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ISSN 1018-5593

European Commission

COST

physical sciences

New steels and manufacturing

processes for critical com ponents

in advanced steam power plants

1996 EUR 16 85 8 EN



European Commission

COST

physica l sciences

New steels and manufacturing

processes for critical components

in advanced steam power plants

K. H. Mayer, C. Berger, R. B. Scarlin

MAN Energy

Nuremberg

Germany

Edited by

P. J -L Mériguet

DG

XII/B 1

COST materials

Rue de la Loi 200

B-1049 Brussels

Supported by the

European Comm ission through Contract No COST 92-0049DE

Directorate-General XII

Science, Research and Development

1996 EUR 16858 EN



Published by the

EUROPEAN COMMISSION

Directorate-General XII

Science, Research and Development

B 1 49

Brussels

LEGAL NOTICE

Neither

the

European Com mission nor any person acting

on behalf of the Com mission is responsible for the use which

might be made of the following information

Cataloguing data can be found at the end of this publication

Luxembourg: Office for Official Publications of the European Com mun ities, 1996

ISBN 92-827-6578-4

© ECSC-EC -EAEC, Brussels · Luxembourg, 1996

Printed in Belgium



New Steels and Manufacturing Processes for Critical

Components in Advanced Steam Power Plant

R.B. Scarlin, ABB Power Ge neration Ltd., Baden

K.H. Mayer, MAN Energy, Nürnberg

C. Berger, Siemens Power Generation KWU, Mülheim

Summary

An increase in the operating temperature and pressure of a steam power plant leads to an in-

crease in the system efficiency. Although the use of austenitic steels would permit such an in-

crease these materials suffer from the disadvantage of high price and susceptibility to thermal

fatigue, caused by the higher coefficient of thermal expansion and low thermal conductivity.

For this reason improved ferritic steels were required to minimise turbine and boiler costs and

provide high flexibility of operation (2 shift operation, frequent start-up/shut down). Such steels

are also the subject of extensive research programmes in Japan an d the USA.

The long-term aim of the COST programme was to develop and evaluate improved creep re-

sistant 9 -12% Cr steels and to manufacture, test and seek service operation of critical com-

ponents required for an advanced steam power plant (steam temperature of 600'C and at su-

percritical pressure). Critical components are:

• High-pressure and intermediate-pressure rotors

• Turbine and valve casings

• Turbine and valve bolting

• Main steam pipes and header sections

• Waterwalls

For each of the critical components a working group was con stituted c omprising:

• steel companies (forgemasters or casting foundries)

• turbine and boiler manufacturers

• utilities and other users

• testing institutes and universities.

The participants are listed in Table 1.

The development goals in terms of required materials properties, fabrication techniques

(forging, casting, welding) and nondestructive examination had been defined by the turbine

and bo iler m anufacturers.

ι -



Within the programme alloy de velopment work was firstly performed on small batches of ma te-

rial (150 - 500 kg), some of which had been be gun w ithin the first round of COST

501.

Subse-

quently for the best steels representative components were manufactured to demonstrate the

feasibility of up-scaling. These components were subjected to nondestructive and destructive

testing.

Steels w ere developed which were able to satisfy the targets set for large com ponents, so that

steam power plant can now be built with operating temperatures up to 600'C; i.e. about 35'C

hotter than was previously possible, with a corresponding relative increase in operating

effi-

ciency of c a. 2%.

Organisation

ABB-Sweden

ABB Powdermet

ABB-Swiizerland

AEG-Kanis

Ahlström

Ansaldo

Austrian Research Centre

Bôhler

Dalmine

ENEL

Energie und Verfahrenstechnik (EVT)

ETH-Zúrich

Forgemasters Engineering Ltd.

Fraunhofer Institute

GEC Alsthom

Georg Fischer

(with Schwe issindustrie, O erlikon)

MAN Energie

(with MW, Darmstadt)

Mannesmann

National Power

NEI-Parsons

Royal Scheide

Saarstahl, Völklingen

Siemens-KWU

Stork Boilers

Sulzer Bros.

