weld on tube (219*32)

-- -- SMAW B323 2h / 760ºC 1 specimen Item 3

Butt weld on tube (219*32)

-- -- SMAW B323 Part as-welded Part 2h / 715ºC

1 specimen

Item 4 Deposit on tube (219*32)

150 (along the tube axis)*250*25

Along the tube axis

SMAW B323 2h / 760ºC 2 specimen

Item 6 Butt weld on plate

150*300 (welding direction)*50

-- CROMOCORD E223 – Air Liquide

2h / 740º 1 specimen

Deposit on tube (219*32)

100 (along the tube axis)*40*20

Along the tube axis

P23 Sumitomo 1h / 715ºC 1 specimen Item 7

Deposit on tube (219*32)

100 (along the tube axis)*40*20

Around the tube axis

P23 Sumitomo 1h / 715ºC 1 specimen

Deposit on tube (219*32)

100 (along the tube axis)*80*16

Along the tube axis

SMAW Bohler-Thyssen

2h / 740ºC 1 specimen Item 8

Deposit on tube (219*32)

100 (around the tube axis)*80*17

Around the tube axis

SMAW Bohler-Thyssen

2h / 740ºC 1 specimen

Item 9 Butt weld on tube (219*32)

¼ of the tube -- KA/06/0150 (B322 + CVNbNi)

2h / 760ºC 1 specimen

Item 10 Butt weld on tube (219*32)

¼ of the tube -- KA/06/0151 (Sumitomo + Nb)

2h / 760ºC 1 specimen

Item 11 Butt weld on tube (219*32)

¼ of the tube -- KA/06/0152 (Sumitomo - Nb)

2h / 760ºC 1 specimen

Item 12 Butt weld on tube (219*32)

¼ of the tube -- KA/06/0153 (Thermanit P24 +Nb)

2h / 760ºC 1 specimen

Table 3.2.4 – Welding samples produced during the project.

Figure 3.2.4: MMA weld geometry. Figure 3.2.5: Pipe G welded joint 3.2.1.4 – Welding consumable development and short term creep testing The limited information available prior to this project on high temperature performance of P23 and P24 weld metals indicated substantial concerns. Stress relief cracking can occur in P23 weld

10º

R5

2.0

2.0

32

200

35

metal and heat-affected zone (HAZ) during post-weld heat treatment (PWHT) [4]. P24 consumables also showed much poorer creep performance than parent P24 steel [6]. There were indications that the fundamental problem could be low weld metal creep ductility, such that cracking can occur at the very low creep strain levels, around 0.5-1%, associated with the relaxation of welding residual stresses on PWHT and/or in service. There was also evidence that weld metals can show anisotropic creep properties, with inferior performance when tested in transverse orientation, across rather than along the direction of welding. However, there was very little information on what causes low creep ductility, whether changes in chemical composition could improve weld properties, or how the interplay between creep strength and ductility affects overall performance. The main aim of weld metal development, section 3.2.1.1, was therefore to improve high temperature performance. Creep testing was carried out to investigate the effects of chemical composition and process variables, and hence develop improved consumables. Consumable development and testing Two somewhat conflicting requirements arose. First, it was necessary to choose a best available “state-of-the-art” consumable at an early stage of the project, and use the selected consumable to make full scale welds for characterisation, long term uniaxial testing, and pressurised component testing. Secondly, it was necessary to investigate a wider range of trial consumable formulations and process variants during the course of the project, with the aim of identifying a better option. Inevitably, therefore, the risk had to be accepted that test data obtained with the original “state-of-the-art” consumable might not be entirely applicable to improved product/s developed by the end of the project. Metrode plc were subcontracted by ISQ to produce small experimental batches of manual-metal-arc (MMA) welding consumables, based on P23 and P24, but with variations in chemical composition to explore and optimise creep behaviour while achieving acceptable microstructures, tensile properties and impact toughness. Each batch was used to make small AWS type welded test coupons suitable for mechanical testing and manufacture of 2-3 longitudinal creep specimens. Five small experimental series, each involving 2-4 new variant consumable batches, were produced over the course of the project. These were supplemented by tests involving MMA and TIG consumables obtained from Sumitomo and Bohler-Thyssen and used to make weld deposits at ISQ, and by testing submerged arc welds made from Metrode and Air Liquide consumables. An exploratory approach was adopted, with creep test results from early batches used to decide on the compositions manufactured in later batches. This had the advantages of matching the availability of consumable production facilities, and allowing test results to guide selection of subsequent compositional variables. For economy and efficiency, short term creep tests were used to rank the weld metals. For each consumable, all-weld-metal tests in the longitudinal orientation were carried out under two standard test conditions, each chosen to produce failure in around 1000h. Tests at 80MPa, 660°C provided data at a moderate stress of reasonable relevance to plant conditions, while tests at 130MPa, 620°C provided a cross-check with less severe temperature acceleration. Crosshead creep strain data provided an adequate measure of minimum secondary creep strain rate, r, in each test. Creep ductility was characterised by means of the Monkman-Grant parameter M= r x t, where t is creep life. The M parameter provides an approximate measure of the creep strain which the material can tolerate prior to cracking. Then, creep life (t) is the product of creep strength, as characterised by the inverse creep rate (1/r) at a given stress and temperature, and creep ductility, as characterised by the parameter M. Development of improved weld metals

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An initial test series on P24 weld metals compared a standard Metrode consumable with a cobalt-bearing variant from the earlier RFCS HITEPR project and a low-carbon weld metal from the Smartweld [6] project. In parallel, the initial test series on P23 weld metals compared a standard mid-range weld composition (Metrode B322) with a highly alloyed variant (Metrode B323). The latter was designed to maximise the content of the creep strengthening alloy elements C, Nb and V within the standard P23 specification range, and also to add B and Ti with the aim of improving grain boundary cohesion and thus enhancing ductility. All the P24 consumables showed very poor creep lives in the range 100-300 hours, with mediocre M values of the order of 1%. The two P23 consumables showed superior creep lives in the range 600 -1200 hours, though with poor M values, typically 0.6%. The data also indicated that B323 achieved a creep life improvement of the order of 40% compared to B322, without substantially affecting M, although its creep ductility at failure was poorer. These results contributed to the decision to concentrate work on steel P23. The decision was also made to select B323 as the “state of the art” project consumable, section 3.2.1.6. A third test series with Metrode trial consumables investigated the effects of separately adding either the “creep strengthening” elements C, Nb, V or the supposed “grain boundary cohesion” element B protected from oxidation by co-alloying with Ti. This showed that whilst the C/Nb/V addition is indeed creep strengthening, the Ti/B addition actually causes weakening, while none of these variations greatly affected creep ductility. A fourth series examined nickel additions to base B322 and B323 compositions. This produced a significant improvement in creep ductility, to around 2% for the lower alloyed consumable based on B322 and for a similar consumable from Sumitomo. However, these lower alloyed consumables were also fairly weak, whereas the higher alloyed variant based on B323 showed lesser improvement in ductility. A hybrid consumable from Bohler-Thyssen, “Thermanit P24”, also showed attractive properties. This consumable, whilst alloyed with Mo instead of W to match P24 parent material, also used niobium (based on the P23 specification) in place of titanium (from the P24 specification) as a carbide forming creep strengthener. It is understood that this formulation was originally designed simply to avoid using titanium, which transfers poorly across the arc, in a welding consumable. However, the resulting creep properties were much superior to standard P24 consumables, again showing the importance of niobium. Compared with the more conventional P23 variants, Thermanit P24 showed broadly comparable creep properties to the Sumitomo-type nickel alloyed consumable, with an acceptable if not particularly high creep strength, combined with relatively good ductility values of around 2%. A fifth and final experimental series aimed to develop high creep strength combined with high ductility, but this proved elusive, Section 3.2.2.3. Niobium was shown to be a key alloying element, with contents above about 0.05% liable to produce low creep ductility, whereas contents below about 0.04% were liable to produce low creep strength. However, attempts to improve on the “Sumitomo” and “Thermanit P24” formulations were not successful. With a “Sumitomo” type nickel-bearing P23 weld metal, higher niobium impaired ductility, while lower niobium produced a ductile but weak weld metal. With the “Thermanit P24” type of formulation, added nickel merely caused weakening. Trials also indicated that varying PWHT temperature in the range 715-760°C had little influence on creep performance, while TIG and submerged-arc welding produced creep data generally in line with MMA. Comparison tests on P23 parent materials, using the same short-term standard test conditions, however produced very different results. The creep lives were quite similar to those of the weld

37

metals, but both creep rate and creep ductility were much higher, with typical M values around 3-5%. Further analysis is given in section 3.2.2.3. Summary The project did not identify a single outstanding optimum consumable formulation with both high strength and high ductility. However, it did indicate routes toward improvement compared with the baseline B323 composition. Increased creep strength and life, but with low ductility, can be produced with high niobium levels. More promisingly, improved creep ductility can be achieved, either via nickel additions, or by alloying with Mo in place of W, in both cases excluding Ti and adding a carefully controlled moderate content of Nb. 3.2.1.5 – Short term creep testing of welded joints The work described here covers short-term creep testing of the main project weld metal B323B and alternative variants in transverse orientation. First, all-weld-metal transversely oriented specimens were machined from weld pads, made by ISQ by weld build-up deposition onto the outer surfaces of P23 pipe segments. These all-weld-metal specimens were tested to ensure failure in weld metal and thus characterise and compare its behaviour in transverse and longitudinal (section 3.2.1.4) orientations. Secondly, cross-weld test specimens were machined from full scale pipe circumferential butt welds made between lengths of pipe G by ISQ. These specimens, each spanning part of the weld, HAZ, and parent material, provided data on the performance of the complete welded joint. It was envisaged that when creep failure occurs in the weld metal, an all-weld-metal transverse test and a cross-weld test would be likely to produce similar data. However, this proved to be a false assumption. The four weld metals transversely tested as all-weld-metal deposits were the strong and highly alloyed project consumable B323, mid-range composition B322, and the more ductile weld metals from Sumitomo and Bohler-Thyssen, section 3.2.1.4. For each weld metal, the two “standard” short term tests were again carried out. Each consumable could then be simply characterised in terms of a life ratio parameter R, the total creep life (summing the two test results) in transverse testing expressed as a percentage of the total creep life in longitudinal testing, section 3.2.1.4. The results varied from 52% for B323 and 57% for B322 up to 64% and 77% for the more ductile Sumitomo and Bohler-Thyssen weld metals. Thus, transverse performance was somewhat inferior, but did not appear to be a major concern. Very different results were obtained when B323B was tested via cross-weld specimens. First, the two “standard” tests were carried out using specimens from the outer part of the weld, near the pipe OD. Both failed in the weld metal with very short test lives, the R ratio being 19% for the 660°C test and only 6% for the 620°C test. Tests were then carried out using specimens from the inner part of the weld, nearer the pipe ID. One failed in parent material, while the other (at 660°C 80MPa) failed in the HAZ and at more than double the life of the outer-weld specimen. The most probable explanation for the unexpectedly poor results is that the B323B pipe butt weld suffered some weld metal stress-relief microcracking in its outer regions, so that the outer-weld specimens contained defects when put on test. These defects (which, as creep-related defects, would have formed by creep cavitation at grain boundaries) then rapidly propagated to cause failure when creep tested. Whilst the B323 and B323B all-weld specimens also received a PWHT, these relatively low-restraint surface build-up deposits will have been subject to lower residual stresses and less triaxiality, and thus at lower risk of stress-relief cracking. Cross-weld tests were also carried out on the four “fifth series” weld metals, section 3.2.1.4, which were used to make separate quadrants of a pipe butt weld which was PWHTd before sectioning. Two moderately ductile weld metals (M values 0.8-1.2) gave R ratios of 92% or higher, while two weak but more ductile consumables (M values around 2% and higher)

38

produced R ratios of approximately 140% and 190%. In the latter two cases the results were unexpectedly good, perhaps because the weaker weld metals were strengthened by dilution with parent material. Generally, therefore, these data give confidence that if a reasonably ductile weld metal is chosen in place of B323B, transverse weld metal creep properties will not be a specific concern. Because these short term failures mainly took place in the weld metal, however, they provide little information on the possible risk of longer term HAZ “Type IV” failure. 3.2.1.6 – Selection of the more promising consumables and parameters The critical consumable selection decision in this project was the choice of the “state-of-the-art” weld metal for the main test programme. An early choice between the variants developed at that stage was essential to obtain long term weld creep data, section 3.2.1.4. The initial data clearly showed that standard P24 weld metal would be weak. However, just two P23 variants had been tested by month 12°, when a choice had to be made: the mid-range composition B322, and the stronger and more highly alloyed variant B323. Metrode found problems in obtaining adequate mechanical properties data for B323. Whereas B322 showed reasonable ambient temperature impact and tensile properties after post-weld heat treatment (PWHT) at 715°C, B323 showed higher tensile strength (779 MPa) with impact and tensile ductility values below acceptable limits. PWHT at the commonly used temperature of 740°C improved impact properties, but PWHT at 760°C was required to achieve a tensile ductility just exceeding the 15% level required by consumables acceptance standards. As the recommended steelmaking tempering temperature for P23 [3] is itself 760°C +/- 15°C, this PWHT temperature is arguably undesirable, in that its overlap with the tempering temperature is liable to weaken the steel close to the weld. However, it is noted that this is common practice with the well known steel P91, and has not apparently caused any practical problems. B323 did offer improved creep life, though with poor ductility. Interrupted creep tests at about 0.75% strain showed only minor (0.05mm) microcracking in B322, but a larger (0.5mm) area of cracking in B323. At this stage, therefore, neither could confidently be declared to be at nil risk of stress-relief cracking. The decision was made to choose B323, with PWHT at 760°C for 2h, as the project weld metal. It was based on two considerations: first, that the higher creep strength could be critical, and secondly, that there was actually merit in investigating a high-risk consumable. On this argument, an R&D project (rather than service application) is the best way to gain experience with a consumable which might either produce improved performance, or, prove unsuitable. Metrode therefore produced a new batch, B323B, to the original B323 specification. The B323B batch was used for all long term creep testing (section 3.3), low cycle fatigue and creep crack growth testing (section 3.3), reheat cracking simulation (section 3.2.1.2), and pressurised tube testing (section 5). Inevitably, it was not quite chemically identical to B323, though the creep data were very similar (section 3.2.1.4). However, B323B did not achieve the required 15% ambient temperature elongation even after 760°C PWHT, again casting doubt on its practicality. ISQ used B323B to weld in several different geometries, including build-up deposits on plate and pipe OD surfaces to produce all-weld-metal creep specimens, and two full-scale girth welds in P23 pipe G, one of which was left as-welded for later PWHT (where required) after sectioning into smaller pieces. In hindsight, the girth welds in pipe G and H which were PWHTd before sectioning would have incurred the greatest creep strains on PWHT, and would thus have been the most likely to suffer stress-relief cracking. Poor experience with B323B in the creep test programme is discussed in section 3.2.1.5, while the identification of probable weld metal stress relief cracking in CT specimens by VTT and in a

39

failed pressurised tube test by IWM are recorded in sections 3.3 and 3.5 respectively. The erratic variability of the cross-weld creep test data, showing a mixture of reasonable and very poor performance, is therefore probably due to the presence of stress-relief cracking and/or damage in some parts of the B323B weld deposit. While this has reduced the value of the long term test data, it has also provided a key result, in that it serves to highlight the necessity to use a more creep ductile weld metal, preferably with creep ductility (M) value well in excess of 1%. Fortunately, two lines of development do both produce weld metals with better creep ductility, sections 3.2.1.4 and 3.2.1.5. These should supersede B323B. Both these types of weld metal appear suitable for application. For the future, longer term testing should be carried out to clarify whether adequate creep ductility is maintained under realistic PWHT and plant service conditions, and to select a single optimum chemistry. Task 3.2.2: Welding qualification 3.2.2.1 – Microstructural analysis of welding in Heat Affected Zone and weld metal The macro section in figure 3.2.6 shows the multipass weld in pipe H including the weld metal, the heat affected zone and the base metal after a post-weld heat treatment at 760°C for 2 hours. The weld metal presents a columnar dendritic microstructure in all passes except the weld root (Figure 3.2.7A). The base metal presents a fully tempered bainitic-martensitic microstructure (Figure 3.2.7C). The coarse-grained, fine-grained and intercritical reheated HAZ microstructures are also shown in Figure 3.2.7 A to C. A Vickers hardness HV1 profile through all material zones of the weld was measured in mid-wall position, figure 3.2.6. Considering the weld metal and the base metal hardness levels, the weld is overmatched. The minimum hardness was measured in the sub-critically reheated heat affected zone.

Figure 3.2.6: Macro section and hardness profile of pipe H weld after PWHT

Figure 3.2.7. A – weld metal and coarse HAZ; B – coarse and fine HAZ; C – base metal

3.2.2.2 – Qualification of the welded joint The welding procedures optimized in the project were established and approved by the welding procedure tests defined in the Standard EN 288-3 [7]. These included visual inspection, NDT by

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final liquid-penetrant, radiographic examination before and after PWHT, longitudinal weld tensile tests, all weld bend tests, Charpy V-notch impact tests, macro examination and hardness tests. 3.2.2.3 – Strain monitored creep testing of weld metals. Section 3.2.1.4 gives a qualitative description of the weld metal creep test programme in terms of consumable development. This section provides a quantitative review and analysis of the creep strain data. Figure 3.2.8 shows the relationship between minimum creep rate r and the Monkman-Grant ductility parameter M, for all the (longitudinal and transverse) all-weld-metal short term creep tests at each test temperature and at both temperatures. Corresponding data for parent materials are also plotted. In general, there is a strong inverse correlation between creep strength and creep ductility. This is not entirely unexpected, since factors causing strengthening can often promote brittleness. Thus, when the matrix is strengthened, greater mismatch stresses may be imposed on grain boundaries, leading to cavitation or decohesion. The level of correlation is high, showing that within the compositional ranges investigated, there is only limited scope for strengthening without impairment to ductility. Nevertheless, some weld chemistries are better than others. For brittle weld metals with M values below 1, there is a substantial spread in r at constant M. Unfortunately, for the more ductile weld metals with M values above 1, there seems to be less scope to increase creep strength while maintaining reasonable ductility. The two different test conditions produce very similar data, Figure 3.2.8. On average, the low stress 660°C test produces a 12% higher creep rate with an almost identical average M value, leading to a 12% lower average creep life. The 620°C test data do also show somewhat more

Figure 3.2.8.

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variability in both creep rate and ductility, so that this test emerges as being marginally the more discriminating. However, the “all data” graph of Figure 3.2.8, in which tie-lines have been drawn to join the two data points from each consumable, shows that performance in the two tests is well correlated. For example, a consumable which is strong and brittle at 660°C, 80MPa is, in almost all cases, quite similarly strong and brittle at 620°C, 130 MPa. It might have been argued that the 660°C test examines resistance to severe thermal degradation during the test, while the 620°C test might be more influenced by initial as-manufactured weld strength. However, there is little indication that any consumable performs markedly better in one test than in the other. Most of the tie-lines are short, suggesting that the differences between linked points are primarily due simply to a normal level of random test variability. It follows that we can have increased confidence in the overall test programme, in that neither a potentially excessive temperature (660°C) nor an unrealistically high stress (tests at 620°C) seems to have caused any particular effect on behaviour.

Figure 3.2.9: All-weld-metal creep test data. Creep life plotted against ductility. The data are replotted, Figure 3.2.9, in terms of creep life t and Monkman-Grant ductility parameter M. This indicates that for the best weld metals, there appears to be an approximate performance “limit line” which represents an inevitable trade-off between creep life and ductility. The data on improved weld metals developed in this project lie close to the line, but if a M value (ductility) substantially above 1% is required to minimise the risk of stress-relief cracking, then it is not also possible to achieve the maximum creep life. It is, however, possible to match the parent creep life. This may well be sufficient for good service performance. Other less well formulated P23 weld metals, such as the original B322 and B323, produce data which lie further below the limit line. The standard P24 consumables produce inferior data, being very weak but with only mediocre ductility. Notably, all the parent materials produced superior data, in that the points lie above the limit line. Whereas parent material creep life and creep rate appear to be broadly comparable with a mid-range weld metal composition such as B322, the parent material creep ductility is nearly an order of magnitude higher. The most probable explanation is, broadly, that the weld metal

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microstructure is inherently more liable to creep failure at low ductility than a parent material microstructure. The weld metal has an inhomogeneous and anisotropic microstructure, promoting anisotropic creep strain accumulation: it may be strengthened and embrittled by strain incurred during the welding cycle and on PWHT: and its coarse columnar structure is liable to promote creep cracking at the columnar grain boundaries. For the future, better modelling and understanding of the key factors could enable further developments in welding consumables and processes to improve weld metal creep ductility. 3.2.2.4 – Selection of most promising material and weldments Welding process investigation, consumable development, and the selection of a state-of-the-art MMA welding procedure have been discussed in sections 3.2.1.3 to 3.2.1.6. This section details the manufacture of project test welds to the state-of-the-art procedure. Pipe welding

Figure 3.2.10: pipe G welded joint: specimen sampling

Plate Welding The test coupon plates were delivered to ISQ in the as manufactured condition and milled to the dimensions shown in figure 3.2.11. A double V weld preparation was selected.

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Figures 3.2.12 and 3.2.13 illustrate the weld preparation and submerged arc welding process.

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Figure 3.2.12: The double V weld preparation.

Figure 3.2.13: Submerged arc welding.

During the welding process, regular non-destructive testing was carried out. Liquid-penetrant and. X-ray examination were performed after welding to the middle thickness. The liquid-penetrant inspection identified some cracks (figure 3.2.14), which were then removed, and after confirmation of sound material, the welding process progressed. After the side change and the first pass had been deposited, a crack appeared in the root region, covering the full length of the weld but only within that pass. In the final two top passes, liquid-penetrant inspection found no defects (figure 3.2.15), and the repair interventions were sufficient to eliminate all the cracks detected during the welding process. No defects were found in the X-ray examination.

Figure 3.2.14: Image of the liquid-penetrant cracking observation.

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Figure 3.2.15: Image of the liquid-penetrant final faces without defects

The properties of Base Metal, HAZ and SAW Weld metal were determined after a PWHT of 2Hrs at 770°C. The cross section of the welded joint is shown in figure 3.2.16. Cross weld hardness data are summarised in figure 3.2.17. The maximum hardness is about 230HV.

Grade 23 : SAW Weld 770°C / 2h / airHV5 hardness cross weld

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Table 3.2.5: Cross weld tensile test results The welded plate was also characterised by impact Charpy V tests after PWHT with notch locations in the base material, weld metal, and heat affected zone (HAZ). The lowest FATT appears in the HAZ, as expected, but the reduction with respect to the base material is only 17°C (figure 3.2.18).

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Grade 23 - Joint SAW ISQ - BM - 770°C / 2Hrs

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3.2.2.5 – Long term cross-weld and all-weld metal creep testing The weld creep rupture test programme produced data out to 10,000 hours at three temperatures, 550, 600, and 650°C. LCF endurance curves have also been produced on all-weld, cross-weld and parent P23. Results are presented in section 3. 3.2.2.6 – Development of dissimilar welding As explained in the introduction of this report, the consortium agreed to focus its activities to achieve the best results on similar Grade 23 joints. Conclusions of Work Package 2 1) Welding procedures have been developed and optimised for manufacture of thick section P23

components. An extensive range of weld metal compositions have been investigated and a comprehensive series of full scale test specimens have been produced.

2) The overmatching, highly alloyed “state-of-the-art” P23 weld metal B323B, selected for its high creep strength and used for the manufacture of long term high temperature test specimens, proved to be a poor choice. This creep-brittle weld metal was found to be liable to reheat cracking during PWHT of full scale thick section butt welds. The occasional presence of pre-existing weld reheat defects in the test specimens therefore led to a pattern of inconsistent and unreliable behaviour.

3) B323B was deliberately chosen as a test of the viability of the high-strength formulation. The negative result is valuable, in that it clarifies the pitfalls that can occur and indicates a compositional range to be avoided.

4) Parallel work using Gleeble weld thermal simulation together with the BWI tensile reheat cracking test showed that the P23 weld HAZ is susceptible to reheat cracking, while confirming that B323B weld metal is extremely susceptible.

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5) Later trials showed that alternative compositions, involving slight deviations from the P23 specification, could be developed to produce substantially less creep-brittle weld metals. These successfully survived PWHT without reheat cracking.

6) The results indicate that the Monkman-Grant creep ductility parameter M provides a reasonable guide to P23 weld metal performance, and that M values above 1% should generally be adequate.

7) Improved weld metals, with M values of the order of 2%, can now be formulated. These should be subjected to long term testing in future work.

8) P23 weld metals nevertheless show intrinsically poorer creep ductility than P23 parent materials. If adequate ductility is to be achieved, some sacrifice in terms of weld metal creep strength appears to be unavoidable. The microstructural reasons for this are not yet clear.

9) The risks of reheat cracking in P23 weld metals and heat-affected zones can thus be reduced by careful selection of consumables and welding procedures. However, the inherent susceptibility of the P23 composition is a disincentive to its selection for thick section power plant components.

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3.3 WP 3 - High temperature testing and data assessment Objectives : • Provision of high temperature mechanical data assessment for base material • Provision of high temperature mechanical data assessment for welded joint • Determination of weld stress reduction factors The output of WP3 is towards WP5 High temperature design and assessment Introduction The main objective of AloAS WP3 is to assess the mechanical high temperature behaviour of the new heat resistant low alloy steels T/P23 and T/P24. The material properties for high temperature service are determined mainly by destructive material testing, such as tensile, creep and fracture toughness testing. Special effort is also put into determining the properties of weld metals (consumables) and the cross welds of the project steels. The material data is assessed with state-of-the-art assessment models and tools. The resulting models are then to be applied in WP5 for selected cases. The mechanical testing, modelling and simulation work of the AloAS project has produced a sound basis for life management of P23/24 steels and the development of their weldments. The collation of mechanical property data has been finalised and the final assessments have been performed. However some creep tests are still running and will be continued by the partners in order to increase the database. The results will be reported separately in international conferences and scientific publications. The main observations from the WP3 subtasks are presented in Table 3.3.1. Note that creep data were also produced in subtask WP2. The materials testing programme has been performed both at moderate (550-600°C) and high temperature regimes (above 600°C). The higher temperatures tests were taken into the testing programme due to the high stress levels needed for sensible durations in creep testing at the intermediate temperatures closer to plant service. At high temperatures the stress levels become relevant for plant operation conditions. The realistic plant stresses are estimated from design codes, based on extrapolated creep rupture data with a safety factor. For T23 steel tube, oxidation may limit application to about 575°C, but for P23, higher temperatures could be viable. The recent ASME Code Case 2199-3, 2006 0 interpolated design stresses are 70MPa and 53MPa at 575°C and 600°C respectively. It was concluded that the creep test programme should include test temperatures enabling data to extend down to around 50–70 MPa. Task 3.3.1 Determination and comparison of mechanical properties A major effort has been made in acquiring high temperature creep data (see 3.3.1.2) needed for the long term life management of the project steels and their welds. In addition the following high temperature fracture mechanical tests have been conducted; Fatigue crack growth tests (FCG, see 3.3.1.1a), low cycle fatigue tests (LCF, see 3.3.1.1b) and creep crack growth tests (CCG, see 3.3.1.3). The steels have been tested as-manufactured. 3.3.1.1 Fatigue crack initiation and crack growth tests Low cycle fatigue (LCF) testing was carried out on P23 steels B, H and ex-service CrMoV steel KA1358 as well as all-weld-metal types B323B and B322, and cross-weld specimens from P23 pipe G welded with B323B. Fatigue crack growth tests were also performed on steels B and KA1358.

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Task Type Materials Main observations

2.1.4 Short term creep (WM) See WP2 B323 longest creep life, selected as main consumable B322 more ductile

2.1.5 Short term creep (CW) See WP1 mainly fracture in BM WSF difficult to acquire

3.1.1 Fatigue crack initiation and crack growth tests

B, H, weld metal, cross-weld, CrMoV (> 200 kh service)

P23 parent satisfactory B323B weld poor, but B322 better

3.1.2 Medium / Long term creep A, B, G, H, J, N, P (T/P23) C, E (T/P24)

weaker than EN10216 [9] acc. VdTÜV 533 [10]

3.1.3 Creep crack growth tests weld metal and HAZ for welded G tube

Creep cracks follow curved HAZ in WM (Stress relief cracks found)

3.1.4 Comparison testing welds and HAZ

Consumable testing, Simulated HAZ testing

Differences between testing direction of welds

Successful production of simulated HAZ microstructures

3.1.5 Comparison; literature data and project material heats -

Comparison to EN10216, VdTÜV 533 and for BM and with Error!

Reference source not found. for CW

3.2.1 Weld creep assessment H, G, P WSF predictions performed, mainly BM fractures, large scatter for H

3.2.2 Data base compilation Aloas materials (BM, WM and CW)

Data base for creep assessments compiled. Assessments performed.

Some tests still running

Table 3.3.1. ALoAS WP3 test, materials and main observations in the order of task. BM = base material, WM=weld material (consumable) and CW = cross weld.

Steel B was tempered at 760°C and had a UTS of 675MPa, but H had been re-normalised and tempered at 780°C to improve its impact properties and had a UTS of 620 MPa. Originally H had been normalised and tempered at 760°C, in which condition its UTS was 694 MPa. Its higher tempering temperature had thus substantially reduced its strength. A third material G was also tempered at 780°C and similarly had a rather low UTS value, 635 MPa. The ex-service CrMoV pipe KA1358 was supplied by E.ON from a UK plant main steam bend and had seen 217,921 hours of service at 565°C with 1378 starts. The compositions of these materials are given in WP1. 3.1.1a Low cycle fatigue testing Plain low cycle fatigue (LCF) specimens of 7 mm diameter were tested at a strain rate of 2.5%/min under fully reversed strain control (Rε, ratio of minimum to maximum strain, = -1). All loops were continuously monitored and the data collected. B and KA1358 were tested at 565°C, while later tests on H were carried out at 575°C for conformity with the pressure tube tests at IFW (see WP5). Further tests on H explored sensitivity to temperature variation at 450°C (representing pipe thermal downshocks) and 600°C. Tests at 575°C were also carried out on B323B weld metal from two separate pad samples deposited on the surface of pipe G, with test specimen axes longitudinal and transverse with respect to the welding direction. The latter simulates “bore” cracking in a pipe butt weld, since the LCF crack itself then grows along the welding direction. Comparison tests were also carried out on the weaker B322 weld metal and on two cross-weld specimens from pipe G. The number of cycles to crack initiation Ni was determined as the departure from linearity of the plot of maximum stress in the cycle versus the number of cycles. For each test the stress range and plastic strain range were determined for the mid-life loop and used to derive the cyclic stress-strain relationship.

49

LCF tests with dwells were also performed on B and H materials and B323B and B322 weld metals at a total strain range of 1.0%, with strain controlled dwells of 5 and 60 minutes at peak tensile strain. Test conditions and results are given in the assessment section 3.3.2.2. 3.1.1b Fatigue crack growth Testing Fatigue crack growth (FCG) tests have been performed for material B and KA1358 with 0.8 compact tension (CT) test specimens with width, W = 40mm and thickness, B = 20mm. The specimens were manufactured with a spark eroded slot to an a/W value of 0.4 (where a is crack depth) and then fatigue pre-cracked in air at a frequency of 50 Hz and load ranges of between 9 and 10kN. The fatigue pre-crack depth was around 2.5mm and took more than 105 cycles in each case. The specimens were manufactured such that the crack plane contained the radial and circumferential directions in pipe B, and the radial and axial directions in KA1358. The FCG tests were carried out at a temperature of 565°C and frequency 1Hz, at a constant load range, with crack monitoring via crack mouth clip gauge. The cracks were grown to an overall depth of approximately 26mm. After testing, specimens were sectioned at 0.25B for metallography, and the remainder broken open in liquid nitrogen for measurement of crack depth. Results are given in the assessment section 3.3.2. 3.3.1.2 Long term creep testing Creep testing for the Aloas steels and welds were conducted in the creep laboratories of the partners. The base material (BM) creep data produced is presented in Figure 3.3.1 and Figure 3.3.2.

40 60 80 100 120 140 160 180 200 220 2400

5

10

15

20

25

Num

ber o

f tes

ts

Stress (MPa)

540 560 580 600 620 640 6600

5

10

15

20

25

30

35

Num

ber o

f tes

ts

Temperature (°C)

Figure 3.3.1: Distribution of BM tests according to stress and temperature.

A B G H J P Z0

5

10

15

20

25

Num

ber o

f tes

ts

Heat designation

0 5000 10000 15000 200000

10

20

30

40

50

60

Num

ber o

f tes

ts

Testing time (h)

Figure 3.3.2: Distribution of tests according to material heat and testing time (Z=other data

approved for use within the AloAS project). 50

For the BM data set the time range is between 23.0 and 16600 h in time for ruptured tests and beyond 20.000 h for running tests (11 running). The stress range for the BM creep tests is between 50.0 MPa and 225.0 MPa. The temperature range is between 550°C and 660°C. For T24 two heats have been tested (C, E). The tests were performed at temperatures 550-660°C with a maximum rupture time of 13.250 h. For the weld metal (WM), for consumable selection, tests have been conducted at three temperatures 575°C, 620°C and 660°C with a total of 65 data points covering 37 different consumables and testing directions (longitudinal and transverse over the weld). The maximum testing time for these tests was 1985.0 h. The stress range is 80 - 152 MPa. In addition to these tests some all weld tests were performed on weld metal from weld pads. For these the test temperatures were 550, 600 and 650°C and the stress range 80 - 170 MPa. The longest creep test is still running at 9500 h. For the cross weld (CW) configuration, 37 creep tests were performed for three P/T 23 heats (G, H and P). The temperature range is 550°C-660°C, stress range 60-225 MPa with the longest ruptured test at over 5000 h (G). The longest running test is over 6366 h. The data distribution is shown in Figure 3.3.3.

40 60 80 100 120 140 160 180 200 220 2400

2

4

6

8

10

Num

ber o

f te

sts

Stress (MPa)

540 560 580 600 620 640 660 6800

2

4

6

8

10

12

14

Num

ber o

f tes

ts

Temperature (°C)

Figure 3.3.3: Distribution of CW tests according to stress and temperature.

3.3.1.3 Creep crack growth test For the creep crack growth testing (CCG) three specimens were extracted from the welded P23 pipe G such that the starter notch was on the HAZ and the direction of the crack growth was from outer diameter inwards. Also one full weld metal specimen was tested. The selected specimen type was W25 mm CT with machined 10% side grooving (both sided). The thickness of the pipe was not quite sufficient for the full W25 mm specimen and therefore a small layer of X20 was welded on top of the pipe by GTAW (preheating at 200ºC, interpass temperature 300ºC). The post weld heat treatment (PWHT) was conducted at 740ºC (2h) for one HAZ specimen and the WM specimen. The two other HAZ specimens were PWHT at 760ºC for 2h. The EDM starter notch was eroded with r =0.07mm. The CCG tests were conducted in a dead weight creep machine in air. The crack length was measured with a PD device using a current of 10A. The polarity of the current was reversed for each measurement in order to compensate for the thermocouple effect. The load line displacement was measured with an LVDT fitted to specially designed frictionless extension arms. Step loading was applied during the start-up in order to verify that the extensometers function as they should. The CCG test details and results are given in section 3.3.2.2

51

3.3.1.4 Comparison testing of welds and HAZ Weld performance in creep (Section 3.3.2.1) and LCF (Section 3.3.1.1) has been detailed elsewhere. Key findings are summarised here. The project selected a strong weld metal, B323B, with the aim of maximising creep rupture strength. This proved unsuccessful. The B323B weld metal, with a Monkman-Grant creep ductility below 1%, was found to be susceptible to reheat cracking on post-weld heat treatment. The indications are that this occurred especially in the most restrained weld which was subject to PHHT before sectioning, leading to some premature creep and pressure tube test failures, and producing very scattered data. Improved weld metals were developed with Monkman-Grant creep ductilities well in excess of 1%, and these should be tested in further research. B323B also exhibited poor room temperature tensile ductility and poor LCF properties. However, alternative weld compositions were readily found to avoid both of these latter problems. The HAZ did not appear to be specifically at risk of LCF cracking, but the mismatch between a strong weld and HAZ and the weaker parent P23 did lead to one failure in a pressure tube test (see WP5) and in the simulated girth weld with the same WM (see creep strain assessment in 3.3.2.2). Because premature weld metal failure took place, only limited data were obtained on the risk of HAZ creep failure. The indications are that P23, while not immune from the HAZ “Type IV” cracking problem which limits weld life in many ferritic steels, could be relatively resistant. This may merit further research. 3.3.1.5 Comparison with literature The base material creep performance of T/P23 has been compared with the European standard 10216-2 0 and published data on HCM2S 0. The T/P24 have identical values for strength to 10 000 and 100 000 h at specified temperatures in the VdTÜV 533 data sheet 0 and in the EN 10216. Data for the T/P24 has also been found in 0. The acquired creep strength of welds is compared with the published cross weld data of Belgian Welding Institute Error! Reference source not found., Vallourec Mannesmann 0 and the UK Thick section welded HCM2S project 0. Also some all weld (consumable) test results have been found in 0. Task 3.3.2 Data base and assessment Task 3.3.2.1 Weld creep data assessment The cross-weld creep strength of ferritic steels is traditionally assumed to be about 80% of the corresponding value for the base material. For some steels however cross-weld data assessments suggest that this does not hold at least at the high temperature end of the testing range. The weld creep strength factors (WSF) tend to fall below 0.8 when extrapolated to typical design life (100.000 h or more). Under such conditions the conventional value of 0.8 would result in a non-conservative (too long) predicted creep life for structures subjected to cross-weld loading.

To acquire well behaving WSF curves, a reasonable relationship between the weld behaviour and the base material master curve is needed. A mismatch of selected weld and base material models can produce e.g. unrealistic cross-over or turn-back in the extrapolation.

For the Aloas cross weld data two procedures for weld creep strength modelling have been applied. First the classical constant weld strength reduction factor has been calculated using heat to heat material strength as base. As a second approach the Rigidity Parameter Correction (RPC 0), method has been calculated for conservative predictions in extrapolation.

52

The welds tested are material H, G and P. Materials H and G were welded with the B323B test consumable and plate material P with a Air Liquide Welding (ALW) submerged arc consumable. The weld strength factor for these materials was calculated as a heat to heat assessment. The varying performance of the B323B welded test materials (H and G) makes the weld strength factor calculation for these rather unreliable. The cause of this is still under investigation, but the fact that reheat cracking have been found in both CT specimens (see creep crack growth) and in the weld metal for cross welds with B323B (see WP2, 3.2.1.5) suggests a preliminary explanation: that pre existing stress reheat cracks cause the premature failures with small failure strains and reduction of area values. It is evident from examination of the weld creep test data that a minority of specimens fail very prematurely. These specimens come from the weld which was post-weld heat treated as a full scale pipe butt weld and hence most likely to suffer reheat cracking due to its high restraint. Creep tests on weld pads, Section 3.2, show better performance, probably because reheat cracking was avoided in this lower restraint configuration. Consequently, the WSF calculations given here are inevitably rather questionable. For a better weld metal that does not suffer reheat cracking, the results of this work might be taken as indicating lower bound behaviour. The weld creep strength factor (WSF) is defined according to ECCC recommendations 0 as:

Ttu

TtWu

RR

WSF//

//)(= (1)

where TtWuR //)( is the predicted strength of the weld at specified time and temperature. The ECCC recommendations further define a convenient fitting efficiency parameter or scatter factor as:

PARAMZ ⋅= 5.210 (2) where PARAM is either the root mean square error (RMS or the standard error of estimate calculated on predicted and true logarithmic rupture times. Classical WSF calculation The weld strength factor for material P and G has been calculated by minimising the root mean square (RMS) for the specific weld in relation to the base material master curve by dividing the stress in the master curve by the WSF. The calculated WSF values are presented in Table 3.3.2.

Heat Consumable WSF Z – efficiency factor All B323B/ALW 0.86 3.44 – ACCEPTABLE P ALW 0.91 1.6 – GOOD G B323B 0.84 1.6 – GOOD H B323B (0.85 or 0.66) >5 – NOT ACCEPTABLE:

Table 3.3.2. Calculated weld strength factors for welded P, G and H. Note that H has low ductility fractures in some test. For material H no valid WSF can be calculated.

The WSF values obtained from the classical calculation perform well against the public domain data Error! Reference source not found.00. The V&M data has an optimal data fit for the weldments at a WSF equalling 0.88 at 550°C against the base material performance reported. The UK thick section HCM2S project welds tested at 575°C had a corresponding WSF equalling 0.81 (with the Chromet 23L consumable) against the same base material properties (since base material data was not reported). The Belgian short term cross weld data (<1000 h) showed inferior results for both the Cr2WV consumable (WSF=0.76 for SMAW and SAW welds) and

53

the AL CROMO SF 223 consumable (WSF = 0.71 for SAW) in relation to the parent material (V&M heat 73.220, 20% stronger than V&M average). It is however to be noted that the base material prediction lines differ, the HCM2S data being stronger in comparison to the AloAS materials. It is also to be remembered that the B323B weld performed in a brittle manner in some tests with the H material and when tested as all weld metal, especially in the transverse direction. The resulting scatter factor was therefore large and the more conservative WSF should be used. It is to be noted that the WSF calculated here is representative to the range of data only since there is so far very little evidence of failure location transition from base metal (or weld metal) to HAZ, which in most heat resistant steels causes a larger drop in the WSF for longer test durations. The RPC method The RPC method introduces conservatism in the WSF predictions, regardless of the used base material model. The RPC approach applies a similar non-linearity formulation as the master curve equation of the Manson’s minimum commitment method (MCM). The RPC states that a rigidity parameter R can be defined to transform the predicted rupture time of the base material to the corresponding cross-weld rupture time by bending it over a pivot point in time. The time transformation is defined as

)log())log()(log(1

)log()log()log( p

pm

pmR t

ttrRttr

tr +−⋅+

−= (3)

where trR is the corrected (RPC) time to rupture (here the welded material prediction), trm the uncorrected (base material master curve) time to rupture and tp the pivot point in time. The correction is zero at the pivot point and reduces the predicted life elsewhere

The term log(tp) can be fitted as a temperature dependent function or used (as here) as a constant. Note however that the pivoting point value when used for weld strength modelling is usually negative to accommodate the transition location strength drop. Any base material creep model can be then modified for the cross-weld data by minimising the root mean square error (RMS) and optimising the values of R and log(tp) that are given for material G and P in Table 3.3.3. The fitting efficiency values are very competitive in comparison to the classical method (see Table 3.3.4).

RPC parameters Heat

log(tp) R

P -34.16 3.14E-04

G -99.37 4.24E-05

Table 3.3.3. Optimised values of log(tp) and R for welded P23 steels P and G

Heat Consumable WSF (100kh, 575°C) Z – efficiency factor P ALW 0.71 1.54 – GOOD G B323B 0.73 1.89 – OK

Table 3.3.4 Calculated weld strength factors for welded P and G using RPC. The values are predictions for 575°C and 100 000 h. The efficiency factor is calculated on the data.

54

The RCP produced conservative predictions for the WSF. In the case of the Aloas welds (G and P) the predicted WSF at 100.000h / 575°C is lower than the classical WSF result. This is totally in line with what is expected from the RPC method since extrapolation beyond the range of data always produces a reduction in the predicted times to rupture and therefore in the WSF. As stated before for the classical WSF calculation also here it is to be remembered that the cross weld data assessed is quite limited and there are only very few failures in HAZ. This indicates that failure location transition has not yet been encountered. Therefore the RPC method could also be non-conservative, but much less so than the classically predicted WSF. 3.3.2.2. Data base Compilation In this section the creep test results are assessed and the results are presented for high temperature LCF and FCG and CCG described in 3.3.1.1a, 3. 3.1.1b and 3.3.1.3 respectively. 3.3.2.2a Creep rupture modelling The model development and assessments for creep rupture and creep strain has been performed on BM, WM and CW data. The results of creep strain modelling for the AloAS materials will be presented in the Creep 2008 conference in Bayreuth in May 2008. The final fine tuning of the models will be done when the still running data are available. Creep rupture: Base Metal The preferred model for the Aloas T/P23 rupture data is the Minimum Commitment master curve (MC 0) defined in Eq. (4),

log(tr)=β0+β1⋅log[σ0]+β2⋅σ0+β3⋅σ02+β4⋅T+β5/T (4)

where tr is the time to rupture T the temperature in K and σ0 the stress in MPa. The model constants β1-β5 are ;

β0 β1 β2 β3 β4 β5 45.3295 -2.6701 0.001439 -4.969E-05 -0.03233 -6611.06

The MC isotherms are plotted together with the BM rupture data in Fig. 3.3.4. For presentation purposes the MC rupture predictions are transformed to Larson-Miller parameter values in most result figures in this report, such as in Fig. 3.3.5 to show property differences of the material heats. It is to be remembered however that the MC model can only be presented in this form isotherm by isotherm. The MC model produces different slopes in stress vs. Larson-Miller parameter presentations at different temperatures. The data at high Larson-Miller parameters are mainly from high temperature tests and are therefore above the minimum commitment 600°C prediction line. Materials A, G and P are well represented by the MC equation. Material J has a different slope from the model at high stresses but converges with the model at long test durations or high temperatures. To enable testing at realistic plant stresses below 100MPa, testing at above 600°C was required to obtain medium term rupture data within the project timescale. An overlapping test matrix at 625°C and 650°C was employed and the results co-plotted using the standard Larsen-Miller time-temperature parameter. This showed no marked or consistent differences between the

55

results at different temperatures, indicating that testing at these high temperatures is not, in itself, liable to cause major errors in predicting behaviour at lower temperatures.

Fig. 3.3.4 Creep strength model (minimum commitment = MC) for P23 base material based on

the Aloas steels A, B, G, H and P.

N NNN

N

N

VM

VMVM

BBB

B BB BB

BB

G GG

G GGG G

G GG

G G GG

GG

PP

PP

PP

P P PP

PPP

PPP

BAA

AAAAAAAAAA

AAA

AAAAAA

JJJJJJ

JJ

JJ

JJ

JJ

J

JJ

JJ J

J

10

100

1000

17.5 18.5 19.5 20.5 21.5 22.5

PLM/1000 (C=20)

Stre

ss (M

Pa)

T/P23 (EN-10216)MC

ALOAS T/P23 MATERIALSA, B, G, H, J, N, P

-20%

Fig. 3.3.5. The Aloas P23 base material creep strength model (minimum commitment) presented

as a Larson-Miller (PLM) plot. The red line presents the 600°C prediction line (MC).

The test programme was more limited for T/P24 where two heats were tested and assessed (see Fig. 3.3.6). The tests programme showed no significant deviation from the VdTÜV 0 or V&M data 0. Both pipes were upper bound as expected for thin walled product forms. The V&M data

56

has been fitted for comparison purposes. Note that the data fit is extracted from plots and might not represent the true data accurately. This model is for comparison only. The best model for this purpose is the SM1 (Soviet model 1, as described in PD6605 0 ). The model is defined in Eq (5):

log(tr)=β0+β1⋅log[Τ]+β2⋅log(σ0)+β3/T+β4⋅σ0/T (5)

The fitting parameters are:

β0 β1 β2 β3 β4 672.9808 -205.831 -0.5432 -53231 -13.6873

10

100

1000

23.5 24.5 25.5 26.5 27.5 28.5

PLM/1000 (C=26.7)

Stre

ss (M

Pa)

Material EMaterial CV&M 550,600°CSM1 model - 575°CVdTÜV 533VdTÜV 80%

Fig. 3.3.6. T/P24 Creep strength comparison of material C (diamonds) and E (boxes) data against

the VdTÜV 533 (continuous black line) and V&M data (open circles). The dash-dotted line is the SM model line for 575°C and the dotted line represents the VdTÜV -20% in stress.

Creep rupture; Cross Weld The material H cross welds (marked as X) in comparison to the H base material (marked as H) are presented as a Larson-Miller (PLM) plot in figure Fig. 3.3.7. The Aloas P23 general base material curve (red continuous line) is the 600°C prediction line and its -20% in stress is the dash-dotted line below it. Note the low ductility outliers (dotted line). The red arrows indicate running tests. The material P cross welds (marked as X) in comparison to the P base material (marked as P) are presented as a Larson-Miller (PLM) plot in figure Fig. 3.3.8. The AloAS P23 general base material curve (red continuous line) is the 600°C prediction line and its -20% in stress is the dash-dotted line below it. The red arrows indicate running tests.

57

XX

XX

X

X

X

X

X

X X

H

H

H

HHH

H HH H

H

10

100

1000

17.5 18.5 19.5 20.5 21.5 22.5

PLM/1000 (C=20)

Stre

ss (M

Pa)

T/P23(EN-10216)

ALOAS

Fractures with low red. area < 5%

Fig. 3.3.7. Welded H

XX XX

X XX

X

PP

PP

PP

P P PP

PPP

PPP

10

100

1000

17.5 18.5 19.5 20.5 21.5 22.5

PLM/1000 (C=20)

Stre

ss (M

Pa)

MC-20%

Fig. 3.3.8. Welded P

The material G cross welds (marked as X) in comparison to the G base material (marked as P) are presented as a Larson-Miller (PLM) plot in figure Fig. 3.3. The Aloas P23 general base material curve (red continuous line) is the 600°C prediction line and its -20% in stress is the dash-dotted line below it. The red arrows indicate running tests. Note that the lower bound test results are from high temperature tests (660 and 620°C). Creep rupture; Weld Metal The weld metal creep results are presented in Fig. 3.3.9 for B323B, B323 and on the more ductile B322 in comparison to the main base material line. The B323B tested from weld pads

58

indicate lowered strength in comparison to the first test batch (B323). The more ductile B322 also shows adequate creep strength.