Techn. Ueberwachungsvereinigung

Vallourec

Vereinigte Schmiedewerke GmbH

Voest Alpine

Forgings

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

Castings

X

X

X

X

X

X

X

X

X

X

X

X

Bolts

X

X

X

X

X

Steam Pipes

χ

χ

χ

χ

χ

χ

ΡΜ Header

χ

χ

χ

χ

Waterwalls

χ

χ

χ

χ

χ

Table 1: List of Participants



Contents

Preface - -

1.

Introduction - 5 -

2. International Developm ent for Adva nced Steam Power Plant - 5 -

3. Crit ical Com pone nts in Advan ced Steam Power Plant - 7 -

4. Materials Developm ent - 8 -

5. New Ferritic-Martensitic Rotor Steels - 10 -

6. New Ferritic-Martensitic Cast Steels - 19 -

7. Creep -Resistant Bolting Material - 24 -

8. Improved Steels for Steam Pipes and Headers - 30 -

9. Con clusions and Future Trends - 48 -

10.

Acknowledgem ents -

4 9

-

11 .

References - 49 -



1.

INTRODUCTION

Increasing fuel costs, the pressure to reduce environmental pollution and the need to

reduce C02-emissions have lead worldwide to the development of power plants with

higher efficiency, greater operating flexibility, improved availability and longer lifetime

[1].

A key role in the further development of power plant technology has be en played by

the materials for highly-loaded turbine components, since basically the aims can only

be achieved through using materials w ith improved strength and toughness [2 - 4]. This

is particularly clear in Fig. 1 [5] which shows schematically the reduction in heat rate for

a steam turbine of up to 10% achieved by increasing the steam temperature from about

540 to 650'C with a simultaneous increase in steam pressure from about 180 to 300

bar, with double reheat.

A major increase in operating efficiency is possible. A temperature increase to 600'C

constituted the first step in the European COST501 Programme on Critical Compo

nents for Advanced Steam Cycles. It is considered possible to make a further increase

to about 620"C, through the use of improved creep resistant ferritic steels. The further

step to about 650'C can only be achieved through the use of highly creep resistant

austenitic steels, i.e. through the use of more expensive steels. In addition to the costs

it is also necessary to consider the effect of the different properties such as higher

coefficient of thermal expansion, lower thermal conductivity and higher susceptibility to

stress corrosion cracking.

2.

INTERNATIONAL DEVELOPMENT FOR ADVANCED STEAM POWER PLANT

2.1 JAPAN

Since there are no fuel reserves in Japan a programme was initiated in 1979 und er the

leadership of the power plant operating organisation EPDC (Electric Power Develop

ment Corporation) and with financial support from the government (MITI). The aim of

this common initiative of operators, manufacturers and the government has been the

development of a 1000 MW plant with a maximum steam temperature of 593'C in the

first phase and of 649'C in the second phase, with single or double reheat and super

critical pressure, see Fig. 2 [6, 7]. The figure shows how the live steam and reheat

steam temperatures are increased in steps, whereby the reheat steam temperature

(where the pressure is lower) in generally raised first. All plant named on this figure are

either comm issioned or under construction.

In Fig. 3 the materials for this development project are shown for the various compo

nents.

The investigation of materials and components in the laboratory and power plant

continued until 1988. At the beginning of the 80's a start was made with the manufac

ture of a 50 MW demonstration plant, Wakamatsu. Subsequently test operation was

successfully performed over a period of several years, in the first stage at 593'C [8].

The start of test operation for the second phase (up to 649'C) was planned for August

1990. The satisfactory progress of the research programme led already in 1989 to the

order for a 700 MW steam power plant by a Japanese operator, with the steam condi

tions 241 bar 538'C/593'C. Commissioning of the plant was in June 1993 [9]. A

1000 MW plant with "ultrasupercritical" steam conditions and b oth live steam a nd single

reheat steam temperature of 593'C was ordered for commissioning in 1997.

In addition EPDC plans to man ufacture a demonstration combi po wer plant with a pres

surised fluidised bed boiler, a gas turbine and supercritical steam conditions. The de

monstration of this concept is also planned to be made in the Wakamatsu power plant,

subsequent to the test ope ration of stage two (649"C).