G GG

GGG

G GG G

G

G G GG

GGG

G

G

GG

GX

X

XX

X

X

10

100

1000

17.5 18.5 19.5 20.5 21.5 22.5PLM/1000 (C=20)

Stre

ss (M

Pa) MC (600°C)

-20%

660°C

620°C

Fig. 3.3.9. Welded G

10

100

1000

17.5 18.5 19.5 20.5 21.5 22.5

PLM/1000 (C=20)

Stre

ss (M

Pa)

B322

B323B

B323

T/P23 (EN-10216)

MC (600°C)

Fig. 3.3.9. Weld metal creep results; B323 (circles), B323 (triangles) and B322 (squares) in

comparison to MB (continuous lines)

The sighting test results from the weld metal development work package (WP2) in comparison to the main base material line at 600°C ±20% in stress are presented in Fig. 3.3.10. Note that the selected project consumable B323B is stronger than both G and H base material.

59

10

100

1000

19 19.5 20 20.5 21 21.5 22

PLM/1000 (C=20)

Stre

ss M

PaB323 CVNbTiB

B323B

B323 trans

B322 base P23

B322 trans

B575B TiB

B576 CVNb

B702 (B322 Ni)

B703 (B323Ni)

Sumitomo trans

Sumitomo

Thermanit P24

23L

subarc

Thyssen P23TIG

Therm P24 trans

P24 C87

HITEPR

HITEPR trans

Fig. 3.3.10. Short term creep results from weld metal development (WP2)

3.3.2.2b Creep strain modelling The development work on the “logistic creep strain prediction” model (LCSP 0) was partly done within AloAS for the special needs of translating and extrapolating creep strain parameters 00 for the needs of WP5 finite element work (local model of girth weld). AloAS creep rupture and creep strain modelling presentations will be presented at the Creep 2008 conference in Bayreuth, Germany. Also the girth weld FE-implementation of the strain model is presented there. The aim to reliably assess components under realistic loading conditions (creep conditions) has been accomplished. The main model used for the Aloas creep strain data is the logistic creep strain prediction method. It is a uniaxial model for predicting full creep curves using the time to creep rupture as base (see section 3.3.2.2a). It predicts primary, secondary and tertiary creep as a function of time to rupture temperature and stress. The LCSP model is presented in Eq. 1 and the LCSP creep shape parameters for P23 steel are presented in Table 3.3.5. The data used for the modelling consists of 839 creep strain value data points in the temperature regime 550-620°C. The bulk of the data was from the primary creep region for better presentation of small strains (important in long term prediction) as shown in Fig. 3.3.11. The heats represented in the data fit were B, G, H and J. The LCSP defining the whole creep curve is defined as;

( ) 0

/1

1)log()log(

log xCtCt

pr

t ⋅⎟⎟⎠

⎞⎜⎜⎝

⎛−

++

=ε

ε

(6)

where tε is the predicted strain at time εt , rt is the rupture time that can be directly taken from the rupture master equation (see creep rupture assessment), i.e the rupture time ( rt ) is the fix point (last point) from which the rest of the curve can be derived as seen in Fig. 3.3.12.

60

5 10 15 20 250

100

200

300

400

500

Num

ber d

ata

poin

ts

Strain %

Fig. 3.3.11. Distribution of data in the creep strain assessment according to accumulated strain.

x0(σ,T) p(σ,T) C

4.21+0.69601⋅log(σ)-7860/(T+273) 14.3-2.271⋅log(σ)+4829/(T+273) 3.5

Table 3.3.5. Creep strain shape equations for Aloas P23. Note: Stress in MPa , temperature in °C.

Fig. 3.3.12. Creep strain simulation with LCSP (red curve) and corresponding Norton minimum strain rate prediction (linear blue line) for P23 at 575°C / 91 MPa (rupture predicted at 100.000 h ).

The model data fit is validated by actual versus predicted time to 1 and 2% strain presentations as in Fig. 3.3.13. The model accurately describes the primary creep for materials G and H. The time to strain fit is also rather good for material B where the model is conservative by a

factor of two in time. For material J the conservativeness is even larger, i.e. a factor of four in time.

Fig. 3.3.13. Actual creep strain data, time to specified strain (here 1 and 2%) compared to predicted time to strain. Note that the curve form of material J is not very well represented by the model.

0 2 .104 4 .104 6 .104 8 .104 1 .1050

0.05

Time (h)

Stra

in (m

m/m

m)

epri σ T,( )

esec σ T,( )

etert σ T,( )

tr σ T,( )

α

61

100

1000

10000

100 1000 10000

True time to strain (h)

Pred

icte

d tim

e to

str

ain

(h)

BGHJ

3.3.2.2c Girth weld simulation using the creep rupture and creep strain models The LCSP has been used in the AloAS project for modelling base material, weld metal and to simulate the creep response of heat affected zones (HAZ). The material shape constants for the weld material B323 were calculated based on the high temperature strain data from EON (Fig. 3.3.14). An example of a girth weld simulation using the AloAS creep strain models for base material, weld material and heat affected zone is presented in Table 3.3.6 and Fig. 3.3.15.

1.0 1.5 2.0 2.5 3.0

1.0

1.5

2.0

2.5

3.0

Pre

dict

ed ti

me

to s

train

, log

(t)

Measured time to strain, log (t)

Fig. 3.3.14. Measured vs predicted time to specified strain for B323B WM at 660°C / 80 MPa.

Material zone in weldment Predicted uniaxial time to rupture tr (BM), h 97 000

tr (HAZ), h 70 000 tr (WM, B323) 280 0001)

1) B323B experimental consumable, life extrapolation with time factor 2.9, shape functions from high temperature short term tests.

Table 3.3.6. Calculated rupture times for single zones in weldment: Girth weld in a pipe with 300 mm inner radius and 28 mm wall thickness with internal pressure of 139 bars (nominal stress = 70 MPa) at 595°C. 3-zone case; base material (BM), heat affected zone (HAZ) and weld metal

(WM). 62

For the heat affected zone (HAZ) it is assumed that the time to rupture is following a 20% reduction in strength (WSF=0.8) and the shape of the creep curve takes this into account as an increased stress (σ/WSF replacing σ in shape equations).

5 .104 1 .105 1.5 .105 2 .1050

0.02

0.04

P23HAZB323

Time (h)

Stra

in (m

m/m

m)

0.2 0.4 0.6 0.8 10

0.02

0.04

P23HAZB323

Normalised time (h)

Stra

in (m

m/m

m)

Fig. 3.3.15. Simulated creep curves in time and normalised time (t/tr) for HAZ (weakest), BM and WM (strongest) used in P23 FEA. As a result of defining the creep strain response of all weld-zones it was possible to simulate large component behaviour (the welded pipe). As commonly known the maximum equivalent strains develop on the inner surface of a pipe whereas the maximum equivalent stresses will redistribute towards the outer surface due to creep. The multiaxial creep ductility again (due to triaxiality constraint) will be reduced towards the outer surface. The Aloas girth weld simulation was modelled using the state of the art LCSP model and the FEA results were filtered with the expression for rigid plastic deformation called Λ filtering as defined in Eq (4) 0[000. The Λ-filtering brings out the impact of constraint and differing material creep properties in the zones of the weld. The Λ-filter is defined as:

mm

fm

fu

hεε

εε

⋅⋅−

=⋅=Λ)5.1exp(65.1

1

(7)

where h is

⎥⎦

⎤⎢⎣

⎡⋅

=q

h kk

3σ (8)

and where kkσ is the trace of the stress tensor ijσ and q the von Mises stress.

The LCSP model in its principal form is uniaxial like many material models, and as such it needs to be generalized into a multiaxial form in order to be able to tackle two- or three-dimensional problems. In addition, issues such as time-integration and coupling with the rate-independent part of the solution need to be solved. The FE-implementation of the LCSP model is built around Eq.(6), i.e. the formulation providing the state of equivalent true-strain as a function of time. Differentiating with respect to time to attain the creep strain-rate leads to a form:

( ) ( ) ( )( )0 1 2 1 3 2, , , ,c

rt t x p C Tε α α α β σ α β= ⋅ + ⋅ +& , (9)

where α is a function of the various parameters and functions of the LCSP model, αi are constants and βi are functions of stress and temperature, respectively. The multiaxiality generalization is performed using an associative flow rule, following the implementation of

63

incremental von Mises type plasticity laws, i.e. the plastic deviatoric strain and equivalent plastic strain are connected via the following equations;

pl plde d nε= , (10)

where

32

Snσ

= ⋅ (11)

and where S is the deviatoric stress tensor. Substituting Eqs. (10) and (11) to Eq. (9) leads to the final multiaxial form of the LCSP model

( )1 3 3 432

cij ijT sε α δα δα β α⎛ ⎞= ⋅ + + ⋅⎜ ⎟

⎝ ⎠&

(12)

where δ, αi and βi are new constants and functions already defined as in Eq. (9). Eq. (12) is implemented using a mixed-formulation for the creep strain rate and displacement finite element method, utilizing p order finite elements (the interpolation orders ranging from 2 to 4). Time-integration was carried out using a variable-order variable-stepsize back differentiation implicit time-stepping routine. The implementation was directly coupled to solution of a simultaneous heat transfer problem to enable a varying temperature field within the solution domain.

In the P23 case study, 2D axisymmetric models of a pipe with a girth weld were made for the material property cases defined in Fig. 3.3.15. In the simulation the pipe is subjected to internal pressure loading and the boundary conditions are specified to present a situation of a long pipe (i.e. continuity conditions are prescribed at the far edge of the models and symmetry at weld centre).

Fig. 3.3.16. FE-mesh for Comsol girth weld simulation.

The girth weld is modelled as a tree material region, the parent material, the heat affected zone and weld metal. The elements in the Comsol model uses p-elements with interpolation degrees ranging from 2 to 4 in a mixed-element formulation (hence the differences in the construction of the meshes). The results of a simulation was studied at the point of the base material strength (consumed life = 100%) and it was shown that the HAZ will cause creep “exhaustion” in a faster rate than the base or weld metal (see Fig. 3.3.17). It is to be noted that the Λ-filter is still under development together with the Comsol strain model implementation. The Comsol application has shown signs of being too sensitive to localisation of stresses and strains. However the LCSP model has also now been successfully transferred to ABAQUS software where these drawbacks seem to have been overcome.

The P23 weld simulations have been performed with data determined from actual data. Earlier simulations on P22 show similar weld behaviour i.e. the highest creep strains are found on the inner surface of the pipe near the weld root within the heat-affected zone. The overmatched B323B used in the simulations causes severe localisation and damage within the HAZ.

64

Material Λ calculated with Mises strain Λ calculated with axial strain

P23

Λmax = 0.10%, indication of

shearing, localisation in HAZ (over matching WM)

Λmax = 0.08%, localisation in HAZ(ID),

εmax= 24% in HAZ

Fig. 3.3.17. axisymmetric Λ-filtered creep exhaustion results for the P23 simulation at 100% of BM life

3.3.2.2d Low cycle fatigue results A plot of LCF endurance (Cycles to initiation Ni) against total strain range (%), Fig. 3.3.18 and Table 3.3.7, indicates that all three parent materials (B, H and KA1358 (CMV)) behave similarly, whereas B323B weld metal is substantially inferior. (Note that three data points, plotted with smaller markers, represent tests with control problems and one pre-existing weld defect: when these are excluded, the variability in the weld data is not abnormally high.) Whereas B323B showed better creep properties in longitudinal than in transverse or cross-weld tests, Section 2, its LCF performance is poorest in longitudinal orientation. Notably, the weaker B322 weld metal is greatly superior to B323B, producing endurance data in line with parent P23. When results are plotted against plastic strain range (see Fig. 3.3.19 and Table 3.3.7), however, the weakest parent material KA1358 (CMV) shows the greatest endurance, while the strongest, B, shows the least. Again, B323B weld metal is inferior, while on this graph, B323B cross-weld specimens show intermediate performance. Fig. 3.3.19 also shows that LCF endurance is significantly dependent on temperature, being approximately a factor of two greater at 450°C than at 600°C, while the data at 575°C are intermediate. Limited fractography and metallography on material B showed no unusual features. The introduction of dwells produced moderate reductions, of the order of 30-50%, in the endurance lives of both the P23 steels and weld metals, Table 3.3.8 and Fig. 3.3.18. (The reduction was still evident when analysed in terms of plastic strain range, Fig. 3.3.19, showing that it was not simply due to increased plastic strain associated with modification of the hysteresis loop.) Hence, damage evidently took place during the relaxation period when elastic tensile strain is converted to tensile creep strain. The softer material H showed more stress relaxation than B during the dwell, but its endurance remained comparable. The softer weld metal B322 showed a slightly greater endurance reduction in the dwell test, but retained a markedly better performance than B323B.

65

0.10

1.00

10 100 1000 10000 100000

Cycles to Initiation

Tota

l Str

ain

Ran

ge (%

)

B CMV HH 600°C H 450C B 5minB 60min H 5min H 60minB323B longit B323B trans B323B XweldB322 Power (B) Power (CMV)Power (H) Power (B323B longit) Power (B323B Xweld)Power (B323B trans)

Fig. 3.3.18 LCF data – Showing cycles to crack initiation versus total strain range.

0.0001

0.001

0.01

10 100 1000 10000 100000Cycles to Initiation

Pla

stic

Str

ain

Ran

ge (m

m/m

m)

B CMVH B322B 5min B 60minH 5min H 60minH 600C H 450CB323B long B323B transB323B Xweld Power (H)Power (B) Power (CMV)Power (H 600C) Power (H 450C)Power (B323B trans) Power (B323B Xweld)Power (B323B long)

Fig. 3.3.19. LCF data – Showing cycles to crack initiation versus plastic strain range.

Cyclic stress strain relationships are shown in Fig. 3.3.20. The soft service exposed CrMoV material KA1358 showed considerable lower cyclic strength than the new P23 materials, while the difference in ambient temperature tensile strength between the two P23 steels was matched by a difference in high temperature cyclic strength. Material H showed substantially greater cyclic strength at 450°C than at 600°C, as would be expected, with the 575°C data points just above those obtained at 600°C. B323B all-weld-metal was significantly stronger than both parent materials, reflecting its high alloy content, but weld B322 was also somewhat stronger than the parent steels, demonstrating the intrinsically higher strength of weld metal.

66

100

1000

0.0001 0.001 0.01

Plastic Strain Range (mm/mm)

Stre

ss R

ange

(MPa

)

H 450CB323B longitudinalB323B transverseB322 longitBHB323B Xweld (mat G)H 600CCMV

Fig. 3.3.20. LCF data – Cyclic stress range as a function of plastic strain range

The two cross-weld tests showed different failure modes, Table 3.3.7. One failed in the weld metal, the other in parent material. The latter also contained a part-through crack located parallel to the fusion boundary in the centre of the reaustenitised HAZ. Data on number of cycles to failure (Nf) are also of interest, Table 3.3.7. It may be noted that the crack growth rate, as indicated by the (Ni – Nf) value, is often greater for low applied strain tests than for the higher strain tests with lower Ni values. In investigating this behaviour, loop data were reassessed to confirm whether the load drop criterion correctly identifies Ni. This showed that toward the end of each test, the loop shape begins to become asymmetric, with a higher peak compressive stress. This may be rationalised on the basis that once a crack forms, the specimen compliance in tension is reduced as the crack opens, but is unaltered in compression when the two halves of the crack are forced back together. The onset of loop asymmetry is therefore a good measure of Ni. Limited work indicated that this criterion would have produced slightly lower Ni values, but would not have affected the above findings in respect of crack growth rate. The data show that a weaker parent material incurs the greater plastic strain, but attains the lesser peak stress, at a given total strain range. The net effect is that there is little difference between the endurance of the different parent materials.

3.3.2.2e Fatigue crack growth results Fatigue crack growth (FCG) tests have been performed with 0.8 compact tension (CT) test specimens with width, W = 40mm and thickness, B = 20mm, Table 3.3.4. The specimens were manufactured such that the crack plane contained the radial and circumferential directions in pipe B, and the radial and axial directions in KA1358.

67

Material Specimen Number

Total Strain Range (%)

Plastic Strain

Range (%)

Stress Range (MPa)

Cycles to initiation

Cycles to failure

2256 1.0 0.674 530 675 1115 2259 0.7 0.382 532 930 2239 2257 0.5 0.189 503 2335 3718 2260 0.4 0.116 484 7887 11615

B (565°C)

2258 0.3 0.032 442 59672 65881 2427 1.0 0.717 511 661 1387 2423 0.8 0.533 454 878 1871 2424 0.5 0.200 457 4313 7607 2425 0.4 0.162 405 7467 9216

H (575°C)

2426 0.31 0.069 409 40751 43740 2430 0.99 0.749 435 469 686 H (600°C) 2431 0.38 0.158 385 4920 7786 2432 0.99 0.574 738 962 1488 H (450°C) 2433 0.39 0.1925 658 8901 14117 2277 1.0 0.778 332 656 - 2278 0.5 0.324 316 2997 -

CrMoV KA1358 (565°C) 2279 0.4 0.227 279 8092 -

2463 0.99 0.611 616 361 529 2470 (weld

defect) 1.00 0.637 570 100 259 2517 1.00 0.616 626 140 346

2464 (test imperfect) 0.70 0.288 671 160 671

2465 0.50 0.127 603 876 2063 2466 (test imperfect) 0.40 0.076 603 835 2263

2518 0.40 0.055 562 3175 8346

2520 0.40 0.060 544 3729 4912

(interrupted) 2467 0.30 0.016 481 10398 11128

B323B longitudinal

(575°C)

2519 0.30 0.020 491 22018 37131 2471 0.99 0.49 620 67 69 2475 1.00 0.50 607 134 281 2472 0.40 0.20 506 3468 3473

B323B transverse

(575°C) 2473 0.30 0.15 450 25070 25957

B322 longit (565°C) 2403 1.00 0.656 587 722 919

2453 Base of

weld 0.399 0.155 419

3224 Failed in

parent

5186 Also crack

across central part of HAZ

B323B cross weld (mat. G) (575°C) 2451

Top of weld 0.296 0.054 371

27196 Failed in

WM 28530

Table 3.3.7. LCF continuous cycle test conditions and results

68

Material Spec.

Number Dwell Time min

Total Strain Range

(%)

Plastic Strain Range

(%)

Stress Range (MPa)

Peak Stress (MPa)

Stress at end

of dwell

Cycles to initiation

Cycles to

failure

2261 60 1.00 0.759 534 259 133 326 492 B (565°C) 2262 5 1.00 0.754 525 256 177 521 966

2428 60 1.00 0.77 494 235 96 404 493 H (575°C) 2429 5 1.00 0.765 464.5 227 135 412 503

2469 60 1.00 0.712 635.5 303 168 70 73 B323B long.

(575°C) 2468 5 1.00 0.696 600 284 180 157 217 B323B trans.

2474 (575°C)

60 1.00 0.733 594 290 155 64 96

B322 long.

2404 (565°C)

60 1.00 0.712 619 297 165 290 326

Table 3.3.8. LCF dwell test conditions and results

Specimens were manufactured with a spark eroded slot to an a/W value of 0.4 (where a is crack depth) and then fatigue pre-cracked in air at a frequency of 50 Hz and load ranges of between 9 and 10kN. The fatigue pre-crack depth was around 2.5mm and took more than 105 cycles in each case. FCG tests were carried out on each specimen at a temperature of 565°C and frequency 1Hz, at a constant load range, with crack monitoring via crack mouth clip gauge. Cracks were grown to an overall depth of approximately 26mm. After testing, specimens were sectioned at 0.25B for metallography, and the remainder broken open in liquid nitrogen for measurement of crack depth. Crack growth was somewhat slower in P23 material B than in CrMoV steel KA1358, Fig. 3.3.21.

Material

Specimen Number

Notch Depth (mm)

Pre-Crack Depth (mm)

Final Crack Depth(mm)

Load Range (KN)

Starting ∆K

(MPa√m)

Final ∆K

(MPa√m)

Number of

Cycles at

565°C F1873 15.85 18.48 25.99 13.67 27.4 57.0 10864 B F1874 15.91 18.12 25.79 9.11 19.0 38.2 20959

KA1358 F1876 15.43 18.70 25.71 9.00 20.0 34.4 12318

Table 3.3.9. Fatigue crack growth data

1.0E-05

1.0E-04

1.0E-03

1.0E-02

10 100Δ K (MPa √ m)

da/d

N (m

m/c

ycle

)

B, P23 KA1358, CrMoV

Fig. 3.3.21. Fatigue crack growth data

69

3.3.2.2f Creep crack growth results The CCG results are presented in Table 3.3.10.

Spec K MPa√m

σref Time

h Time to da=0.5

mm max LLD Notch loc. Crack path

HAZ1 13 89.6 1159 666 h 0.7029 FL WM

HAZ3 15 103 287.2 172 h 0.5682 FL FL/CGHAZ-WM

HAZ4 15 103 2606 2189 h 0.9233 CGHAZ WM

WM2 15 103 91.4 57.5 h Wide open WM WM

Table 3.3.10. CCG test details for HAZ and WM CT specimens.

After the testing a 1.5mm thick slice was removed from the middle of the specimens by wire erosion for metallographic investigation leaving the rest of the specimen intact for final crack length measurement from the fracture surface. The metallographic observations have been summarised in Table 3.3.11. Specimen Remarks

HAZ1 Crack path curved (ending almost perpendicular), following a HAZ in WM, one side crack turning back towards FL/CGHAZ. Some cavities.

HAZ3 Fairly straight crack close to FL/CGHAZ, moving away from FL to WM, not many side cracks. Some cavities. The tip of the side crack is oxidised.

HAZ4 Stress Relief cracks in WM far away from the crack. Crack first followed a HAZ in WM, one side crack, which became the main crack, grown to FL. The crack has possibly followed stress relief cracks. Slag inclusions all over.

WM2 Some very small side cracks near the notch tip, some stress relief cracks mid way. The specimen broken into two pieces. Deformed grains near the surface where fast fracture.

Table 3.3 11. Summary of the metallographic investigation of the ALOAS P23 CCG tests

The results of the tests can be summarized as follows:

• In all cases the crack grew in the WM, indicating that this was the most creep brittle zone • The fracture was intergranular in all cases • In two HAZ specimens the crack followed the HAZ in WM, following the shape of the

weld bead and curving strongly, becoming almost perpendicular to FL, see Fig. 3.3.22. • Stress relief cracking was observed in areas far away from the crack plane • The calculated Vc/V-rates varied between 0.01 (at a0) and 0.3 (at af) indicating that the

fracture was creep brittle and the C* parameter is not valid • Crack growth does not seem to be correlated by K either • The peculiar shape of the PD and LLD curve of the specimen HAZ1 was reflected as a

strange shape in the C* curve, see Fig. 3.3.23. • All four C* curves, however, fall nicely on a single line with relatively narrow scatter

band

70

Fig. 3.3.22. Micrographs showing the crack path following the weld bead shape in specimens

HAZ1 and HAZ4. ALOAS CCG da/dt vs. C*, ALL

1.E-04

1.E-03

1.E-02

1.E-01

1.E+00

1.E+01

1.E-03 1.E-02 1.E-01 1.E+00 1.E+01 1.E+02C* [N/mmh]

da/d

t [m

m/h

]

HAZ1HAZ3HAZ4WM2

Fig. 3.3.23. The C* plot of all four CCG tests.

Discussion and conclusions Creep (BM) Materials B (tempered at 760°C, UTS 675MPa), G and H (tempered at 780°C, UTS 635MPa -G and 620MPa -H) were tested. Material H showed poorest creep properties, with B intermediate and G generally best. Hence, whilst high temperature 780°C tempering sharply reduces UTS, its effect on creep performance appears to be lesser. A review of ECCC P23 data showed that the few (650°C) data points at below 100MPa showed a downturn in rupture life, compared with the trend at higher stresses. Further, 18 casts had been tested at 98MPa, but of the 6 casts tested below 80MPa, all but one had been chosen from the 6 that were strongest at 98MPa. The effect of this biased selection was to conceal and minimise the downturn in the rupture data below 100MPa. Figure 3.3.25 illustrates the ALoAS results at 650°C, compared with trend lines based on the ECCC data at higher stresses and (after adjustment to remove bias) at lower stresses. The results confirm a marked downturn at realistic plant stresses. The ALoAS results fall below the ECCC data, probably because they relate to thicker section materials.

71

Life

Stre

ss M

Pa

ECCC >100MPa trendECCC < 100MPa trendBGH

90

50

70

Fig. 3.3.25. Logarithmic stress – life plot of 650°C creep rupture data.

The adverse trend is, however, at least partly an artefact due to oxidation during testing. This acts to increase effective stress and reduce creep life, especially for low stress tests, which (at all test temperatures) suffer greatest metal thickness loss. To investigate this, VTT carried out comparison tests in argon at 650°C 78MPa, finding that an air atmosphere reduced creep life by 38%. E.ON found reasonable agreement with a simple oxidation model, suggesting that it might be possible to apply correction factors to air test data. However, Komai 0 obtained similar results, and suggested that decarburisation beneath the oxide also plays a role. If this is correct, a modelling correction is unlikely to be adequate, and only inert atmosphere testing will be capable of producing reliable design stresses for P23. It is understood that ASME have recently accepted a similar view, and have now withdrawn design values for P23 steel within the lower stress ranges, until reliable data can be obtained e.g. by inert atmosphere testing. For the plant manufacturer and end user, the lack of reliable design data is a major barrier to the application of P23 steel.

0.01

0.1

1

10

10 100 1000 10000 100000

Cycles to initiation

Hys

tere

sis

ener

gy -

Stre

ss ra

nge

(MPa

) x p

last

ic

stra

in ra

nge

(mm

/mm

), m

id-li

fe lo

op

B CMVH B322H 600C H 450CB323B longit B323B transB323B Xweld Power (all parent 565/575C)

Power (B323B longit) Power (B323B trans)Power (B323B Xweld) Power (H 450C)Power (H 600C)

Fig. 3.3.24. LCF data – Showing strong correlation for all parent material casts between nominal

“hysteresis energy” and cycles to initiation.

72

0.00

0.20

0.40

0.60

0.80

1.00

1.20

100 1000 10000 100000

Nominal total hysteresis energy - mid-life loop stress range (MPa) x plastic strain range x no. cycles to initiation

Tota

l ela

stic

str

ain

rang

eBHCMVB322H 600CH 450CB323 longB323 transB323 XweldPower (CMV)Power (H)Power (B)

Fig. 3.3.25. LCF data – Nominal total hysteresis energy as a function of elastic strain range.

A simplified version of this hypothesis, based on nominal total hysteresis energy, is examined in Fig. 3.3.25. For total strain range values of 0.5% and above, nominal total hysteresis energy to crack initiation is indeed fairly constant. However, as the strain range falls below 0.5% and the number of cycles to initiation rises sharply, nominal total hysteresis energy required for crack initiation also tends to increase, albeit much less sharply. This implies that the hypothesis is largely correct, but cycles with very low energy input do become rather less effective in causing damage. Contrary to common assumption, LCF performance does vary significantly with temperature. Improved endurance at the lower temperature of 450°C, Fig. 3.3.19 and Fig. 3.3.24, may be linked with higher cyclic strength, Fig. 3.3.20. Comparing the 575°C and 600°C data, the cyclic strength change is marginal, but the reduction in endurance is rather more significant, perhaps because the 600°C cycle also introduces some creep relaxation damage. The dwell test data show a significant creep-fatigue interaction, but do not suggest a major concern. Comparing the parent materials, the weaker H apparently compensates for its higher relaxation strain with higher creep ductility, thereby broadly matching the creep-fatigue endurance of B. The weld data are perhaps the most important of the LCF results. The poor performance of B323B provides a further reason to be wary of using such a strong, brittle weld metal. However, the notably better performance of alternative P23 weld metal variant B322 (which was not particularly soft, and did not especially perform well in the creep tests) indicates that a suitable weld metal can readily be chosen to avoid any major concern with LCF performance and bore cracking resistance. The cross-weld results indicate that the HAZ is not a particular concern in LCF. When the composite specimen is cycled, the stronger weld metal is intrinsically more brittle than the parent material, but because the weaker parent material accommodates most of the imposed strain, the weld accrues little plastic strain. Hence, although the all-weld-metal data are greatly inferior to the parent data, the composite cross-weld specimen does not necessarily fail in the weld metal.

73

In the simple uniaxial geometry, both weld metal and parent failure were observed. The plastic strain and hysteresis energy data reflect the high energy absorption by the parent material part of the cross-weld specimen. The LCF tests on ALoAS pressurised tube butt welds by IfW showed a different failure mode, with cracking in material at the outermost edge of the weld HAZ. This can be attributed to the effects of pressure cycling in radially expanding the weak parent tube, while the strong weld and HAZ resist radial expansion, hence producing mismatch cracking at the edge of the HAZ. The uniaxial test does not, therefore, fully simulate this failure mechanism in a real component. The apparently anomalous finding that fatigue crack growth rate can be greater in a specimen subjected to a lower imposed strain cycle is of interest. The LCF test continuously creates fatigue damage throughout the specimen. When a crack initiates, therefore, it then grows through highly pre-damaged material. By contrast, a conventional fatigue crack growth test measures the rate at which a crack grows into largely unexposed material. It is not surprising, therefore, that the two tests do not result in identical crack growth behaviour. It may be surmised that the low imposed strain cycle, which inputs more energy into the specimen before a crack initiates (Fig. 3.3.25), may produce a high and uniform level of fatigue damage in the specimen prior to crack initiation. Hence, even though the imposed strain level remains low, the crack once initiated grows rapidly to cause failure. By contrast, it can reasonably be surmised that the higher imposed strain cycle produces less general damage throughout the specimen before a crack first initiates. Hence, even though the imposed strain level remains high, the crack once initiated may not grow so rapidly to failure. If this is correct, the implication is that standard FCG test data could be non-conservative in certain applications. FCG data should be reasonably applicable when a fatigue crack grows from a stress-raiser into material which has not sustained prior fatigue damage, as for example in the typical bore cracking case or thermal fatigue situation. However, the use of standard data may be less valid when the fatigue cycle has damaged the full pipe wall thickness, for example where the fatigue loading is of mechanical origin. CCG tests The CCG tests showed that the creep crack preferred to grow in the WM, indicating that this was the most creep brittle zone. The fracture was intergranular in nature for all cases and followed the heat affected zone between beads within the WM. In some cases it even curved away from the direction of the fusion line (perpendicular to maximum stress). In one of the cases stress relief cracking was observed in areas far away from the crack plane. The CCG results gave calculated Vc/V-rates between 0.01 (at a0) and 0.3 (at af) indicating that the fracture was indeed creep brittle, and that the C* parameter is not a valid parameter for describing the fracture mechanics. However, the crack growth also does not seem to be well correlated by K. In spite of the above, looking at the four C* curves it is evident that the results fall on a single line with a relatively narrow scatter band. Conclusions of Work Package 3 A comprehensive high temperature materials database has been generated on P23 steels and weldments. These data provide support for high temperature power plant design with P23 steel. Background data on factors such as weld performance and defect development under plant cycling conditions will also support the lifetime management of operational power plant. The results show how important it is to have a good creep ductility for the weld metal of the newer high strength low alloy steels like P23 and P24, in addition to promising creep strength in preliminary qualification. The reasons are related to mechanisms that lead to strain localisation

74

and cracking towards longer term creep exposure, and emphasise the importance of ductility criteria in short term testing and initial selection of candidate compositions. The results have also established an excellent basis for further development of suitable compositions that would perform better, and fundamental understanding on the mechanisms of the creep behaviour of multibead weld metal for high strength low alloy steels like P23 and P24. Further development of the weld metal is clearly needed, and there are new avenues to explore for the purpose. For this class of steels, the project has provided new tools and understanding for faster evaluation and further improvement of high performance welds.

75

3.4 WP4: Microstructural Assessment This work package (WP4) provides the means and tools for evaluation and modelling of the base material (BM) and welded joint microstructural evaluation. It is divided into two tasks: Task 4.1: Microstructural modelling and assessment Task 4.2: Characterization of aged material The output is towards WP5 High temperature design and assessment WP4 Objectives: • Evaluation of the microstructural evolution of the base material and welded joint • modelling of microstructural evolution The main objective of WP4 is the evaluation of the microstructural evolution of the base material and welded joint for usability and life management planning for high temperature applications. The intended target for the characterisation is power plant components such as steam lines, water walls, etc. The knowledge acquired from aging, material property testing and the microstructural investigation is then to be utilised for modelling. The material has been aged by CSM and VTT. The material codes and more information about the aged samples are reported in WP1. Introduction The mechanical properties of advanced low alloyed steels like T/P 23 and T/P 24 are highly dependent on the microstructure. Their about 1.8 times higher creep strength compared to conventional T/P 22 steel is achieved by a martensitic-bainitic structure containing iron-chromium carbides at grain boundaries and lath boundaries and very small precipitation strengthening carbonitrides inside the laths. This advantageous microstructure is not, however, thermodynamically stable, and during high temperature service the microstructure gradually evolves toward a more stable structure with lower creep strength. The first changes in carbide structure occur as gradual coarsening of M7C3 and M23C6 carbides located at grain boundaries and lath boundaries. The kinetics of carbide growth is governed by the Ostwald ripening (coarsening) law d3 – do3 = kt, where d is the particle diameter after exposed time t, do is the initial particle diameter and k is a coarsening constant having an Arrhenius equation type of temperature dependency. Finally M7C3 and M23C6 carbides gradually change to the most stable carbide type M6C, which grow at a considerably slower rate. M6C carbides are molybdenum and tungsten rich, and development of these carbides depletes solution strengthening molybdenum and tungsten from solution, which is detrimental to creep strength. Moreover, precipitation of M6C carbides may cause solution of surrounding precipitation strengthening fine MX carbonitrides, causing additional deterioration of creep strength. Due to this strong effect on creep behaviour, microstructural stability is one of the most important properties of modern precipitation strengthened boiler steels. The aim of this study is to compare the microstructures and stability of T/P 23 and T/P 24 steels by microscopy of the steels after long term high temperature annealing. Task 3.4.1: Microstructural modelling and assessment 3.4.1.1. Microstructural assessment of materials, weld metals and heat affected zone structures in order to study the evolution of employed and/or tested material to predict the microstructural changes that will occur in power plant steels, TEM analysis of precipitates: type, chemical analysis, dimension, density The summary of the aging tests carried out at CSM and VTT is shown in table 3.4.1. Crept

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specimens have been also used.

Heat Treatment Grade and

ALoAS Code OD x WT

[mm x mm] Quenching [°C] Tempering [°C]

Temperature & Ageing Time [°C - hours]

P23 B 219x31.75 1070 760 550 -3,000 T23 H 76x12.5 1050 780 600 - 3,000 P23 B 219x31.75 1070 760 550 - 6,000 T23 H 76x12.5 1050 780 600 - 6,000 P23 B 219x31.75 1070 760 550 - 10,000 T23 H 76x12.5 1050 780 600 - 10,000 T24 C 44.5x6.3 1000 750 600 - 3000 T24 C 44.5x6.3 1000 750 600 - 6,000 T24 C 44.5x6.3 1000 750 600 - 10,000

P23 G 219x31.75 1070 780 625, 675 - 100, 300, 1000, 3000, 10000

T24 C 44.5x6.3 1000 750 625, 675 - 100, 300, 1000, 3000, 10000

Table 3.4.1. Aging conditions for grades 23 and 24. Hardness measurements Measured hardness (HV 10) on aged samples showed, as expected, decreasing values as a function of time and temperature (table 3.4.1). Samples from grade 23 pipes showed lower hardness values than the tubes, presumably due to their slower cooling rate during manufacture, as a consequence of their larger size. The T24 hardness, initially higher than that of T23, dropped down below the T23 values after 10.000 hours of aging at 600°C. This could be an effect of a different precipitate evolution sequence.

Figure 3.4.1: Hardness values as function of time and temperature of aging for grade 23 and 24 All the aged specimens of grade 23 and 24 have been analysed by LM and SEM. Some specimens have also been selected for further more detailed analysis by STEM+EDS, and for X Ray Diffraction (X-RD) analysis of the extracted precipitates.

Microstructural evolution Figure 3.4.2 shows grade 23 microstructures observed by SEM after aging at 550°C for 3-6-10.000 hours, compared with the as received material. It is possible to observe a progressive increase of the sizes of precipitates at grain boundaries. Figure 3.4.3 shows the microstructures after aging at 600°C. The coarsening of the precipitates is more pronounced.

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ness

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0) T23 500°CT23 600°CP23 550°CP23 600°CT24 550°CT24 600°C

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As received after 3000h After 6000h after 10000h

Figure 3.4.2: SEM microstructure of grade 23 after aging at 550°C

As received after 3000h After 6000h after 10000h

Figure 3.4.3: SEM microstructure of grade 23 after aging at 600°C

EDS analysis was performed on the white particles at grain boundaries and inside the grains. The EDS spectra with the composition of the particles at the grain boundaries and inside the grains are shown in Figure 3.4.4. The particles at grain boundaries contain W, while the particles inside the grains contain Nb. This difference in composition indicates that the two types of precipitates belong to different families. Further TEM investigation and analysis is reported later.

Figure 3.4.4 EDS analysis of precipitates

Similar microstructural analysis has been carried out on grade 24 and Figure 3.4.5 shows the effects of aging to 3000h.

as received After 3000h at 550°C After 3000h at 600°C

Figure 3.4.5: Microstructure of Grade 24 as function of time and temperature of aging 3.4.1.2 Microstructural modelling to predict changes as a function of time and temperature and the production of a time temperature accelerated ageing matrix

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The TEM microstructures in thin foils and extraction replicas after the heat treatment are shown for both grades in the WP1 section of this report. Precipitate evolution by TEM analysis TEM investigations were performed on both grades. Specimens aged at different temperatures and times with increasing Larson-Miller Parameter values [(PLM= (T(°K))x(C+log(time(Hours))x1000] were used as shown in table 3.4.2.

Grade Temperature (°C) Aging Time (hours) PLM (C=20) 23 As received 23 600 6425 20,78 23 600 15000 21,11 23 625 3000 21,08 23 625 10000 21,55 23 660 1130 21,51 23 675 10000 22,75 24 As received 24 625 10000 21,55 24 675 10000 22,75

Table 3.4.2: aged specimens analysed by TEM Extraction replicas were prepared and analysed by STEM+EDS, and Selected Area Diffraction Patterns (SAPD) have been used to identify the compositions and types of precipitates. The phases predicted by JMatPro tools at the equilibrium stage in the temperature range 550-700°C are showed in figures 3.4.6 for both grades: Grade 23: only M6C, M(C,N) are predicted; Grade 24: M(C,N); M6C jointly with M23C6 are predicted.

Grade 23 Grade 24

Figure 3.4.6: equilibrium diagram by JMatPro On the basis of EDS analysis and X-ray diffraction, the following precipitates were identified: Grade 23:

1. as heat-treated material: carbides M2C, M6C, M23C6 and MC are present after the heat treatment. M2C are practically W-V carbides, while M6C precipitates have a very high chromium content compared with the usual value. The size distribution (see Figure 3.4.7) shows that the sizes of M2C and MC precipitates are below 100 nm, with M6C in the range 200-300 nm, and M23C6 spread within the range 100-400 nm.

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Figure 3.4.7: Frequency of precipitates Figure 3.4.8: Frequency of precipitates vs. Vs. size in Grade 23 after heat treatment size in Grade 23 after 6425 hours at 600°C

2. After 6425 hours at 600°C, the following modifications in size distribution of the

precipitates occur (fig. 3.4.8): • M2C do not change size • MC change size and their frequency is reduced • The M6C already present in the as heat-treated material shifts towards a higher

size range, while new M6C forms with a lower Cr content and smaller size • M23C6 and MC shift towards a higher size range.

3. After 15000 hours at 600°C: the following variations were observed (see Figure 3.4.9): • all M6C precipitates shift towards higher size • M23C6 and MC continue the coarsening process • M2C do not change size

4. After 1130 hours at 660°C: the size distribution is quite similar to that observed after the heat treatment (see Figure 3.4.10). It seems that an incubation time is necessary to start microstructural evolution:

• only a difference in the frequency of M2C and MC may be noted.

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(%)

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M2C

MC

M6C

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M2C

MC

M6C

Figure 3.4.9: Frequency of precipitates vs. Figure 3.4.10: Frequency of precipitates vs. size in Grade 23 after 15.000hours at 600°C size in Grade 23 after 1130 hours at 660°C Figures 3.4.11 shows the evolution of the precipitates in term of composition. The MC and M2C compositions are very similar and in these figures cannot be separated. The formation of new M6C during the aging seems to be the most interesting observation. The pre-existing M6C (with high Cr content) shifts towards higher sizes, while new M6C forms (as visible in Figure 3.4.8) with lower Cr content. The hypothesis is that in the as heat treated material, M6C forms on pre-existing M23C6 carbides, while during ageing, new particles form. The M23C6 disappears after 10.000 hours of aging as predicted by the thermo-dynamical tool JMatPro.

P23 600°C 6425 hours

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precipitates size (nm)

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y (%

)

m23c6m2cmcm6c

P23 base

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m23c6m2cmcm6c

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0.00 0.25 0.50 0.75 1.00

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P23 as treated

Fe+W

MC M6C M23C6

Mo+V

+Ti+

Nb

Cr0.0 0.2 0.4 0.6 0.8 1.0

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MC M6C M23C6

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+Ti+

Nb

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Fe+W

MC M6C M23C6

Mo+

V+Ti

+Nb

Cr

0.0 0.2 0.4 0.6 0.8 1.0

0.0

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P23 625°C 10000h

Fe+W

MC M6C

Mo+V

+Ti+

Nb

Cr0.0 0.2 0.4 0.6 0.8 1.0

0.0

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P23 660°C 1130h

Fe+W

MC M6C M23C6

Mo+V

+Ti+

Nb

Cr0.00 0.25 0.50 0.75 1.00

0.00

0.25

0.50

0.75

1.00 0.00

0.25

0.50

0.75

1.00

P23 675°C 10000h

Fe+W

MC M6C

Mo+V

+Ti+

Nb

Cr

Figure 3.4.11. Evolution of the composition of the precipitates for Grade 23 from as treated material up to 10.000 hours aged at 675°C.

Figure 3.4.12 shows the frequency and dimensions of extracted precipitates from differently aged specimens.

a) P23 625°C – 3000h b) 625°C – 10000h c) 675°C – 10000h

Figure 3.4.12: examples of extraction replicas from aged grade 23 specimens

For grade 23 the quantitative analysis of particle dimensions for each precipitate type has been summarised in figure 3.4.13. It has to be taken into account that this diagram shows a tendency of microstructural evolution, but it is important to note that the prediction of the thermo-dynamical model seems correct. This activity will be continued after the end of the present program, when other more aged specimens from creep tests will be available, in order to increase the microstructural database.

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0

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400

450

20,00 20,50 21,00 21,50 22,00 22,50 23,00

PLM

Mea

n D

iam

eter

M23C6MXM2XM6C

Figure 3.4.13: evolution of the diameter (nm) of the particles as function of PLM (time and

temperature of aging): the values at PLM 20 are the as treated dimensions Similar analysis has been made on grade 24 specimens for comparison. Grade 24 (figures 3.4.14):

a. as heat-treated material: carbides M6C, M23C6 and MC are present after the heat treatment. MC are practically V-Ti carbide, while M6C precipitates have a very high chromium content.

b. After 10000 hours at 625 and at 675°C: the presence of the M23C6 is strongly reduced. For grade 24, the results of the TEM analysis also fit quite well with the predictions of the thermodynamic tool. During grade 24 aging, it appears that the M6C precipitates change their compositions with respect to the as heat-treated particles. The pre-existing M6C have a higher Cr content than the M6C precipitates after aging. In figure 3.4.15, it can be seen that the dimensions of the precipitates are very small after the heat treatment, but coarsening is faster than in grade 23. Comparing figures 3.4.12c and 3.4.15c, both from aged samples at 675°C for 10000hours, it may be noted that in grade 24, some particles reach a 3-5μm diameter. Table 3.4.3. shows the mean diameter values (nm) of all the precipitates analysed. The dimensions of M6C seem to be quite stable, but the initial dimension is larger then in Grade 23. Also, the coarsening of M23C6 is significant.

PLM MC M23C6 M6C As received 48 144 376

21,55 50,5 334 22,75 183 546 376

Table 3.4.3: evolution of dimensions for grade 24 (dimensions nm)

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0.00 0.25 0.50 0.75 1.00

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Fe+Mo

MC M6C M23C6

V+Ti

Cr0.0 0.2 0.4 0.6 0.8 1.0

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P24 625°C 10000h

Fe+Mo

MC M6C

V+Ti

Cr0.0 0.2 0.4 0.6 0.8 1.0

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P24 675°C 10000h

Fe+Mo

MC M6C M23C6

V+Ti

Cr

Figure 3.4.14. Evolution of the composition of the precipitates for Grade 24 from as treated

material up to 10.000 hours aged at 675°C.

a) P24 as treated b) 625°C – 10000h c) 675°C – 10000h

Figure 3.4.15: examples of extraction replicas from aged grade 24 specimens Task 3.4.2: Characterization of aged material 3.4.2.1 Ongoing assessment of microstructural evolution during ageing at 550, 600 and 650°C at 1k, 3k, 10k hours: type, chemical analysis, dimension, density, The material aged by VTT is presented in table 3.4.4 and their respective specimen (micrograph) codes in table 3.4.5. The material codes are reported in work package WP1. An example of the microstructural ageing of P23 is presented for material G in Figure 3.4.16. Results As new material In as new condition the microstructure in both examined steel types was tempered martensite or bainite. The inherent austenitic grain size of steel P23 was about ASTM 6 and in the examined material P24 the austenitic grain size was about ASTM 8. The lath structure in both steels was fine and the structure contained small carbides at lath boundaries and considerably larger carbides at the prior austenite grain boundaries. The resolution power of microscopy was not high enough to reveal the precipitation strengthening small MX-type carbonitrides inside the laths. Polished and etched metallographic cross sections were prepared from the annealed samples as well as from the as received material. The cross sections were examined and microphotographed with an optical microscope and a scanning electron microscope (SEM) equipped with an energy dispersive X-ray analyzer (EDS). To detect the carbide changes in as early a stage as possible, the microstructure of P24 steel was studied from the coarse grained HAZ, in which the effect of original tempering is smaller than in the base material of the tube.

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For the material G, it is now known that ageing causes coarsening of the lath structure and growth of the carbides located at austenite grain boundaries. Furthermore, it causes partial disappearance of lath boundaries and corresponding lath boundary carbides. The carbides located at the remaining lath packet boundaries do not show clear indications of coarsening.

Type Temperature Exposure time (h) 675°C 100, 300, 1000, 3000, 10 000 Aging test 625°C 300, 1000, 3000, 10 000

Table 3.4.4. Ageing test matrix for materials G, C and M

PLM 20.18 20.65 20.86 21.08 21.31 21.55 21.80 22.26

Temperature (°C). 625 625 675 625 675 625 675 675

Time (aging, h) 300 1000 100 3000 300 10000 1000 3000

G G-3 G-5 G-1 G-7 G-2 G-8 G-4 G-6

C C-3 C-5 C-1 C-7 C-2 C-8 C-4 C-6

M M-3 M-5 M-1 M-7 M-2 M-8 M-4 M-6

Time (sim., years) 0.7 2.6 4.5 8.3 15.3 29.7 58.9 201.0

Table 3.4.5. Specimen/micrograph specific codes for the ageing tests in Table 4.4. The table is in the order of ascending Larson-Miller time temperature parameter PLM (C=20, PLM/1000), showing increasing “age” (=simulated time in years) in power plant service conditions (575°C).

Micrograph maps, arranged in terms of as temperature compensated time (Larson-Miller), have been produced in VTT laboratories. The Aloas parent materials G (P23) and C (T24) are compared with material M parent, weld and HAZ (an older heat of welded P24) in Appendix 1. Changes in carbide composition EDS-analysis of the metallographic samples showed in the originally iron-rich carbides located at the austenitic grain boundaries of P23 steel an increasing content of tungsten (W) after long term annealing in 675°C. Scanning electron microscopy using backscattered electron (BSE) techniques did not, however, reveal separate carbide types before being annealed in 675°C for 10000h. In this sample BSE images revealed two types of carbides having clearly different average atom weight (Figures 3.4.17). Part of the carbides have due to a high average atom weight a white appearance in Figure 3.4.17, while the appearance of the majority of the carbides is gray, indicating a smaller atom weight. EDS-analysis of the sample revealed that carbides with white appearance in Figure 4.17 are tungsten-rich carbides, apparently M6C-type carbides, whereas the gray carbides are iron-rich, apparently M23C6-type carbides. Microscopy of samples prepared from P24 steel showed precipitation of separate different types of carbides in a clearly earlier stage. In this steel, carbides with clearly different atom weight were observed after annealing in 675°C for 1000h. The carbides having a high atom weight and white appearance in BSE-image were in EDS-analysis observed to be due to different composition of this steel molybdenum-rich carbides, apparently also in this case M6C-type of carbides and the lighter carbides with gray appearance iron rich M23C6-type carbides. Carbide coarsening Optical microscopy of the steel samples showed clear carbide growth after annealing at high temperature, but due to the limited resolution power of an optical microscope, the observations were only qualitative. Examples are shown in Figures 3.4.16. Quantitative measurements of carbide coarsening were performed with scanning electron microscopy and an image analyzer program. Usually in particle size analysis the average values

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are very sensitive to magnification and other details of microscopy due to the strong effect of large number of the smallest particles in the structure. An additional difficulty in the carbide size analysis was the observation that carbides often grow very rapidly along the grain boundary line. To avoid these difficulties it was decided to use the width of the largest carbides as a measure of carbide coarsening. Mapping and measurements of carbide growth at the grain boundaries indicate continuous initiation (or emergence across the detection limit) during aging. This will complicate modelling of the size distribution. However, clear growth of the apparent effective diameter of the grain boundary carbides was observed, as can be seen in Figure 3.4.. More detailed distributions have been obtained with TEM at higher resolution on selected specimens (see 3.4.1.2). These two techniques; SEM and TEM, to analyse the precipitates are not directly comparable; therefore the two techniques have to be taken as complementary.

PLM/1000 Microstructure PLM/1000 Microstructure

as received

20.65

20.86

21.08

21.31

21.55

21.80

22.26

Figure 3.4.16. Microstructural evolution during simulated high temperature service for material G, shown as a series with increasing PLM.

As received (effective diameter 0.34 µm) 10.000 h aged (effective diameter 1.05 µm)

Figure 3.4.17. Precipitation coarsening at the grain boundaries for P23 (material G), as received versus 10 000 h aged at 675°C

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A new methodology has been recently developed by CSM for analysis of precipitates in steels and alloys. The iron matrix is electrolytically dissolved by a solution of HCl in ethyl alcohol, then the solution is filtered and then dried. The extracted particles are then analysed by X Ray Diffraction. The results of the X-RD analysis obtained on as heat-treated grade 23 are compared in figures 3.4.18 and 3.4.19 with the particles extracted from the samples aged for 10.000 hours at 550°C and 600°C respectively. The red arrows indicate precipitate types which give an increased number of counts after aging, while the blue arrows show precipitate types which give a reduced number of counts. It is possible to observe that: - the analysed type of precipitates are mainly the same as predicted by JMatPro and observed by STEM+EDS; only the M7C3 was not analysed by TEM on the as received specimen. Since the M7C3 is a meta-stable precipitate, it evidently dissolves during aging - the difference between the peaks of the as heat-treated material and the aged specimens confirm the evolution of the particles. The amplitude of M6C peaks in the aged specimens is a clear indication of the increase of number and coarsening of this kind of carbide. Currently this methodology is still under development with the aim to generate quantitative analysis of the extracted particles.

36 41 46 51 562 theta

coun

ts

QT_to550_10000

M23C6

M23C

M7C3

M6C

M23C6

M7C3

M6C

M6C

M6C M7C3M23C6

M7C3

M23C6

M7C3

M7C3

M2C

M2CM6C

MC ?

MC ?M6C

Figure 3.4.18: comparison of X-RD particle analysis of grade 23 specimens as treated and aged

at 550°C for 10.000 hours 3.4.2.2. Characterisation of the aged materials mechanical properties as above: aging at 550 - 600°C up to 10k hours. Tensile, impact tests. Specimens of both grade 23 and 24 have been machined to perform mechanical tests after aging for different times and temperatures. Hardness tests have been reported in Figure 3.4.1. Tensile tests and impact tests results can be summarised in the following figures:

- Figure 3.4.20 shows that the mechanical properties at room temperature of grade 23 have a reduction after 10.000 hours of aging at 550°C; grade 24 shows little increase of tensile properties during aging to reach the same values of the as treated material after 10000 hours;

- Figure 3.4.21 and 3.4.22: the diagram shown the trends of UTS and YS results at 550°C and 600°C for Grade 23 specimens, aged up to 10.000h at 550°C and at 600°C.

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36 41 46 51 562 theta

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M6C

M6C

M6C

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M7C3

M23C6M7C3

M7C3M7C3

M23C6

M7C3

M23C6

M7C

M2C

M2C

MC ?M6C

MC ?

M6C

Figure 3.4.19: comparison of X-RD particle analysis of grade 23 specimens as treated and aged

at 600°C for 10.000 hours Figure 3.4.20: T23 and T24 tensile test results at 20°C after aging at 550°C up to 10.000hours

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Figure 3.4.21: T23 tensile data at 600°C after aging at 550°C up to 10.000hours

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Figure 3.4.22: T23 tensile data at 600°C after aging at 600°C up to 10.000hours

Grade 23 material shows better impact properties after heat treatment than grade 24, and also maintains better impact properties after aging. The grade 24 impact energy values drop sharply toward zero after aging up to 10.000 hours at temperature of 550° and 600°C, with a tendency toward fully brittle fracture (figures 3.4.23, 3.4.24). These behaviours may be consequences of the different dimensions of the stable precipitates observed after aging.

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Figure 3.4.23: Comparison of impact values at RT of grade 23 and 24 specimens aged at 550 and

600°C up to 10.000 hours

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Figure 3.4.24: Comparison of brittle fracture at RT of grade 23 and 24 specimens aged at 550

and 600°C up to 10.000 hours 88

Conclusions of Work Package 4 The assessment of the microstructural evolution of both grades 23 and 24 has been performed by aging tests carried out in the range of temperatures 550-675°C up to 10.000 hours. For both grades, the results obtained in term of microstructural observation seem quite well in agreement with the phases predicted at equilibrium by the JMatPro thermodynamic tool. However due to differences in precipitation and coarsening kinetics also metastable phases may develop during aging. In grade 23, all the metastable types of precipitates like MC, M23C6 and M7C3 show a tendency to change their size distributions and to form the M6C carbide, which is the more stable phase at equilibrium. In grade 24, the situation is a little bit different because at equilibrium the M23C6 carbide is also a stable phase with a faster growth then M6C. In addition to changes in carbide structure increasing thermal exposure causes gradual coarsening in the lath structure. Obviously, during the aging times performed, the true equilibrium stage has not been reached. However, the aging tests at 675°C for 10.000 hours may approximately be equated, on a time-temperature parameter basis, to an exposure time of 200 years at 575°C. This is the maximum proposed service temperature for these low Cr grades, due to their oxidation rate. Regarding the variation of the mechanical properties during the aging at 550°C up to 10000 hours, the tensile properties at RT (YS and UTS) are practically not affected by exposure. However, the tensile properties measured at 550°C and 600°C shown a reduction of about 200MPa in each case. Different behaviour is observed for the impact values. Grade 23 appears to be much better then Grade 24 after heat treatment and also after aging: the Grade 24 material reaches values lower then 50 Joule with brittle fracture. This could be related to the larger dimensions of the precipitates that are present in the Grade 24 steels after aging. The availability of more aged specimens from crept samples which are still running in this project, and in other EU programs such as COST 536, will be very useful to obtain more information on the microstructural evolution of both Grade 23 and 24.

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3.5 Work Package 5: High Temperature Design and Assessment

Introduction

This work package provides an assessment of the durability of proposed new welding consumables for P23 and P24 steels. Where possible this assessment takes into account pragmatic experiences on UK plant with CMV material specified systems. The scope of this assessment is shaped by recent UK experiences with thermal ‘bore cracking’ on CMV pipework systems, which in 2001 was originally confirmed as a threat to pipework integrity in the UK. During early June 2001, a new type of main steam pipe weld cracking was found on a UK station. The cracking initiated from the weld root or adjacent machined recess corner and grew radially outwards, in the worst cases, to a roughly uniform depth in excess of 20mm, fully circumferentially around the bore. The cracking has been found in CMV main steam pipework, which is of nominal 240mm bore and 60mm wall thickness. Subsequent metallurgical investigations suggested that the crack growth mechanism occurs initially due to thermal fatigue, with cracks initiated relatively early in the plant life. The pipe girth welds had been inspected during previous statutory outages, however this particular type of weld defect was undetected, despite its size. This was attributed to the nature of the cracks, which meant that they were difficult to find using normal CMV inspection procedures, hence enhanced procedures for detection and sizing were developed. On the initial detection of these defects an extended inspection campaign was initiated across all UK power plants. This revealed many instances of bore cracking, which necessitated the removal of affected welds or in some instances continued operation underpinned by a safety case. The instances of weld bore cracking tended to mainly occur in clusters either at the top of the main steam line, adjacent to the main boiler stop valves, or further downstream adjacent to the steam chests. In some stations the cracking was localised, in others it occurred at several welds along the main steam line. Crack growth rates of 2mm per year due to thermal fatigue are not uncommon. Of more significant concern are cracks that have already propagated a significant way through the pipe wall, which are the subjected to a high risk of accelerated crack propagation to creep crack growth.

At the time of writing, main steam bore cracking still occurs in the UK, but is now considered a ‘managed’ issue, with strategies in place to manage the integrity of affected plant with defects of various sizes [25, 26]. However, the experiences with main steam bore cracking on UK main steam CMV pipework has again emphasised the following important issues:

• Successful outage inspection and return to service requires knowledge of expected defects, hence an understanding of the risk of defect initiation and propagation is a prerequisite.

• Generating utilities will not necessarily operate plant in accordance with design intent, particularly if the plant is operated in a commercial environment.

• Safety is paramount in the eyes of the plant owner. • There is no substitute for real service experience, and it is essential that such

experiences are used to shape the requirements for the development of future materials, and construction/repair methods.

• Even considering the extensive service experience gained with CMV pipework and weldments; there are many aspects associated with ‘integrity’ management that are still being addressed as emerging issues.

Hence, the above experiences has emphasised the need within the scope of this project to undertake some assessment of the likely durability of the proposed weldments and to compare this against plant experiences with CMV pipework installations. It is within this framework that the tasks in this work package have been undertaken; hence comprise a mixture of analytical and pragmatic assessments utilising service experience.

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Overview of Work Package3.5 Tasks The work package has been split into three specific areas, which are intended to provide an assessment of the durability of the low alloy steels under realistic plant operational conditions. Where possible, comparisons against CMV durability have been made. Task 3.5.1: High Temperature Design Review. This is intended to review components that the partners consider are the most suitable for simulation purposes. The overall aim of the work package is to provide evidence that the use of P23 or T23 materials in practical power plant installations will be at least as durable as materials such as CMV, under typical plant operation conditions. In the originally submitted plan it was intended to undertake a design review of:

• Task 3.5.1.1: Seamless girth welded pressure vessels and pipework • Task 3.5.1.2: Seam welded pressure vessels and pipework

During the project meeting in Rome on the January 29th, 2004 it was agreed that it was of less importance to devote time to the assessment of seam welded installations, due to the predominance of girth welded designs in the European Union. Hence task 3.5.1.2 was not pursued. Task 3.5.2: Structural Assessment of the effects of typical plant operating conditions. This work package task will provide an assessment of the durability of selected components highlighted from the review in Task 5.1. The intention was to model nominated components, and assess their durability to withstand;

• Repeated thermal cycling based on typical plant loads, including fault events. • Assess the effects and sensitivity of pipe supports and their location, which can have a

significant influence on creep crack growth rates in damaged girth welds. The work package task will draw upon recent experiences in the UK with main steam bore cracking. In addition boiler flow models can be used to assess the effects of temperature and pressure distributions on the structural response. The subsequent work schedule was identified as:

♦ Task 3.5.2.1: Assessment of the effects of typical plant operating cycles on the structural integrity of pipework.

♦ Task 3.5.2.2: Assessment of pipework system supports on the structural integrity of pipework.

♦ Task 3.5.2.3: Assessment of operational temperature and pressure distribution and flow, in a selected boiler case.

Task 3.5.3: Experimental Validation of Component Simulation The intention was to provide verification of the generic simulation models assessed as part of Task 3.5.2. These tests were conducted on tube material forms. The subsequent work schedule was identified as:

♦ Task 3.5.3.1: Material modelling for P23 and P24 materials (base metal, HAZ, weld metal)

♦ Task 3.5.3.2: Verification of the simulation work by simulation (stresses and deformation) and lifetime prediction of one tube test.

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♦ Task 3.5.3.3: 6 tests on tubes with similar and/or dissimilar welds under internal pressure and cyclic thermal loading over 1 to 2 years (depends on component availability).

3.5.1: High Temperature Design Review

Initially the following components/geometries were considered for detailed analysis within the work package: • Main Steam girth welds, which are subject to established damage mechanisms such as bore

cracking, creep damage and Type IV weld cracking. In addition, there is also a large amount of available data on plant operational loads and service experience. There is a potentially large market for pipework replacements due to the age of UK plant and the requirements of the Large Coal Plant Directive (LCPD) that imposes operational limits on many stations. Consequently the UK will shortly enter a phase of New Build and re-lifing of existing plant that have not opted out of the LCPD.

• Headers on UK plant are regularly prone to interstub ligament cracking, in addition there

have recently been an increasing number of problems regarding the life expectancy and performance of P91 steels.

• Small bore and large bore branch geometries are susceptible to cracking at the weld,

primarily due to the action of external pipework loads, and also on the internal bore where the branch penetrates the main pipe, which is a similar mechanism to header interstub ligament cracking.

Following a review of the above it was agreed that the best use of available resources would be to concentrate on the analysis of main steam girth welds. It is useful at this juncture to compare some of the typical main steam pipework dimensions on UK conventional and combined cycle gas turbine (CCGT) plant.

• Conventional Units: Outside diameter 360mm, wall thickness 60mm • CCGT Units: Outside diameter 284-356mm, wall thickness 20-46mm

An extensive review of UK plant operational data from the E.ON UK archive has been undertaken, both for steady state and cyclic operation. This is a useful review since it provides some of the core operational data used to set the framework for subsequent material testing undertaken in other work packages. The data presented here is not intended to be restrictive, more importantly it is intended to show the far ranging operational conditions that in-service plant is expected to cope with. In many cases these operational data is outside original design specifications. Figure 3.5.1 shows a thermal cycle for illustration, based on typical operational transient steam temperature data obtained from within the main steam flow, at the top of each of the four main steam lines as they exit the final superheater header stage. Thermal transients of this type are the main cause of thermal bore cracking. This data review has shown that very severe thermal transients can occur, and some of the main findings of the survey are outlined below: Boiler Stop Valve (BSV) Opening These transients occur due to opening the BSV when the steam temperature is cooler that the surrounding CMV metal temperature, which is typical of hot starts.

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The largest transient rate recorded was 380oC/minute, over a duration of 30 seconds, consequently the through wall temperature differential is conservatively estimated to be 190oC.

Figure 3.5.1. Typical Main Steam Line Temperature Transient On-Load Cycling The effects of thermal cycling whilst the plant is operational at full load are effectively attenuated due to the cyclical frequency and relatively modest thermal transient rates. The usual cause of on-load cycling is due to poor control of attemperator sprays (overspray followed by underspray) and firing (overfiring followed by underfiring). On-load cycling due to poor spray control can be quite severe immediately downstream of the attemperators. However, by the time the steam has reached the final superheater outlet the effect has been mitigated by the mixing of the steam, additional heat input from the final superheater stages and heat being absorbed by the mass of metal between the attemperators and final outlet. The on-load cycling illustrated in figure 3.5.1 is relatively benign because of its cyclical nature, relatively modest rates (20oC/minute) and durations of circa 2-3 minutes. It should be noted that there is a ‘transit time’ [27] for any thermal transient acting on the pipe inner bore surface to travel through the pipe wall and thereby affecting the outer surface temperature. Consider a typical CMV pipework section of 360mm outside diameter and 60mm wall thickness subjected to an on-load thermal transient cycle of 20oC/minute, of two minutes duration. The full 40oC temperature differential would not occur, because the thermal shock load applied to the pipework bore surface will take approximately 45 seconds to travel across the pipe wall, at which time the outside wall also begins to cool. In addition, at the end of the thermal transient downshock the subsequent increase (upshock) in steam temperature would tend to further reduce the effective through wall temperature differential. There have been no observed occurrences where an on-load transient has resulted in a through wall temperature differential that exceeds those from operation of the BSV. However poor attemperator/firing control has led to some on-load transients being observed that results in a conservative estimate of through wall temperature differential of 120oC. It should be noted that this event occurred at a rate of 12oC/minute over a 10 minute period. The effect of the long transient duration would tend to alleviate the effects due to a reduction in the outer surface temperature as described above.

Downshock

On-load Oscillation

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Forced Cool (Depressurisation) One instance has been recorded where a severe downshock of 130oC/minute occurred over a 1 minute period occurred due to the depressurisation of the boiler during a forced cool. Drains Operation Incorrect operation of pipework drains is commonly associated with internal pipework cracking at penetration features. The most severe transient observed due to incorrect operation of the drains was 21oC/minute and lasted for 11 minutes. Again, the effect of the long transient duration would tend to alleviate the effects due to a reduction in the outer surface temperature. In summary, this review of E.ON UK’s experience with plant transients has shown that all plant, CCGT and conventional fossil fired, can be subjected to sharp thermal transients, which can be of the order of 75-100oC/min, over one minute duration. This should be of less concern for the thinner CCGT pipework geometries; however it is necessary to compare weld performance of potential P23/P24 replacements against existing weld materials and geometry.

E.ON UK was tasked to provide recommendations on the realistic plant stress level which should be covered within the test programmes. The following describes the recommendations made and the rationale used to obtain them. A simple way to determine the realistic plant stress is to decide what temperature future plant will be designed to operate at, and then refer to design codes which govern the conditions to be specified for a design life of, for example, 100,000 hours. For high temperature operation, these are based on extrapolated creep rupture data, and incorporate a safety factor. ASME design data have been cited by V&M [28]. A common current assumption in the industry is that because of oxidation, the limiting plant temperature for P/T23 is 575°C. This may be unduly conservative for thick section P23, where oxidation should be a lesser issue. However, the 575°C temperature can be used as a starting point for planning. The ASME mean 100,000 hour creep strength for T23 at 575°C is 106 MPa. This is of course based on considerable data extrapolation and may be somewhat unreliable. The ASME maximum allowable stress at 575°C (interpolated) has recently been updated in the 2007 edition, as per code case 2199-3, to 70 MPa, which includes the safety factor. As E.ON UK steam pipework operates at up to 568°C, the 575°C design stress values should be slightly conservative in respect of potential plant replacements. However, there is a case for extending the temperature limits to be considered, both to examine whether P23 might be a suitable alternative e.g. for P91 at temperatures above 575°C, and because in practice, it may sometimes be necessary to use steels in plant at higher temperatures and lower stresses. It should be noted here that it is not unusual to find that on a typical 500MW Unit, which has four main steam lines; that steady state temperatures can vary by 5-10oC between steam legs. Hence, considering the impact that this will have on creep life consumption, it should not be assumed that inspecting or monitoring the condition of one of the main steam lines will be indicative of the condition of the other legs. This fact accentuates the need to understand the risks associated with the accumulation of material damage or defect initiation and propagation for both parent material and weldments with consideration of real plant operating conditions. This emphasises the requirement to consider performance at temperatures beyond anticipated design levels.

3.5.2: Structural Assessment of the effects of typical plant operating conditions.

The aim of this task is to provide a comparison of and to comment on the structural performance of a typical pipework system constructed in P23 material and one from conventional CMV. The following has been considered:

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• Defect tolerance, with respect to thermal fatigue cracking. • Effect of pipework hanger loads and consideration of mal-operation, as experienced on

operational plant. • Refinements in the approach to modelling steam flow transients, and the subsequent

effect on pipe stress profiles and the persistence of a thermal transient along a steam pipeline.

3.5.2.1: Assessment of the effects of typical plant operating cycles on the structural integrity of pipework A defect tolerance assessment of a typical P23 pipework girth weld containing fully circumferential radial bore cracks, of different depths has been completed and compared with the response of a typical CMV pipework geometry. In this assessment the following pipework section geometry has been used:

CMV 360mm OD 60mm wall thickness P23 356mm OD 36mm wall thickness

The P23 pipework geometry, (Ref Dalmine internal order 1013732/1), is used because it had the closest internal diameter to the current main steam lines, and under operating pressure would have approximately a similar creep life to the current design. Additionally, from a preliminary elastic system analysis of the main steam lines this pipe produces the closest allowable stress ratios to the current system and has been used in the subsequent pipework system model described in Task 5.2.2. The pipework welds were subjected to the following loads, which represent typical plant conditions:

Operating pressure 176bar Pipework system load 10MPa Crack face pressurisation Steam thermal transients ranging 50-200oC/min

Figure 3.5.1 illustrates a typical steam transient that occurs when the boiler comes onto full load and results in a tensile stress on the pipe bore. Table 3.5.1 details expected magnitudes of axial stress due to the thermal shock for both the P23 and CMV geometries. The 75oC/min transient represents what may be described as a ‘normal’ operational transient, typical of UK plant experience. The 150oC/min transient represents a ‘severe’ fault transient, which again has been observed on UK plant. Figure 5.2 shows the typical response of the P23 and CMV pipework to these ‘normal’ and ‘severe’ thermal transients. The factor of safety is obtained from a series of analyses using the R6 defect assessment code [29].

Transient Magnitude (oC/min) P23 CMV 75 135 142

150 270 285 Table 3.5.1. Peak Axial Stress at Pipe Bore

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Figure 3.5.2: P23 versus CMV - Circumferential defect tolerance to typical plant transients, based on a factor of safety Clearly for the ‘normal’ operational transient, in this case 50oC/min, the P23 pipework has a higher factor of safety for all crack depths. However, when the pipework is subjected to a ‘severe’ or fault transient then the benefits are markedly reduced, in particular for deeper cracks beyond half wall thickness.

3.5.2.2: Assessment of pipework system supports on the structural integrity of pipework. Pipework systems in service experience a range of operating conditions, which can be considered outside the original design intent. In particular, experience during outage overhauls has shown that the adjustment and proper maintenance of pipework hanger supports, in addition to the operational conditions, has a significant effect on the life of the parent pipe and weldments during service. Consequently models using CMV material and P23 have been generated with the purpose of comparing maximum stresses and bending moments in the pipes, and the effects of these on the support system. In addition the effect of some limited cases of ‘hanger’ failure, based on typical scenarios experienced on UK plant has also been assessed. Initially, two CMV main steam lines were identified for use, representing real pipework systems on an E.ON UK coal fired power station and covering the full run of pipework from the boiler superheater outlet header to the turbine inlet. The pipework runs from the boiler superheater header outlet to the turbine inlet leads to a drop in height of around 42m and extends away from the boiler by a distance of around 40m. The pipework is supported at regular intervals by a total of 13 supports and is fully insulated and clad. For the purposes of this assessment, a P23 pipe of outer diameter 355.6mm and a wall thickness of 35.71mm were chosen for the reason described in Task 3.5.2.1. The preliminary analysis was undertaken using PSA5 [30], which is a pipework stress analysis code, used for static and dynamic analysis and code compliance checks of pipework systems. Subsequently pipework hanger loads were calculated, which therefore provided all the boundary conditions for the construction of a detailed finite element model of the P23 pipework system.

Potential problems can occur when the hangers are operating incorrectly and pipework loads can become significantly higher than the design conditions. There are a number of different problems that can occur with the different hanger types, but with the constant effort hangers three of the typical problems are:

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1. Broken tie rod between hanger and pipework, i.e. the hanger provides no support,

effectively removed from system, 2. Hanger is damaged or incorrectly set, i.e. the hanger provides wrong level of support,

constant effort is higher, or lower than design, 3. Locking pins not removed from hanger, i.e. the hanger acts as a rigid support, no vertical

movement. Additional finite element models of the pipework systems were developed in order to provide additional insight into the potential in-service performance of a P23 pipework system. The intention was to develop comparable models of a P23 and CMV pipework system, which would support the service data review and defect tolerance assessments summarised in earlier sections. Where possible and within the scope of the project, the modelling accounts for known observations associated with plant performance. Finite Element Modelling VTT generated a pipe shell model in Abaqus software [31] that is based on typical UK pipework installations. E.ON UK provided representative geometry and operational data to support the analyses. The aim was to analyse the pipe structural response to typical operational conditions. In addition one load case was analysed that represented fault conditions, caused by extreme plant operating conditions and instances where the supporting pipework hangers have malfunctioned. This provided data to support evidence for the service tolerance of a P23 pipework system. First, the temperatures and pressures for both CMV and P23 model were calculated with APROS 5.05 process simulation code (see Section 3.5.2.3). For the FE modelling of the typical steam transient specified for the case study, the specified pipe line shown in figure 3.5.3 was modelled. The preliminary analyses with hypothetic loads showed satisfactory performance of the model. The pressure and temperature transients were transferred from the APROS code to the Abaqus code and preliminary elastic structural analyses with those transient loads were successfully performed and they showed satisfactory performance of the model. Both a coarse and a fine shell element mesh were created but the final analyses were conducted only with the coarse mesh. It is fine enough to give correct global results and quite accurate local stresses and strains. Elastic-plastic material properties were gathered partly from EU norms and partly from the data provided by E.ON. No creep was included in the material model. Two different kinds of initial states were considered. The initial temperature was either operational or room temperature. In fact, the initial temperature means also the stress-free temperature in these kinds of analyses. Thus, the stresses due to thermal expansion were different in the two cases. However, gravitation seems to have a greater effect on the structural behaviour of the pipeline than the temperature variations. In this kind of analyses it is always troublesome to decide, whether the stress-tree temperature is room temperature or operational temperature, consequently, idealising the whole pipework system is not that straightforward. For example, the pipes have residual stresses due to the manufacturing process and during service operation the stresses even out and relax at high temperatures. These analyses were not dependent on time, since there is no creep behaviour included and the heat transfer was modelled already with APROS process simulation code. The time-varying temperatures from the APROS code were applied to the FE model in a way that they were constant throughout the thickness of the pipe wall. That is not a completely realistic assumption, but a close one. Especially in the case of CMV model which has a thicker pipe wall there can be

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significant temperature gradients through the wall. Those gradients induce thermal stresses which are not considered in these analyses. In other words, the temperature in these analyses varied in time and in the axial direction (the pipeline being divided into 28 parts in that sense), but not in the radial direction. The heat transfer analyses can be conducted with Abaqus, but this is time-consuming and did not fit in with the scope of this project. The simulations were conducted in two steps. Initially, there were no forces or pressures applied, and the temperature is either RT (20oC) or 568oC. The model was completely stress-free. In the first step, the temperatures and inner pressures were shifted to the initial state of the operational transient simulated by APROS (approx. 355oC and 16.5 MPa, respectively). These conditions are representative of those presented in figure 1, obtained from a UK Power Station during service. In this same step, the gravitation and the hanger forces were ramped up linearly to their given values, which again represent in-service values. In the second step (duration 9960 seconds), the above mentioned gravitation and hanger support forces were retained, and temperature and pressure transient values were applied to the sections. Naturally, the first step gave already the highest stresses. After that, they just oscillated according to the transients. The valve was modelled simply by making the pipe wall thickness in that section very large. That made it “locally rigid” so that it still could freely move in space. The stresses near that section were not very reliable. It should be noted that loads associated with ‘cold pulls’, which is an additional static load applied to the pipework during final installation on site, were not included in this study. The pipeline was not fixed completely rigidly at its ends, since in that case the ends would have had notable plastic deformations and also would have had a slight effect on the energy balance in the analysis. One accident case was considered for the both pipework models. The hanger at position number 14 on the pipework section, which corresponded with the highest hanger force, was removed. The CMV model became a mechanism when the pipe ends became pivots by having very large plastic deformations. Otherwise the results were sensible. Some main results of the both models are compared below. Figure 3.5.4 shows displacement magnitudes [m] at the end of the steam transient of CMV model (left) and P23 model (right). The displacements are scaled up by a factor of 10 in the deformed shape of the model for illustrative purposes. Figure 3.5.5 shows the distribution of Von Mises stress [Pa] on the inner surface of CMV model (left) and P23 model (right) at the end of the steam transient. The location is from the boiler end. The initial condition in this case was the operating temperature 568oC. Figure 3.5.6 shows the Von Mises stresses on the outer surface of the pipe wall on a certain element (8872), as functions of time. Figure 3.5.7 shows the vertical displacements of a certain node (11963) as functions of time. Different cases are compared with each other.

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Figure 3.5.3: The model geometry with hanger forces on the left. Part of the finite element mesh

on the right. The model presented can still be improved further by consideration of the following. First of all, a model with pipe and elbow elements would be the most adequate for this kind of problems. Cold pulls and hanger forces could then be modelled much more easily and the analyses would be much less time consuming. Shell or solid element models are necessary when accurate stress distributions have to be evaluated. Stress free temperature should preferably be (closer to) the operational temperature (550oC). The loads could be separated to different consecutive steps in order to better monitor their effect on the pipe system and to simulate more realistically the situation in the power plant. However, this has only a minor effect on the final state of the system (that is mainly used for the comparison of the different models). The heat transfer through the thickness of the pipe wall could also be modelled. This study has shown that although the bulk stresses for the P23 pipework system shown in figure 3.5.6 are higher than for an equivalent CMV system, the implications are that they are still acceptable, when considering the creep response of P23 parent material and the defect tolerance review provided in previous sections. In addition, it should be noted that the implications of pipework hanger malfunctions are no worse than have been experienced on UK plant on CMV pipework systems.

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Figure 3.5.4: Displacement magnitude [m] at the end of the steam transient of CMV (left) and P23 (right) models. The displacements are scaled up by a factor of 10 in the deformed shape of

the model for illustrational reasons.

Figure 3.5.5: Distribution of von Mises stress [Pa] on the inner surface of CMV model (left) and

P23 model (right) at the end of the steam transient. The location is from the boiler end. Note: In the legend for figures 3.5.6 and 3.5.7:

RT = initial state; room temperature (in all the other cases it is the operating temp.) ACCIDENT = accident case; one hanger (number 14) removed

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Figure 3.5.6: Von Mises stresses in the outer wall in element 8872

Figure 3.5.7: Vertical displacements of the node 11963. 3.5.2.3: Assessment of operational temperature and pressure distribution and flow, in a selected boiler case. The main steam line from the boiler superheater to the turbine inlet was modelled with APROS 5.05 simulation code. The six equation thermal hydraulics model of the code was used. Heat pipe and point components of APROS were used to describe the steam line. The steam line was divided into 270 nodes. The steam line model is presented in Figure 3.5.8 and the specifics are as listed in Table 5.2. Both CMV and P23 steam lines were modelled. The temperature and pressure transients were transferred to the FE code Abaqus.

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Model Parameters:

o 30 points defining pressures and temperatures o 29 pipe elements consisting nodes, branches and heat structures o Total number of flow nodes is 231 (average length ~ 0.5 m) o Heat structure nodes: 5 in pipe wall, 2 in isolation and 1 in liner

Boundary conditions: o Steam source, 165 bar, 568°C, 99 kg/s (variable) o Condenser 0,8 bar, 100 °C

6-equation model calculation speeds: o near steady state 2 times faster than real time o transient 2 times slower than real time

Table 3.5.2: APROS model parameters

A pipe insulation layer of 50 mm was modelled, which represents in-service conditions. Heat transfer to the surrounding environment was not assumed; hence the system was adiabatic. The temperature of the pipe was 350ºC at initial state of the transient, steam mass flow was 3 kg/s and inlet pressure was 165 bar. After the transient was started, steam mass flow and inlet steam temperature were controlled. Inlet pressure was constant during the transient. It was assumed that steam mass flow was constant for the first 3000 seconds and inlet temperature of steam was 430ºC. During the first 3000 seconds the steam temperature at the outlet of the steam pipe was increased from 350ºC to 400ºC. At a time of 3001 seconds a boiler start up was assumed. The steam temperature was reduced and steam at a lower temperature was fed in the pipe for 300 seconds. After that the steam temperature was increased to an operating temperature (540 ºC) in 4200 seconds.

Figure 3.5.8: APROS steam line model

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In Figure 3.5.9, the steam temperatures of a P23 pipework model are shown at seven points along the steam line. Location is shown as a distance from the pipe inlet. Steam pressures can be seen in Figure 3.5.10. The corresponding results for CMV model are quite similar.

Figure 3.5.9: Steam temperatures in the pipe (P23)

Figure 3.5.10: Steam Pressures (P23)

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Because the pipe wall was thicker in CMV model, the pipe had a larger heat capacity and the steam temperature at the pipe outlet increased slower when steam mass flow was low for the first 3000 seconds of simulation. Steam temperatures at the pipe outlet are compared in Figure 3.5.11.

Figure 3.5.11: Pipe outlet temperatures with CMV and P23 geometries

This flow simulation demonstrates that any imposed thermal transient will persist along the length of a typical pipework system. The persistence of the thermal transient has important implications for service inspection strategies. This suggests that it is not possible to focus girth weld inspections based on the assumption that the highest risk area is adjacent to the rigid BSV. On some UK stations, incidents of thermal bore cracking has occurred at seemingly random locations along the pipeline, which tends to support the potential effect of the persistence of the thermal transient along the pipeline.

3.5.3: Experimental Validation of component simulation

3.5.3.1. Material modelling The material model aims to describe the deformation and lifetime behaviour under typical thermo-mechanical fatigue loading conditions. The lifetime model is based on fatigue crack growth where the crack growth rate can be described in terms of the cyclic plastic deformation of the material. For many steels, the lifetime under cyclic or thermo-cyclic loading is controlled by the formation and growth of microcracks. For high temperature applications Riedel [32] proposed the DCF-parameter (CF for Creep-Fatigue). The damage parameter DTMF is an extension of DCF for non-isothermal cycles of arbitrary shape. Its functional form is

FD ine

e

IeffITMF ⎟

⎟⎠

⎞⎜⎜⎝

⎛Δ

ΔΔ

++

Δ= ε

σσσ

cy

2

cy

2,

σN314,2

Eσ45,1

.

where σcy is the cyclic yield stress, N the Ramberg-Osgood hardening exponent of the cyclic stress-strain curve and E the Young’s modulus. These parameters can be determined on the basis

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of the LCF-hysteresis. In order to characterize the cyclic plasticity, so called complex low cycle fatigue (CLFC) experiments where performed at different temperatures. These experiments cover strain rates from 10-5 to 10-3 1/s, hold times for relaxation and different strain amplitudes. A full characterization was performed on material H (Tenaris Dalmine heat 109194), which had been selected for the component test at an early stage of the project. A typical strain control and stress results for different temperatures are shown in Figure 3.5.12.

A B

Figure 3.5.12: Applied strain (A) and resulting stress at 550, 575 and 600°C (B) A Chaboche-type model for cyclic viscoplasticity was adjusted to the experimental results at different temperatures. By linear interpolation of the Chaboche model parameters a description of the cyclic plasticity of a thermo-mechanical fatigue (TMF) experiment could be obtained. The lifetime parameter DTMF [33] was adjusted to the experimental results. A uniform description of the LCF and TMF experiments lifetime in the range 1,000 to 15,000 cycles to failure in a scatter-band significantly smaller than factor two was achieved and shown in Figure 3.5.13.

Figure 3.5.13: Experimental and modelled cycles to failure of the specimen test on material H

with a scatter band of factor two. Spot sample of weld metal B323 where investigated longitudinal to the welding direction. They displayed a similar behaviour for LCF deformation and lifetime as material H.

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3.5.3.2. Lifetime prediction of the component tests A finite element simulation was performed to predict the cycles to failure of the welded component. The model was set up using three different material zones: base metal (BM), heat affected zone (HAZ) and weld metal (WM). The geometry of the zones was taken from the macro-section presented in section 2.1.1 of overall report. The material properties of the BM where taken from the material model for material H. The material properties of the HAZ where derived from the BM properties by increasing the equivalent model parameter of the ultimate tensile strength of the BM gradually across the weld thus matching the increasing hardness measured across the HAZ. The WM was modelled based on the BM material model with a reduced fatigue resistance. Two lifetime predictions where performed, one for each component test. For the first component test 6.000 cycles were predicted to reach a macroscopic (about 2 mm) crack. The predicted failure location is in the BM on the outer surface of the pipe, close to the HAZ. For the second component test the predicted lifetime (reaching a macroscopic crack) was between 110.000 and 230.000 cycles depending on the formulation of the damage model using the elastic-plastic and the pure elastic formulation, respectively. Again, the assumed failure location is the BM close to the HAZ on the tube outside. It should be mentioned, that for the second component test the extrapolation of the lifetime estimation is far outside the range for the experimental basis and therefore, the model is not validated for this range of lifetime. 3.5.3.3 Component tests The component test specimens consisted of two girth welded pipe H segments of same length. In order to reach the axial load levels with the capacity of the machine, the welded pipe walls were machined after post-weld heat treatment to a wall thickness of 6 mm. The pipes were loaded by internal pressure, superimposed cyclic axial load and a constant temperature. The pipes were heated by three heating cartridges with a power of 4 kW to obtain the desired test temperature of 450°C, which is representative of plant conditions associated with thermal bore cracking, figure 1. Internal pressure was applied from the opposite side to act in the gap between filler body and tube. The specimens were fixed in the testing rig via flanges with pre-tensioned screws for the application of cyclic loading and are surrounded by a bursting protection in case of unexpected break before leak. The complete testing facility consists of the servo hydraulic testing machine with the tube as well as gas pressure intensifier, the temperature control device and the computer system for registration of the data (load, displacement, temperature). Figure 3.5.14 shows the testing machine with the pipe together with a graph of the loading scenario.

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Figure 3.5.14: Testing machine with the pipe and loading scenario. Two pipes were tested with different load levels. The first pipe was loaded with relatively high loads corresponding to the LCF tests on specimens that failed after 4.686 and 6.500 cycles, respectively. In the first test the internal pressure was 336 bar and the superimposed axial load varied between +/- 365 kN (R= -1) producing in total a maximum tensile stress of 370 MPa and a minimum compressive stress of -218 MPa. After the first cycles the cyclic axial loading was switched from load to displacement controlled loading. In the second test, both the internal pressure and the axial load were reduced by about 35% to obtain at least a 10 times higher lifetime. The first tube failed after 5118 cycles which is in the range of the LCF test on the specimens. The failure has occurred by circumferential cracks on both sides at the interface between weld and base metal (Figure 3.5.15). Fractographic investigations reveal a pattern of circular load markers characteristic for a fatigue crack, which initiated at the outer surface. In all cases, the fatigue cracks were located in the intercritically reheated affected zone about 3 mm from the fusion line of the weld. The crack position related to the macroscopic visible deformation of the tube also bulging indicates that the high deformations induced by bulging has influenced the damage process.

Figure 3.5.15: Overview of tube 1 after testing

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The second tube loaded by an internal pressure of 238 MPa and a cyclic axial load of +/-238 kN failed after 111.903 cycles by loss of internal pressure. In this case no macroscopically visible deformation of the tube walls was observed. Small surface cracks of 0.5 mm and 1.5 mm length were observed in the weld metal at about 2 mm distance. Fractographic investigations revealed a half circular shaped crack in the weld of the tube. Detailed scanning electron microscope investigations show spherical inclusions containing oxygen, mineral elements and some of the metallic elements of the weld metal (Figure 3.5.16). The shape and the chemical composition of these defects indicate that they are so-called fish eyes induced by welding. Mineral elements originate from the welding electrode wrap. The crack propagation patterns clearly show crack initiation at the respective defects and crack propagation in direction of the outer and the inner surface. Also, the aspect in the surrounding of the fish eyes including dendritic structure indicate that cracking first occurred during welding.

Figure 3.5.16: Secondary electrons scanning electron microscope pictures of one weld metal

crack fracture surface Discussion The analytical and test work undertaken within this work package has been devised to provide some comment on potential issues associated with operating a P23 pipework system, subject to realistic operational conditions and working practices. Opportunity has been taken to compare, where relevant, operational experiences accumulated over many years on CMV pipework systems on UK plant. This approach has utilised operational experience and data from E.ON UK, in particular with respect to incidences of girth weld thermal bore cracking on plant, and analytical, modelling and component testing expertise from partners VTT and IWM. The benefits of using operational experience to inform the testing and analytical work has provide a means of assessing the potential performance of P23 pipework systems against the known performance of equivalent CMV pipework systems. The analytical work on pipework modelling and defect tolerance under ‘service’ conditions has shown that an equivalent P23 pipework system is potentially more durable than an equivalent CMV system, providing a number of caveats are recognised. The component tests have been performed successfully and the material model gives a good prediction of lifetime, within the range of specimens tested. Prediction is conservative for smaller loading, i.e. higher lifetime. Despite the defect in the weld the lifetime is higher than expected from a Woehler-curve with the slope of 5 normally used for the assessment of un-welded components (FKM guideline).

Outer tube surface

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Conclusions of Work Package 5 Subjected to modest plant thermal transients, associated with a station that is being operated in accordance with good practice, a P23 pipework girth weld with a circumferential crack has a higher factor of safety than an equivalent CMV pipework girth weld. When subject to severe plant thermal transients, associated with poor operational practice, there is no significant benefit associated with a P23 pipework system in preference to a CMV pipework system. P23 pipework systems are at least as tolerant of mal-adjustment and poor maintenance practices associated with pipework hanger systems as CMV systems. Thermal transients due to poor operation have been shown to persist along the whole length of a typical pipework system. This influences the pipework weld inspection strategy, since the assumption that weld locations adjacent to the Boiler Stop Valve are most at risk is not true for every station or pipework system. This finding supports inspection data found on some UK stations. The material model developed gives a good prediction for lifetime, within the range of specimens tested. This work has shown that a P23 pipework system will be more durable than an equivalent CMV system providing that:

• Good operational practice in maintained, thereby minimising the risk of severe operational transients;

• Pipe support maintenance and adjustment is undertaken at regular intervals; • Pipe weld integrity will be enhanced if the above two points are adhered to; • Inspection regimes, for the detection and management of thermal bore cracks, should

consider all weld locations along the pipeline.

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4 Conclusions The main conclusions on the base of the activities done and the obtained results can be summarised as follow: - The Grades 23 and 24 can be produced with the conventional industrial process routes in

tubular and plate form. For the thick components the quenching after austenitization is necessary to avoid the formation of ferrites island that have a negative effect on the mechanical and creep properties;

- Welding procedures have been developed and optimised for manufacture of thick section P23 components. An extensive range of weld metal compositions have been investigated and a comprehensive series of full scale test specimens have been produced.

- The overmatching, highly alloyed “state-of-the-art” P23 weld metal B323B, selected for its high creep strength and used for the manufacture of long term high temperature test specimens, proved to be a poor choice. This creep-brittle weld metal was found to be liable to reheat cracking during PWHT of full scale thick section butt welds. The occasional presence of pre-existing weld reheat defects in the test specimens therefore led to a pattern of inconsistent and unreliable behaviour. B323B was deliberately chosen as a test of the viability of the high-strength formulation. The negative result is valuable, in that it clarifies the pitfalls that can occur and indicates a compositional range to be avoided.

- Parallel work using Gleeble weld thermal simulation together with the BWI tensile reheat cracking test showed that the P23 HAZ could be susceptible to reheat cracking, while confirming that B323B weld metal is extremely susceptible.

- Later trials showed that alternative electrode compositions, involving slight deviations from the P23 specification, could be developed to produce substantially less creep-brittle weld metals. These successfully survived PWHT without reheat cracking.

- The results indicate that the Monkman-Grant creep ductility parameter M provides a reasonable guide to P23 weld metal performance, and that M values above 1% should generally be adequate. Improved weld metals, with M values of the order of 2%, can now be formulated. These should be subjected to long term testing in future work.

- The tested P23 weld metals nevertheless show intrinsically poorer creep ductility than P23 parent materials. If adequate ductility is to be achieved, some sacrifice in terms of weld metal creep strength appears to be unavoidable. The microstructural reasons for this are not yet clear.

- The risks of reheat cracking in P23 weld metals and heat-affected zones can thus be reduced by careful selection of consumables and welding procedures.

- A comprehensive high temperature materials database has been generated on P23 steels and weldments. These data provide support for high temperature power plant design with P23 steel. Background data on factors such as weld performance and defect development under plant cycling conditions will also support the lifetime management of operational power plant. The results show how important it is to have a good creep ductility for the weld metal of the newer high strength low alloy steels like P23 and P24, in addition to promising creep strength in preliminary qualification. The reasons are related to mechanisms that lead to strain localisation and cracking towards longer term creep exposure, and emphasise the importance of ductility criteria in short term testing and initial selection of candidate compositions. The results have also established an excellent basis for further development of suitable compositions that would perform better, and fundamental understanding on the mechanisms of the creep behaviour of multibead weld metal for high strength low alloy steels like P23 and P24. Further development of the weld metal is clearly needed, and there are new avenues to explore for the purpose. For this class of steels, the project has provided new tools and understanding for faster evaluation and further improvement of high performance welds.

- The assessment of the microstructural evolution of both grades 23 and 24 has been performed by aging tests carried out in the range of temperatures 550-675°C up to 10.000 hours. For both grades, the results obtained in term of microstructural observation seem

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quite well in agreement with the phases predicted at equilibrium by the JMatPro thermodynamic tool. However due to differences in precipitation and coarsening kinetics also metastable phases may develop during ageing. In grade 23, all the metastable types of precipitates like MC, M23C6 and M7C3 show a tendency to change their size distributions and to form the M6C carbide, which is the more stable phase at equilibrium. In grade 24, the situation is a little bit different because at equilibrium the M23C6 carbide is also a stable phase with a faster growth then M6C. In addition to changes in carbide structure increasing thermal exposure causes gradual coarsening in the lath structure. Obviously, during the aging times performed, the true equilibrium stage has not been reached. However, the aging tests at 675°C for 10.000 hours may approximately be equated, on a time-temperature parameter basis, to an exposure time of 200 years at 575°C. This is the maximum proposed service temperature for these low Cr grades, due to their oxidation rate.

- Regarding the variation of the mechanical properties during the aging at 550°C up to 10000 hours, the tensile properties at RT (YS and UTS) are practically not affected by exposure. However, the tensile properties measured at 550°C and 600°C shown a reduction of about 200MPa in each case. Different behaviour is observed for the impact values. Grade 23 appears to be much better then Grade 24 after heat treatment and also after aging: the Grade 24 material reaches values lower then 50 Joule with brittle fracture. This could be related to the larger dimensions of the precipitates that are present in the Grade 24 steels after aging.

- Subjected to modest plant thermal transients, associated with a station that is being operated in accordance with good practice, a P23 pipework girth weld with a circumferential crack has a higher factor of safety than an equivalent CMV pipework girth weld. When subject to severe plant thermal transients, associated with poor operational practice, there is no significant benefit associated with a P23 pipework system in preference to a CMV pipework system. P23 pipework systems are at least as tolerant of mal-adjustment and poor maintenance practices associated with pipework hanger systems as CMV systems.

- Thermal transients due to poor operation have been shown to persist along the whole length of a typical pipework system. This influences the pipework weld inspection strategy, since the assumption that weld locations adjacent to the Boiler Stop Valve are most at risk is not true for every station or pipework system. This finding supports inspection data found on some UK stations. The material model developed gives a good prediction for lifetime, within the range of specimens tested.

This work has shown that a P23 pipework system will be more durable than an equivalent CMV system providing that:

• Good operational practice in maintained, thereby minimising the risk of severe operational transients;

• Pipe support maintenance and adjustment is undertaken at regular intervals; • Pipe weld integrity will be enhanced if the above two points are adhered to; • Inspection regimes, for the detection and management of thermal bore cracks, should

consider all weld locations along the pipeline.

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5 Exploitation and impact of the research results a) Actual applications: the activities carried out and the results obtained in the project by the

Consortium will have positive effects on the EU steel industries. Tubular and plate components for new power and petrochemical plants, as well as, for plant repowering have been consistently characterised.

b) Technical and economic potential for the use of the results: Welding consumables, welding procedures and welding new simulations methodology are now more investigated and characterised form the point of view of microstructural and mechanical point of view. A consistent database, including also creep on base, weld metal and welded joint is now available. The first approach to the investigation of the microstructural evolution of these grades has been realised. Obviously the analysed specimens had a limited aging time, but the availability of more aged specimens from crept samples, which are still running in this project, and in other EU programs such as COST 536, will be very useful to obtain more information on the microstructural evolution of both Grade 23 and 24. These investigations will be mandatory to guarantee the creep data assessment validity, that will be carried out in the near future in the ECCC group, as it was done in the past for the 9%Cr grades. In fact up to now the microstructural evolution seem to follow the prediction made by the thermodynamic tools, but information after >50.000 hours are required to verify if unpredicted phases could precipitate or transformed. The new methodology developed for the XRD analysis of the extracted precipitates will be applied in the future to analyse other ferritic/martensitic creep resistant steels and will be also calibrated for the investigation of the austenitic steels that will be more extensively used in the next USC power plants.

c) Any possible patent filing: the grade 23 is free from patent. The grade 24 is patented in Europe by Vallourec & Mannesmann. Potentially the consumables could be patented. The Consortium will evaluate this opportunity.

d) Publications/conference presentations resulting from the project: the creep data on base materials and welded joints will be shared with the ECCC (European Creep Collaborative Committee) founded in 1991 to collect, generate and assess EU creep data to supply the creep values to CEN to be introduced into EN standards. Publications have been scheduled for the next “Creep 2008” Int. Conference to be held in Bayreuth (D) in May 2008. Other publication will be defined later.

e) Any other aspects concerning the dissemination of results: The approach taken to develop and validate the material model should be considered for future projects, especially if more representative test conditions and durations can be accommodated. This would require further refinement and idealisation of load conditions, both mean, cyclic range, sequence and duration. Provision of site service data and experience should be considered to support similar projects in the future, since it allows comparison against both an experience and analytical base.

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List of figures and tables Figure 3.1.1: Grade 23 pipes (OD219 x WT31.75mm) and plate (1000 x 400 x 50mm) Figure 3.1.2: Examples of microstructure for grade 23 and 24 Figure 3.1.3: T24 code C, external surface Figure 3.1.4: P23 cast G (OD 219mm x WT 31.75mm): SEM microstructure Figure 3.1.5: Quenched and tempered casts G and C (TEM on Thin Foil) Figure 3.1.6: precipitates (TEM extraction replica) on P23 Cast G quenched and tempered Figure 3.1.7: size distribution of precipitates Figure 3.1.8: Microstructure of plate in different positions Figure 3.1.9: Hot tensile data on P23 pipe casts G and J Figure 3.1.10: comparison of the mechanical properties of the Pipe A and the Plate P Figure 3.1.11: FATT curve for pipe Code G Figure 3.1.12: Charpy transition curves - transverse (a) and longitudinal orientation (b) of plates Figure 3.1.13: mastercurve Figure 3.1.14: Examples of broken specimens Figure 3.2.1: Metrode weld deposit sample geometry Figure 3.2.2: Reheat cracking in P23 weld metal subsequently reheated into the coarse-grained-temperature region (Tp = 1340 °C): CG(R)-WM – filler metal B323B. Figure 3.2.3: Reheat cracking in P23 weld metal discovered from a CT specimen extracted from a laboratory-scale welded sample (PWHT) – B323B filler metal Figure 3.2.4: MMA weld geometry. Figure 3.2.5: Pipe G welded joint Figure 3.2.6: Macro section and hardness profile of pipe H weld after PWHT Figure 3.2.7: A – weld metal and coarse HAZ; B – coarse and fine HAZ; C – base metal Figure 3.2.8: All-weld-metal creep test data: Creep rate plotted against ductility Figure 3.2.9: All-weld-metal creep test data. Creep life plotted against ductility. Figure 3.2.10: pipe G welded joint: specimen sampling Figure 3.2.11: Test plate dimensions Figure 3.2.12: The double V weld preparation Figure 3.2.13: Submerged arc welding Figure 3.2.14: Image of the liquid-penetrant cracking observation Figure 3.2.15: Image of the liquid-penetrant final faces without defects Figure 3.2.16: Plate cross weld section Figure 3.2.17: Hardness profile of plate welded joint Figure 3.2.18: Charpy V results Figure 3.3.26: Distribution of BM tests according to stress and temperature Figure 3.3.27: Distribution of tests according to material heat and testing time Figure 3.28.3: Distribution of CW tests according to stress and temperature. Figure 3.3.29: Creep strength model (minimum commitment = MC) for P23 base material based on the Aloas steels A, B, G, H and P. Figure 3.3.30. The Aloas P23 base material creep strength model (minimum commitment) presented as a Larson-Miller (PLM) plot. The red line presents the 600°C prediction line (MC). Figure 3.3.31: T/P24 Creep strength comparison of material C (diamonds) and E (boxes) data against the VdTÜV 533 (continuous black line) and V&M data (open circles). The dash-dotted line is the SM model line for 575°C and the dotted line represents the VdTÜV -20% in stress. Figure 3.3.32: Welded H Figure 3.3.33: Welded P Figure 3.3.9: Welded G Figure 3.3.34: Weld metal creep results; B323 (circles), B323 (triangles) and B322 (squares) in comparison to MB (continuous lines) Figure 3.3.35: Short term creep results from weld metal development (WP2) Figure 3.3.36: Distribution of data in the creep strain assessment according to accumulated strain. Figure 3.3.37: Creep strain simulation with LCSP (red curve) and corresponding Norton minimum strain rate prediction (linear blue line) for P23 at 575°C / 91 MPa (rupture predicted at 100 000 h ). Figure 3.3.38: Actual creep strain data, time to specified strain (here 1 and 2%) compared to predicted time to strain. Note that the curve form of material J is not very well represented by the model. Figure 3.3.39: Measured vs predicted time to specified strain for B323B WM at 660°C / 80 MPa.

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Figure 3.3.40: Simulated creep curves in time and normalised time (t/tr) for HAZ (weakest), BM and WM (strongest) used in P23 FEA. Figure 3.3.41: FE-mesh for Comsol girth weld simulation. Figure 3.3.42: axisymmetric Λ-filtered creep exhaustion results for the P23 simulation at 100% of BM life Figure 3.3.43: LCF data – Showing cycles to crack initiation versus total strain range. Figure 3.3.44: LCF data – Showing cycles to crack initiation versus plastic strain range. Figure 3.3.45: LCF data – Cyclic stress range as a function of plastic strain range Figure 3.3.46: Fatigue crack growth data Figure 3.3.47: Micrographs showing the crack path following the weld bead shape in specimens HAZ1 and HAZ4. Figure 3.3.48: The C* plot of all four CCG tests Figure 3.3.25: Logarithmic stress – life plot of 650°C creep rupture data. Figure 3.3.49: LCF data – Showing strong correlation for all parent material casts between nominal “hysteresis energy” and cycles to initiation. Figure 3.3.50: LCF data – Nominal total hysteresis energy as a function of elastic strain range. Figure 3.4.1: Hardness values as function of time and temperature of aging for grade 23 and 24 Figure 3.4.2: SEM microstructure of grade 23 after aging at 550°C Figure 3.4.3: SEM microstructure of grade 23 after aging at 600°C Figure 3.4.4: EDS analysis of precipitates Figures 3.4.5: Aspect of the microstructure of Grade 24 as function of time and temperature of aging Figure 3.4.6: equilibrium diagram by JMatPro Figure 3.4.7: Frequency of precipitates vs size in Grade 23 after heat treatment Figure 3.4.8: Frequency of precipitates vs size in Grade 23 after 6425 hours at 600°C Figure 3.4.9: Frequency of precipitates versus size in Grade 23 after 15.000hours at 600°C Figure 3.4.10: Frequency of precipitates versus size in Grade 23 after 1130 hours at 660°C Figure 3.4.11: Evolution of the composition of the precipitates for Grade 23 from as treated material up to 10.000 hours aged at 675°C. Figure 3.4.12: examples of extraction replicas from aged grade 23 specimens Figure 3.4.13: evolution of the diameter of the particles as function of PLM (time and temperature of aging): the values at PLM 20 are the as treated dimensions Figure 3.4.14: Evolution of the composition of the precipitates for Grade 24 from as treated material up to 10.000 hours aged at 675°C. Figure 3.4.15: examples of extraction replicas from aged grade 24 specimens Figure 3.4.16: Microstructural evolution during simulated high temperature service for material G, shown as a series with increasing PLM. Figure 3.4.17: Precipitation coarsening at the grain boundaries for P23 (material G), as received versus 10.000 h aged at 675°C Figure 3.4.18: comparison of X-RD particle analysis of grade 23 specimens as treated and aged at 550°C for 10.000 hours Figure 3.4.19: comparison of X-RD particle analysis of grade 23 specimens as treated and aged at 600°C for 10.000 hours Figure 3.4.20: T23 and T24 tensile data at 20°C after aging at 550°C up to 10.000hours Figure 3.4.21: T23 tensile data at 600°C after aging at 550°C up to 10.000hours Figure 3.4.22: T23 tensile data at 600°C after aging at 600°C up to 10.000hours Figure 3.4.23: Comparison of impact values of grade 23 and 24 specimens aged at 550 and 600°C up to 10.000 hours Figure 3.4.24: Comparison of brittle fracture of grade 23 and 24 specimens aged at 550 and 600°C up to 10.000 hours Figure 3.5.1: Typical Main Steam Line Temperature Transient Figure 3.5.2: P23 versus CMV - Circumferential defect tolerance to typical plant transients, based on a factor of safety Figure 3.5.3: The model geometry with hanger forces on the left. Part of the finite element mesh on the right. Figure 3.5.4: Displacement magnitude [m] at the end of the steam transient of CMV (left) and P23 (right) models. The displacements are scaled up by a factor of 10 in the deformed shape of the model for illustrational reasons. Figure 3.5.5: Distribution of von Mises stress [Pa] on the inner surface of CMV model (left) and P23 model (right) at the end of the steam transient. The location is from the boiler end.

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Figure 3.5.6: Von Mises stresses in the outer wall in element 8872 Figure 3.5.7: Vertical displacements of the node 11963. Figure 3.5.8: APROS steam line model Figure 3.5.9: Steam temperatures in the pipe (P23) Figure 3.5.10: Steam Pressures (P23) Figure 3.5.11: Pipe outlet temperatures with CMV and P23 geometries Figure 3.5.12: Applied strain (A) and resulting stress at 550, 575 and 600°C (B) Figure 3.5.13: Experimental and modelled cycles to failure of the specimen test on material H with a scatter band of factor two. Figure 3.5.14: Testing machine with the pipe and loading scenario. Figure 3.5.15: Overview of tube 1 after testing Figure 3.5.16: Secondary electrons scanning electron microscope pictures of one weld metal crack fracture surface List of tables Table 3.1.1: ASTM Standards and chemical analysis of Grade 23 and 24 heats Table 3.1.2: ASTM Standards for heat treatment Table 3.1.3: product forms and heat treatments Table 3.1.4: example of the results obtained Table 3.1.5: minimum values for ASTM standards Table 3.1.6: Examples of mechanical properties obtained on different products and grades Table 3.1.7: mechanical test results as function of cooling rate and tempering conditions Table 3.1.8: Charpy V impact results on plate P with different cooling rate and tempering conditions Table 3.1.9: Summary of the impact toughness results Table 3.1.10: To values Table 3.2.1: All weld metal analysis and aim compositions, wt %. Table 3.2.2: Ambient temperature all weld metal tensile, hardness and impact properties Table 3.2.3: Reheat cracking test results for the simulated HAZs of P23 steel G, P22 reference steel, and the thermally Gleeble re-heated microstructures of the B323B multipass weld metal. Table 3.2.4 – Welding samples produced during the project. Table 3.2.5: Cross weld tensile test results Table 3.3.12. ALoAS WP3 test, materials and main observations in the order of task. BM = base material, WM = weld material (consumable) and CW = cross weld. Table 3.3.13. Calculated weld strength factors for welded P, G and H. Note that H has low ductility fractures in some test. For material H no valid WSF can be calculated Table 3.3.14. Optimised values of log(tp) and R for welded P23 steels P and G Table 3.3.15 Calculated weld strength factors for welded P and G using RPC. The values are predictions for 575°C and 100 000 h. The efficiency factor is calculated on the data. Table 3.3.16. Creep strain shape equations for Aloas P23. Note: Stress in MPa , temperature in °C. Table 3.3.17. Calculated rupture times for single zones in weldment: Girth weld in a pipe with 300 mm inner radius and 28 mm wall thickness with internal pressure of 139 bars (nominal stress = 70 MPa) at 595°C. 3-zone case; base material (BM), heat affected zone (HAZ) and weld metal (WM). Table 3.3.18. LCF continuous cycle test conditions and results Table 3.3.19. LCF dwell test conditions and results Table 3.3.20. fatigue crack growth data Table 3.3.21. CCG test details for HAZ and WM CT specimens. Table 3.3 22. Summary of the metallograpic investigation of the ALOAS P23 CCG tests Table 3.4.1. Aging conditions for grades 23 and 24. Table 3.4.2: aged specimens analysed by TEM Table 3.4.3: evolution of dimensions for grade 24 (dimensions nm) Table 3.4.4. Ageing test matrix for materials G, C and M Table 3.4.5. Specimen/micrograph specific codes for the ageing tests in Table 4.4. The table is in the order of ascending Larson-Miller time temperature parameter PLM (C=20, PLM/1000), showing increasing “age” (=simulated time in years) in power plant service conditions (575°C). Table 3.5.1. Peak Axial Stress at Pipe Bore Table 3.5.2: APROS model parameters

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Allen, D. and Siegele, D. ’Development of weldments of Grade 23 steel for new generation power plants’ Proc.Conf. ”11th International Conference on Creep and Fracture of Engineering Materials and Structures (Creep 2008)”, Bayreuth, 4-9 May 2008. Germany (accepted for publication)

[2] Nevasmaa, P., Salonen, J., Holmström, S. and Caminada, S. IIW-Doc. IX-2237-07 / VTT Research Report No VTT-R-02973-07. VTT, Espoo 2007.

[3] Dhooge, A. & Vekeman, J. Welding in the World 49(2005)9/10. pp. 75-93. [4] Nawrocki, J. G., DuPont, J. N., Robino, C. V. and Marder, A. R. Welding Journal 79(12), (2000),

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[7] EN 288-3 – Specification and approval of welding procedures for metallic materials – Part 3: Welding procedure tests for arc welding of steels.

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[10] DATA, VdTÜV 533 (12.2003) Warmfester Stahl 7 CrMoVTiB 10-10, Werkstoff-Nr1.7378 [11] Holmström, Stefan; Auerkari, Pertti, Robust prediction of full creep curves from minimal data and

time to rupture model, Energy Materials. Institute of Materials, Minerals and Mining and W.S.Maney & Son Ltd. Vol. 1 (2006) No: 4, 249-255, doi-link:10.1179/174892406X173594

[12] Hurst R.C, Rantala J.H. Influence of multiaxial stresses on creep and creep rupture of tubular components, ASM Handbook, Vol. 8, 2000 p.

[13] Holmström, S., Laukkanen A., Calonius, K., Visualising creep exhaustion in a P22 girth weld. International Conference on Life Management and Maintenance for Power Plants. Helsinki-Stockholm-Helsinki, 12-14 June 2007. Vol 2. pp.208-221

[14] KOMAI, N, and IMAZATO, T, (2006), “Effect of Tempering Times on Creep Strength in ASME Gr.23 (2.25Cr – 1.6W) Steel”, Proc. 8th Liege Conf. “Materials for Advanced Power Engineering 2006”, publ. Forschungszentrum Julich, Book 2, pp. 997-1009.

[15] Bendick, B., Gabrel J., Hahn B., Vandenberghe B., New alloy heat resistant ferritic steels T/P23 and T/P24 for power plant application. Pressure Vessels and Piping 84 (2007) 13-20. doi:101016/j.ijpvp.2006.09.002.

[16] CARBON ABATEMENT TECHNOLOGIES PROGRAMME, Advanced modelling and testing: thick section welded alloy HCM2S (P23). Project report R293 for the cleaner fossil fuels programme. (http://www.berr.gov.uk/energy/sources/renewables/publications/page29152.html)

[17] PD 6605, 1998. Guidance on methodology for assessment of stress rupture data, Part 1 and 2, BSI, London 51 + 27 p.

[18] Holmström S., Auerkari P. Predicting weld strength reduction for 9Cr steel. Pressure Vessel and Piping Journal: 10.1016/j.ijpvp.2006.08.007

[19] ECCC recommendations, 2005. Volume 5 Part IIb [Issue 1]. Guidance for the assessment of creep rupture, creep strain and stress relaxation data, Recommendations for the assessment of weld creep rupture data.

[20] Holmström, S., Auerkari, P., Holdsworth S. Predicting creep strain response from rupture data and robust creep curve model. International Conference on Life Management and Maintenance for Power Plants. Helsinki-Stockholm-Helsinki, 12-14 June 2007. Vol 1. pp.185-195.

[21] S.R.Holdsworth, M.Askins, A.Baker, E.Gariboldi, S.Holmström, A.Klenk, M.Ringel, G.Merckling, R. Sandström, M.Schwienheer, S.Spigarelli. 2005, Factors influencing creep model equation selection. Creep and Fracture in High temperature Components, Design and life assessment issues. ECCC creep conference, 12-14 Sept. London, UK. p.380, ISBN No. 1-932078-49-5

[22] Holmström, S., Laukkanen A., Calonius, K., Visualising creep exhaustion in a P22 girth weld. International Conference on Life Management and Maintenance for Power Plants. Helsinki-Stockholm-Helsinki, 12-14 June 2007. Vol 2. pp.208-221.

116

[23] Holmström S, Laukkanen A, Calonius K, Weldment matching for creep life of P22 and P91 girth welds. International Conference on Integrity of High Temperature Welds, 24-26 April 2007, London, UK. Conf. Proc. IOM Communications publication, ISBN-1-86125-166-1. p. 245-255.

[24] TERANISHI, H, and MC EVILY, A J, (1979), “A Comparison of the Elevated Temperature Low-Cycle Fatigue Behaviour of 2¼ Cr-Mo-V Steel”, Proc. Intl. Conf on Low Cycle Fatigue Strength, 1979, DVM, Stuttgart, pp. 25-38.

[25] Fenton S F., Morris A., Mulvihill P., ‘Investigation and analysis of CrMoV Bore Cracking’, Institute of Materials 2nd International Conference “Integrity of High Temperature Welds”, November 2003. London

[26] Morris A., ‘Application of FE Analysis to the Ongoing Management of Main Steam Pipework Bore Cracking’, NAFEMS World Congress, May 2005. Malta

[27] Biot M A., ‘New Methods in Heat Flow Analysis with Application to Flight Structures’, J.Aeronaut. Sci.,24 (1967), 857-873

[28] ARNDT, J, et al. (1998), “The T23 / T24 Book”, Vallourec and Mannesmann Tubes [29] R6 Assessment Code, V4.3, British Energy Generation Ltd. [30] PSA5 code – Whessoe Computing Systems [31] Abaqus finite element software, Simulia, Dassault Systems [32] Riedel, H. (1987), Fracture at High Temperatures, Springer Verlag, Berlin, Heidelberg, New York [33] Schmitt, W.; Mohrmann, R.; Riedel, H.; Dietsche, A.; Fischersworring-Bunk, A.: Modelling of

fatigue life of automobile exhaust components. Fatigue 2002 (Ed. A.F. Blom), Engineering Materials Advisory Services Ltd., Cradley Heath, U.K., 2002, pp. 781-788.

117

Appendix 1:

Welding procedure specification (WPS)

118

Weld Procedure Record ALoAS P23 Ref: ALoAS/B322(AW)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B322)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: As-welded

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B322 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.056 0.55 0.29 0.011 0.012 2.38 0.03 0.12 0.24 0.05 1.74 0.05 <0.001 0.004 0.02 12

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa -- 0.2% proof MPa -- Elongation, 4d % --

Elongation, 5d % -- Reduction of area % --

Hardness, HV (10) Weld Metal

Cap Max = 342 (average = 312)

Mid thickness Max = 336 (average = 322)

Orig. GBH Date Nov 2004

119

Weld Procedure Record ALoAS P23 Ref: ALoAS/B322(715)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B322)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 715°C/2 hours

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B322 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.056 0.55 0.29 0.011 0.012 2.38 0.03 0.12 0.24 0.05 1.74 0.05 <0.001 0.004 0.02 12

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa 721 59 0.95 0.2% proof MPa 639 35 0.63 Elongation, 4d % 20.5

Weld C/L VWT0/1.25

55 0.82

This would be the bead overlap region in a 2 bead

per layer weld. Elongation, 5d % 17 52 0.73 Reduction of area % 61 117 1.55

Weld offset C/L by 4mmVWT4/1.25 54 0.68

This would be the bead centre of a 2 bead per

layer weld.

Hardness, HV (10) Weld Metal

Cap Max = 281 (average = 273)

Mid thickness Max = 245 (average = 229)

Orig. GBH Date Nov 2004

120

Weld Procedure Record ALoAS P23 Ref: ALoAS/B322(740)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B322)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 740°C/2 hours

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B322 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.056 0.55 0.29 0.011 0.012 2.38 0.03 0.12 0.24 0.05 1.74 0.05 <0.001 0.004 0.02 12

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa 666 140 1.83 0.2% proof MPa 581 111 1.49 Elongation, 4d % 22.5

Weld C/L VWT0/1.25

118 1.73

This would be the bead overlap region in a 2 bead

per layer weld. Elongation, 5d % 20 160 2.07 Reduction of area % 66 132 1.87

Weld offset C/L by 4mmVWT4/1.25 126 1.70

This would be the bead centre of a 2 bead per

layer weld.

Hardness, HV (10) Weld Metal

Cap Max = 243 (average = 238)

Mid thickness Max = 249 (average = 241)

Orig. GBH Date Nov 2004

121

Weld Procedure Record ALoAS P23 Ref: ALoAS/B323(AW)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B323)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: As-welded

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B323 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.078 0.59 0.39 0.008 0.010 2.36 0.03 0.12 0.30 0.09 1.75 0.04 <0.001 0.025 0.03 48

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa -- 0.2% proof MPa -- Elongation, 4d % --

Elongation, 5d % -- Reduction of area % --

Hardness, HV (10) Weld Metal

Cap Max = 397 (average = 372)

Mid thickness Max = 376 (average = 350)

Orig. GBH Date Nov 2004

122

Weld Procedure Record ALoAS P23 Ref: ALoAS/B323(715)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B323)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 715°C / 2 hours

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B323 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.078 0.59 0.39 0.008 0.010 2.36 0.03 0.12 0.30 0.09 1.75 0.04 <0.001 0.025 0.03 48

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa 779 18 0.31 0.2% proof MPa 711 18 0.34 Elongation, 4d % 10

Weld C/L VWT0/1.25

24 0.44

This would be the bead overlap region in a 2 bead

per layer weld. Elongation, 5d % 8.5 16 0.58 Reduction of area % 15 46 0.48

Weld offset C/L by 4mm VWT4/1.25 17 0.30

This would be the bead centre of a 2 bead per

layer weld.

Hardness, HV (10) Weld Metal

Cap Max = 325 (average = 313)

Mid thickness Max = 283 (average = 281)

Orig. GBH Date Nov 2004

123

Weld Procedure Record ALoAS P23 Ref: ALoAS/B323(740)

Material: CMn base material Weld Details

Filler Metal: ALoAS P23 (batch B323)

Classification: --

Process: MMA Flux: NA

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 740°C / 2 hours

10-12mm

20°

12.5mm

CMn

Buttering

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min

Heat Input

kJ/mm Procedural Comments:

1-12 3.2 115-125 21-25 ~120-180 0.9-1.6 Buttered with two layers of B323 electrode.

Bead sequence was 6 layers of 2 runs each, 12

runs in total.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al Ti N B* Deposit 0.078 0.59 0.39 0.008 0.010 2.36 0.03 0.12 0.30 0.09 1.75 0.04 <0.001 0.025 0.03 48

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS MPa 723 31 0.50 0.2% proof MPa 657 61 0.85 Elongation, 4d % 9

Weld C/L VWT0/1.25

103 1.43

This would be the bead overlap region in a 2 bead

per layer weld. Elongation, 5d % 7.5 38 0.62 Reduction of area % 15 46 0.68

Weld offset C/L by 4mm VWT4/1.25 24 0.45

This would be the bead centre of a 2 bead per

layer weld.

Hardness, HV (10) Weld Metal

Cap Max = 276 (average = 267)

Mid thickness Max = 264 (average = 258)

Orig. GBH Date Nov 2004

124

Weld Procedure Record P23 Sub-Arc Weld Ref: SAW/P23/1

Material: CMn base material Weld Details

Filler Metal: Metrode 2CrWV (WO21521)

Classification: --

Process: SAW Flux: LA491 (WO19632)

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 715°C / 2 hours

28mm

20°

19mm

CMn

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min ROL mm

Heat Input

kJ/mm Procedural Comments:

1-26 2.4 300 30 350 - ~1.5 Plate was not buttered.

Bead sequence was 2 layers of 3 runs each, and

then 5 layers of 4 runs each. Specimens taken

from centre where dilution was minimal.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al N B* Wire 0.072 0.44 0.28 0.002 0.012 2.13 0.13 0.07 0.23 0.06 1.64 0.12 0.01 0.01 15

Deposit 0.058 0.51 0.36 0.002 0.015 2.09 0.12 0.07 0.19 0.05 1.49 0.11 0.01 0.01 8

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS 695MPa 36 0.54 0.2% proof 630MPa Weld 50 0.68 Elongation, 4d 16% 20 0.21 Elongation, 5d 12% Reduction of area 38%

Hardness, HV (10) Weld Metal

Cap 216-270 (average = 250)

Mid thickness 235-281 (average = 260)

Orig. GBH Date Nov 2004

125

Weld Procedure Record P23 Sub-Arc Weld Ref: SAW/P23/2

Material: CMn base material Weld Details

Filler Metal: Metrode 2CrWV (WO21521)

Classification: --

Process: SAW Flux: LA491 (WO19632)

Current: DC+ Position: ASME 1G

Preheat / Interpass Temperature: 200 / 250°C

PWHT: 740°C / 2 hours

28mm

20°

19mm

CMn

Run No

ø mm

Current Amp

Arc Volts

Travel Speed

mm/min ROL mm

Heat Input

kJ/mm Procedural Comments:

1-26 2.4 300 30 350 - ~1.5 Plate was not buttered.

Bead sequence was 2 layers of 3 runs each, and

then 5 layers of 4 runs each. Specimens taken

from centre where dilution was minimal.

Analysis C Mn Si S P Cr Ni Mo V Nb W Cu Al N B* Wire 0.072 0.44 0.28 0.002 0.012 2.13 0.13 0.07 0.23 0.06 1.64 0.12 0.01 0.01 15

Deposit 0.058 0.51 0.36 0.002 0.015 2.09 0.12 0.07 0.19 0.05 1.49 0.11 0.01 0.01 8

* ppm Charpy +20°C Tensile (all-weld): Impact J mm

UTS 645MPa 48 0.71 0.2% proof 572MPa Weld 218 2.37 Elongation, 4d 22% 259 1.98 Elongation, 5d 17.5% Reduction of area 53%

Hardness, HV (10) Weld Metal

Cap 206-245 (average = 233)

Mid thickness 238-254 (average = 244)

Orig. GBH Date Nov 2004

126

Welding procedure specification (WPS)

Location: ISQ Welding Process: SMAW Joint type: Weld pad on flat plate Welding Position: PA Method of preparation and cleaning: Joint Design: Welding details:

Run Process Size of filler metal

Current A ± 10

Voltage V ± 10

Type of current / Polarity

Travel speed

Heat input

1 SMAW 4.0 190 24 DC/+ 17 21 Others SMAW 5.0 220 26 DC/+ 23 14

Designation of welding consumables: HITEHR Preheat temperature: 220ºC Interpass temperature: 330ºC Post-weld heat treatment: Stress relief Time/Temperature/Method: 6h/730ºC ± 10ºC / Furnace Heating and cooling rates: 300ºC/h max (above 320ºC)

150 mm

150 mm

30 mm

Weld Pad

Welding direction

127

TABLE 1– AMBIENT ALL WELD METAL TENSILE AND HARDNESS PROPERTIES OF THE FIRST SERIE

OF METRODE COMSUMABLES

Elongation %

Hardness, Hv(10Kg)

Cap Mid thickness

PWHT

ºC/h

Rm MPa

Rp0.2%

MPa 4D 5D

Z %

Average Max Average Max 23L AW 870 19 16 50 290 - 350 -

AW - - - - - 312 342 322 336 715/2 721 639 20.5 17 61 273 281 229 245 B322 740/2 666 581 22.5 20 66 238 243 241 249 AW - - - - - 372 397 350 376

715/2 779 711 10 8.5 15 313 325 281 283 B323 740/2 723 657 9 7.5 15 267 276 264 258 AW - - - - - 287 322 287 322

715/2 695 630 16 12 38 250 270 260 281 SAW 740/2 645 572 22 17.5 53 233 245 244 254

TABLE 2– AMBIENT ALL WELD METAL TENSILE AND HARDNESS PROPERTIES (PWHT 705ºC/10HR) OF HITHER ELECTRODES

Elongation %

Hardness, Hv10

Cap Mid thickness

Rm MPa

Rp0.2%

MPa 4D 5D

Z %

Average Max Average Max MH1 672 585 18 16 32 240 343 226 235 MH2 - - - - - 232 240 218 227

Sumitomo 803 750 21 -- -- -- -- -- -- B-Thyssen 620 520 18 -- -- -- -- --

TABLE 3 AMBIENT ALL WELD METAL TENSILE AND HARDNESS PROPERTIES OF THE ELECTRODES

FOR LARGE WELDING

Elongation %

Hardness, Hv(10Kg)

Cap Mid thickness

PWHT

ºC/h

Rm MPa

Rp0.2%

MPa 4D 5D

Z %

Average Max Average Max 740/2 745 668 8.5 8 20 - - 263 272 B323B(3.2)

WO22919 3.2mm 760/2 725 637 15 15 30 - - 241 249

740/2 688 640 9.5 9 17 - - 255 270 B323B(4.0) WO22920 4.0mm 760/2 686 621 10 9 16 - - 240 251

TABLE 4 AMBIENT ALL WELD METAL TENSILE AND HARDNESS PROPERTIES OF THE SECOND OF

METRODE COMSUMABLES

Elongation %

Hardness, Hv(10Kg)

Cap Mid thickness

PWHT

ºC/h

Rm MPa

Rp0.2%

MPa 4D 5D

Z %

Average Max Average Max A71 740/2 725 640 18 16.5 55 250 254 248 254 A72 740/2 732 652 20.5 18 62 246 251 247 251 A84 740/2 704 620 20.5 18 52 239 242 240 242 A85 740/2 748 672 19.5 17.5 66 251 258 249 256

128

TABLE 5– ALL WELD METAL IMPACT PROPERTIES OF THE FIRST SERIE OF METRODE ELECTRODES Impact Energy at +20ºC

Process Batch PWHT ºC/h Location CVN, J

Average (min) LE, mm

Average (min) AW --- 22 0.39 SMAW 23L

715/0.5 --- 70 --- WELD CENTRELINEA 50 (35) 0.80 (0.63) 715/2 BEAD CENTREB 74 (52) 0.99 (0.68) WELD CENTRELINEA 123 (111) 1.68 (1.49)

B322 740/2

BEAD CENTREB 139 (126) 1.88 (1.70) WELD CENTRELINEA 20 (18) 0.36 (0.31) 715/2 BEAD CENTREB 26 (16) 0.45 (0.30) WELD CENTRELINEA 65 (31) 0.93 (0.50)

MMA

B323 740/2

BEAD CENTREB 36 (24) 0.58 (0.45) 715/2 WELD CENTRELINEA 35 (20) 0.48 (0.21) SAW 740/2 WELD CENTRELINEA 175 (48) 1.69 (0.71)

CVN – Charpy energy LE – Lateral Expansion

TABLE 6– ALL WELD METAL IMPACT PROPERTIES (PWHT 705ºC/10HR) OF HITHER ELECTRODES +20ºC 0ºC -20ºC -30ºC -35ºC -40ºC -45ºC -50ºC -60ºC -70ºC

CVN J

LE Mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE Mm

CVN J

LE Mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE mm

MH1 158 165 171

2.23 2.34 2.38

149 151

2.16 2.21

132 145

1.84 2.09

146 1.86 11 0.10

MH2 188 2.60 163 1.84 138 1.93 143 149 2.13 21 0.43 14 0.30 135 1.86 6 0.14 B-Thyssen

120

CVN – Charpy energy LE – Lateral Expansion TABLE 7– ALL WELD METAL IMPACT PROPERTIES (PWHT 705ºC/10HR + STEP COOLING) OF HITHER

ELECTRODES

+20ºC 0ºC -20ºC -30ºC -35ºC -40ºC -45ºC -50ºC -60ºC -70ºC CVN

J LE mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE mm

CVN J

LE Mm

CVN J

LE Mm

CVN J

LE mm

CVN J

LE Mm

CVN J

LE mm

CVN J

LE mm

MH2 170 2.30 164 2.06 61 0.98 51 0.72 146 2.06 177 2.35 11 0.20 9 0.12

CVN – Charpy energy LE – Lateral Expansion

TABLE 8 ALL WELD METAL IMPACT PROPERTIES OF THE ELECTRODES FOR LARGE WELDING

IMPACT ENERGY AT +20°C

BATCH PWHT °C/H LOCATION (1) CVN, J

AVERAGE (MIN)

LE, MM AVERAGE

(MIN)

740/2 WELD CENTRELINE A 70 (60) 0.93 (0.85) B323B(3.2)

WO22919 3.2mm 740/2 WELD CENTRELINE

A 99 (70) 1.25 (0.99)

740/2 WELD CENTRELINE A 61 (51) 0.79 (0.69) B323B(4.0)

WO22920 4.0mm 740/2 WELD CENTRELINE

A 91 (45) 1.24 (0.74)

CVN = Charpy energy LE = Lateral Expansion

TABLE 8 ALL WELD METAL IMPACT PROPERTIES OF THE SECOND SERIE OF METRODE ELECTRODES

129

IMPACT ENERGY AT +20°C IMPACT ENERGY AT -20°C

BATCH PWHT °C/H LOCATION (1) CVN, J

AVERAGE (MIN)

LE, MM AVERAGE

(MIN)

CVN, J AVERAGE (MIN)

LE, MM AVERAGE (MIN)

A71 740/2 WELD CENTRELINE A 94 (80) 1.39 (1.25) 32 (26) 0.48 (0.42)

A72 740/2 WELD CENTRELINE A 127 (112) 1.85 (1.66) 66 (57) 0.96 (0.87)

A84 740/2 WELD CENTRELINE A 136 (130) 1.88 (1.78) 95 (74) 1.34 (1.07)

A85 740/2 WELD CENTRELINE A 125 (119) 1.74 (1.68) 67 (53) 0.96 (0.84)

CVN = Charpy energy LE = Lateral Expansion

TABLE 9 – AMBIENT ALL WELD METAL TENSILE AND IMPACT PROPERTIES OF SUMITOMO ELECTRODES

Elongation %

Charpy Impact value J PWHT

ºC/h

Rm MPa

Rp0.2%

MPa 4D 5D

Z % Nº1 Nº2 Nº3 Average

Sumitomo P23 715/1 803 715 21 -- -- 71 73 71 72

A

B

Figure 1 – Localization of the impact tests.

130

Appendix 2: LM and SEM microstructural investigation of aged specimens Steel G (P23) Opt. 1000x SEM 2000x

625ºC 675ºC

100 h

300 h

1000 h

3000 h

10 000 h

131

Opt. 1000x SEM 2000x

625ºC Steel C (T24) Steel M (P24)

Steel M (P24) Coarse grained HAZ Weld metal

Coarse grained HAZ (SEM)

675ºC Steel C (T24) Steel M (P24)

Steel M (P24) Coarse grained HAZ Weld metal

Coarse grained HAZ (SEM)

100 h

300 h

1000 h

132

3000 h

10 000 h

SEI, 2000x BEI, 2000x Grade 23 Pipe G 675ºC 3000 h

Grade 23 Pipe G 675ºC 10.000 h

Grade 24 Pipe M HAZ 675ºC 1000 h

133

Technical Annex Proposal No: PR-02032

Contract No:

TITLE: APPLICATION OF ADVANCED LOW ALLOY STEELS FOR NEW HIGH

TEMPERATURE COMPONENTS - (ALoAS)

1. OBJECTIVES

The long term performance of similar and dissimilar welded joints in the creep resistant steels is often life-limiting in the design and operation of high temperature power and process plant. For the new low alloy steels, long term creep data on welded joints are not currently available, and the microstructural evolution of welds and base material is not well understood. It is necessary to investigate these factors to develop design data, welding consumables and welding procedures to minimise the risks of plant service failures at similar and dissimilar welds.

Two specific development requirements are addressed in the proposed project. First, problems of poor weld creep ductility have been reported. Welding consumables with proven high temperature properties therefore need to be developed and validated by creep rupture testing, microstructural and mechanical properties assessment. Secondly, the new steels and welds must be shown to have improved resistance to the in-service “bore cracking” phenomenon now identified in existing UK CrMoV steam pipework. This will require comparative low cycle fatigue crack initiation and growth testing, together with creep strain rate, creep rupture and creep crack growth testing, on new and original pipework steels and weldments. In parallel, microstructural modelling and characterisation techniques will be employed to predict the microstructural changes that will occur in steel components operating in high temperature plant components. This will be used to suggest suitable combinations of time, temperature and possibly stress that could be used to accelerate the ageing process in such a way as to provide material, within the timescales of the project, which could be used to determine the important mechanical properties pertaining to defect development. Because of the requirement for ever more efficient coal fired power plant and petrochemical plant, the demands on the properties of the steels from which the plant is constructed are also increasing. This has resulted in the development of more highly alloyed steels to meet the needs of improved creep and corrosion resistance. A consequence of such developments is steels with less stable microstructures than those currently used. For new plant, the changes in microstructure and properties can be expected to be of little significance. However, as the plant becomes older, service induced microstructural degradation may have significant effects on defect initiation and propagation rates and critical defect sizes. Modelling of welded components will also be required to analyse the effects of geometry, system support and static and transient plant loading on performance in service. This is necessary to determine the design advantages of new materials by properly assessing the interactions between materials properties and plant performance. For example, materials that are stronger in creep can be designed with reduced thickness, thus lowering thermal gradients in service conditions, and thereby improving materials performance under the thermal fatigue loadings that arise in cyclic and flexible operation of power plant.

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Summarising, the aims of the project are the following: - Mechanical and microstructural assessment of base materials, - Development of weld material and welding procedure to avoid the “bore cracking”

phenomenon, - Mechanical and microstructural assessment of similar and dissimilar welding, - High temperature design and assessment of welded components including welded pressure

vessels and pipework, - Microstructural modelling to predict changes as a function of time and temperature in

service operation. - Piping integrity assessment under realistic loading conditions by combined thermal

hydraulic system analysis and stress analysis tools. 2. WORK PROGRAMME AND DISTRIBUTION OF TASKS

The project is organized in the following work-packages: WP1: Characterisation of the ‘as-received’ steels (WP Leader: CSM) WP2: Development of weldments (WP Leader: ISQ) WP3: Mechanical data assessment (WP Leader: VTT) WP4: Microstructural assessment (WP Leader: CSM) WP5: High temperature design and assessment (WP Leader: Powergen) WP6: Project management (WP Leader: CSM) WP1: Characterisation of the ‘as-received’ steels This WP deals with the selection of the materials and is constituted by 3 tasks: Task 1.1: material production Task 1.2: microstructural characterization of base material Task 1.3: mechanical characterization of base material The output is towards WP3 Mechanical data assessment WP2: Development of weldments This WP deals with the definition of welding procedures, including consumables. The WP is divided into two tasks: Task 2.1: welding trials and modelling Task 2.2: welding qualification The output is towards WP3 Mechanical data assessment WP3: Mechanical data assessment This WP intends to provide high temperature mechanical data for materials and welded joints. It is devided into two tasks: Task 3.1: Determination and comparison of mechanical properties Task 3.2: Data base and assessment The output is towards WP5 High temperature design and assessment WP4: Microstructural assessment This WP intends to evaluate, also by means of modelling, the microstructural evolution of the material and welded joints. It is divided into two tasks: Task 4.1: Microstructural modelling and assessment Task 4.2: Characterization of aged material The output is towards WP5 High temperature design and assessment

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WP5 High temperature design and assessment This WP aims at reviewing design requirements for welded components, giving also the final assessment of the whole results. It is divided into three tasks: Task 5.1: High temperature design review Task 5.2: Structural assessment of the effects of typical plant operating conditions Task 5.3: Experimental validation of component simulation WP6: Project management This WP deals with the overall management of the project. It is divided into one task: Task 6.1: Project management and reporting The project flow chart is represented in the following page, while work packages are detailed with the specific forms.

PROJECT FLOW CHART Material selection Welding selection

Task 1.1

Task 1.2 Task 1.3

Task 2.1

Task 2.2

Task 3.1

Task 3.2

Task 4.1 Task 4.2

Task 5.1 Task 5.2

Task 5.3

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WP1: Characterisation of the ‘as-received’ steels (WP Leader: CSM)

WORK PACKAGE DESCRIPTION WP N° 1

Workpackage Title Characterisation of the ‘as-received’ steels EstMan/Months

WP Leader CSM 7

Contractor (s) PG (1), UsI (11), Dal (3), ISQ (3), IWM (0,5), VTT (1) 19,5

Etc.

Total 26,5

Objectives : • Selection of test materials and plate/tube geometries • Microstructural characterisation of base material in normalised and tempered condition of base material • Mechanical properties of base material including: strength, toughness, fatigue crack initiation and crack growth, minimum creep and creep crack growth rates,

Description of work /tasks including results : Task 1.1: material production 1.1.1 Production of material Grade 23 and 24 steel by industrial routes (UsI, Dal) or on the

laboratory scale (CSM) 1.1.2 heat treatment of base material for testing and conventional mechanical tests for

product qualification, (UsI, Dal) Task 1.2: microstructural characterization of base material 1.2.1 Characterisation of microstructures using light (LM) and electron microscopy (SEM, TEM) with EDS analysis and other analytical techniques to obtain background information on the grain size, dislocation density, type, dimension and density of precipitates present in the matrix, (CSM, UsI, ISQ, IWM, VTT) Task 1.3: mechanical characterization of base material Characterisation of as-received materials properties by the following tests: 1.3.1.Tensile strength from room temperature up to 650°C, (UsI, Dal)

1.3.2 Hardness, (UsI, Dal, PG) 1.3.3.Impact and FATT curve (UsI, Dal, PG) 1.3.4 Short term creep test at 550, 600, 650°C with 2 stress levels to obtain: 1k, 3k hours rupture. (CSM, Dal, VTT ) 1.3.5 Selection of the most promising material (All partners) 1.3.6 Fracture toughness (UsI, ISQ,)

Interrelation with other workpackages (please give WP No) : 3

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WP2: Development of weldments (WP Leader: ISQ)

WORK PACKAGE DESCRIPTION WP N° 2

Workpackage Title Development of weldments EstMan/Months

WP Leader ISQ 19

Contractor (s) CSM (5), PG (7), UsI (5), Dal (1), IWM (0,5), VTT (11) 29,5

Etc.

Total 48,5

Objectives : • development of welding consumables • set up of welding procedure (welding technology and parameters) • qualification and simulation (using Gleeble testing) of welded joints • availability of mechanical properties of welded joints Description of work /tasks including results : Task 2.1: welding trials and modelling: 2.1.1 Development of welding consumables (max 3 composition) for P23 and P24 steel to avoid premature low creep ductility failure, (ISQ, PG) 2.1.2 Welding simulation and modelling, using Gleeble testing for various thermal cycles and modelling of heat treatments (IWM, VTT) 2.1.3 Welding trials to select the welding technology (SMAW, TIG, FCAW, NGTIG, or submerged arc) and the welding parameters (max.3 conditions) (ISQ, VTT) 2.1.4 Welding consumables and short term creep testing as a first filter to reject unsuitable developments and select consumables for further assessment (CSM, PG, Dal, ISQ, VTT (1 Temperature 4 stress levels) 2.1.5 Short term creep testing (1-3kh) of welded joints (CSM, Dal) 2.1.6 Selection of the more promising consumables and parameters (one-two conditions) (CSM, PG, ISQ, IWM, VTT)

Task 2.2: welding qualification: 2.2.1 Microstructural analysis of welding in Heat Affected Zone and weld metal (ISQ, IWM, VTT, PG) 2.2.2 Qualification of welded joint including: tensile, bending, impact and hardness profile of welded joint, (ISQ, VTT) 2.2.3 Strain monitored creep testing of weld metals (3 temperatures, stress to obtain 1k, 3k, 5k hours rupture) and analysis of creep ductility and damage. (CSM, PG, ISQ, VTT) 2.2.4 Selection of most promising material and weldments (CSM, PG, ISQ, IWM, VTT) 2.2.5 Long term cross-weld and all-weld metal creep testing (3 temperatures with stresses to obtain 10k, 20k, 30k, >30k hours rupture) to improve our capability to predict the risks of plant service failures at welds, (CSM, UsI, Dal, VTT , PG) 2.2.6 Development of dissimilar welding – P23 (or P24) welded to P91, to low alloy steel (e.g. P22), and to austenitic steel (nickel-based filler) – Hardness, impact, and short/medium term creep testing (1-3-5kh) (CSM, VTT)

Interrelation with other workpackages (please give WP No) : WP1, WP3, WP4

138

WP3: Mechanical data assessment (WP Leader: VTT)

WORK PACKAGE DESCRIPTION WP N° 3

Workpackage Title High temperature testing and data assessment EstMan/Months

WP Leader VTT 9

Contractor (s) CSM (9), PG (9), UsI (6), Dal (1), ISQ (14) 39

Etc. -

Total 48

Objectives : • Provision of high temperature mechanical data assessment for base material • Provision of high temperature mechanical data assessment for welded joint • Determination of weld stress reduction factors Description of work /tasks including results : Task 3.1: Determination and comparison of mechanical properties: High temperature testing and characterisation of as-manufactured materials by: 3.1.1 fatigue crack initiation and crack growth tests (PG, UsI, ISQ) 3.1.2 long term creep tests in the range 550-650°C with 4-5 stress levels to obtain rupture up to 30kh, with strain measurement to have the minimum creep rate, (CSM, UsI, Dal, ISQ, VTT); 3.1.3 creep crack growth tests, (PG, UsI, ISQ, VTT) 3.1.4 comparison testing of welds (PG) and heat affected zones (CSM, VTT) 3.1.5 Comparison with literature data and with original pipe steels and weldments (CSM, PG, ISQ, VTT) Task 3.2: Data base and assessment: 3.2.1 Weld creep data assessment to improve capability for long term design and operation of advanced high temperature plant using low alloy ferritic steels, (CSM, PG, UsI, ISQ, VTT) 3.2.2 Database compilation on high temperature properties of steels and weldments – Creep rupture, creep strain data, creep crack growth, creep-fatigue, low cycle fatigue initiation and growth, (CSM, PG, UsI, ISQ, VTT)

Interrelation with other workpackages (please give WP No) : WP1, WP2 and WP5

139

WP4: Microstructural assessment (WP Leader: CSM)

WORK PACKAGE DESCRIPTION WP N° 4

Workpackage Title Microstructural assessment EstMan/Months

WP Leader CSM 16

Contractor (s) PG (1), ISQ (3), IWM (4), VTT (2) 10

Etc.

Total 26

Objectives : • Evaluation of the microstructural evolution of the base material and welded joint • modelling of microstructural evolution

Description of work /tasks including results : Task 4.1: Microstructural modelling and assessment: 4.1.1. Microstructural assessment of materials, weld metals and heat affected zone structures in order to study the evolution of employed and/or tested material to predict the microstructural changes that will occur in power plant steels, (CSM, ISQ, IWM, VTT): TEM analysis of precipitates: type, chemical analysis, dimension, density, 4.1.2 Microstructural modelling to predict changes as a function of time and temperature and the production of a time temperature accelerated ageing matrix (CSM, VTT) Task 4.2: Characterization of aged material: 4.2.1 On-going assessment of microstructural evolution during ageing at 550, 600 and 650°C at 1k, 3k, 10k hours (CSM, IWM,VTT): type, chemical analysis, dimension, density, 4.2.2. Characterisation of the aged materials mechanical properties as above (CSM, PG, ISQ, VTT): aging at 550 - 650°C up to 10k hours. Tensile, impact tests.

Interrelation with other workpackages (please give WP No) : WP1

140

WP5: High temperature design and assessment (WP Leader: Powergen)

WORK PACKAGE DESCRIPTION WP N° 5

Workpackage Title High temperature design and assessment EstMan/Months

WP Leader PG 7

Contractor (s) IWM (15), VTT (8) 23

Etc.

Total 30

Objectives : High temperature design and assessment of welded components including seam welded pressure vessels and pipework. Description of work /tasks including results : Task 5.1: High temperature design review: 5.1.1 High temperature design review of welded components for seamless butt welded pressure vessels and pipework. (PG, IWM, VTT) 5.1.2 High temperature design review of welded components for seam welded pressure vessels and pipework. (IWM, VTT) Task 5.2: Structural assessment of the effects of typical plant operating conditions 5.2.1 Assessment of the effects of typical plant operating cycles on the structural integrity of pipework. (PG, IWM, VTT). 5.2.2 Assessment of the effects of pipework system support on the structural integrity of pipework (PG, VTT). 5.2.3 Assessment of in service effects of temperature and pressure distribution and flow in a selected boiler case. Correlations to structural studies (task 5.1 and manufacturing WP 2 and 3) – VTT contribution 4 man months (VTT Processes) using eg. APROS and 4 man-months (VTT) using ABAQUS. Task 5.3: Experimental validation of component simulation 5.3.1 Material modelling for P23 and P24 materials (base metal, HAZ, weld metal) (IWM). 5.3.2 Verification of the simulation work by simulation (stresses and deformation) and lifetime prediction of one tube test. Simulation of the behaviour of one welded component under plant loading conditions and lifetime prediction. (IWM). 5.3.3. 6 tests on tubes with similar and/or dissimilar welds under internal pressure and cyclic thermal loading over 1 to 2 years (depends on component availability). Accompanying characterisation of microstructural evolution during the tube tests is performed in WP4 (IWM)

Interrelation with other workpackages (please give WP No) : WP3

141

WP6: Project management and reporting (WP Leader: CSM)

WORK PACKAGE DESCRIPTION WP N° 6

Workpackage Title Project management and reporting EstMan/Months

WP Leader CSM 4

Contractor (s) PG (2), Dal (2), UsI (2), ISQ (2), IWM (2), VTT (2) 12

Etc. -

Total 16

Objectives : Project management.

Description of work /tasks including results : Task 6.1: Project management and reporting: • Project co-ordination and meetings (CSM) • Participation in the ECSC official meetings (CSM) • WP co-ordination and meetings (CSM, VTT, PG, ISQ, IWM) • Half year reporting (CSM, PG, UsI, Dal, ISQ, IWM, VTT) • Annual reporting (CSM, PG, UsI, Dal, ISQ, IWM, VTT) • First Milestone for selection of material (18th month) (CSM, PG, UsI, Dal, ISQ, IWM, VTT) • Second Milestone for selection of consumables and welding parameters (24th month) (CSM, PG, UsI, Dal, ISQ, IWM, VTT) • Final report (CSM, PG, UsI, Dal, ISQ, IWM, VTT)

Interrelation with other workpackages (please give WP No) : WP1-5

142

European Commission

EUR 23598 — Applications of advanced low-alloy steels for new high-temperature components

A. Di Gianfrancesco, D. Venditti, D. J. Allen, A. Morris, S. Caminada, S. Pillot, M. M. Rodriguez, V. Friedman, P. von Hartrott, D. Siegele, S. Holmström, J. Rantala, J. Salonen, P. Nevasmaa, K. Calonius, P. Junninen

Luxembourg: Office for Official Publications of the European Communities

2009 — 142 pp. — 21 × 29.7 cm

Research Fund for Coal and Steel series

ISBN 978-92-79-10006-2

ISSN 1018-5593

Price (excluding VAT) in Luxembourg: EUR 20