2.2 USA

In 1978 the American power plant operator company EPRI (Electric Power Research

Institute) initiated two feasibility studies for the following steam conditions:

Phase 0

Phase 1

Phase 2

double reheat 316 bar, 566'C

double reheat 316 bar, 593'C

double reheat 352 bar, 649'C.

Both studies led to recommendations for further development to 316 bar and 593'C

with double reheat, since with these parameters it is possible to achieve the maximum

improvement in heat rate without a loss of reliability, at relatively low research and de

velopment costs [10]. These studies led in 1986 to the initiation of a five year EPRI Re

search Programme in which certain Japanese and European power plant manufac

turers also p articipated. The essential aims of the research p rogramme are summ arised

in Fig. 4. In addition to these aims further requirements were also specified for the

turbines:

• short start-up times (cold start 10 - 12h, warm start-up 4h and hot start-up 2h),

• suitability for peak load opera tion,

• improved reliability,

• improved efficiency,

• improved control and monitoring devices.

The progress of the development work was reported in 1986 [11], 1988 [12] and

April 1991 [13]. No plant with temperatures above 566'C have as yet been ordered in

the USA.

2.3 EUROPE

There is already extensive operating experience with smaller plant which was built in

the 50's particularly for the chemical industry e.g. [14 - 16], The plant built in Europe in

the 50's and early 60's is summarised in Fig. 5. The steam temperatures lie between

600 and 65 0'C and pressures between 180 and 330 bar. The power rating of the

plants, which were mostly built for the chemical industry, is comparatively low (3 to 125

MW).

For the highly loaded components austenitic steels and creep resistant ferritic

steels such as X21CrMoV121 or G- X22CrMoV121 were mostly used, sometimes with

cooling, see Fig. 6.

So far recent plans for power plant with advanced steam inlet temperatures are known

from English [17], Danish [18 - 20] and German power plant operators. A particular in

terest in pulverised coal fired plant with higher steam temperatures is present in Den-

marie. In a study by ELSAM [18] the development potential for pulverised coal fired

power plant is shown, which according to Fig. 7 is only exceeded by oil or gas fired

combined cycle processes w ith respect to overall efficiency. A pow er plant with a steam

temperature of 580'C and overall efficiency of 47.5% has been ordered [19]. The pro

cess 3 with an efficiency of well over 50% and a steam temperature of about 640'C

could be achieved after the year 2000, if new high temperature materials become

available.

Planning by the German Power Plant Operators for higher steam inlet temperatures

and large power plants is also gaining momentum. In a recent paper [21] it is stated

that coal and lignite will be used increasingly in Europe for the generation of electrical



power, since they are readily available and are preferable to oil and gas, for which re

serves are smaller and transport distances are generally greater. Since electricity re

quirements will continue to rise and CO2 emissions must be limited it will be nece ssary

to move to more efficient, high tempe rature, coal or lignite fuelled power plant. Current

capability is believed to be sufficient to construct plant with a live steam temperature of

580"C and pressure of 275 bar, along with a reheat temperature of 600'C. A VGB

conference in 1993 was dedicated to the subject of fossil fired power plant with ad

vanced design parameters. The views of the European turbine manufacturers concer

ning material selection and appropriate designs have been presented [22- 24].

In a number of literature references [16, 25 - 27] reports have been presented on fea

sibility studies performed for plant with advanced steam conditions (580'C to 650'C).

The materials selection was based on the available proven creep-resistant materials

[16, 28], whereby components manufactured of ferritic steels would require cooling. In

order to avoid the need for cooling such components the COST 501 Programme on

"Critical Components for Advanced Steam Cycles" has been aimed at the development

of 9 -12% Cr steels with improved creep properties at about 600'C. A major part of the

work was based on the results of a review of the high temperature properties of known

9-12% Cr steels, performed with the backing of the COST organisation [29]. The

review concluded that primary importance should be placed on obtaining a stable

microstructure through alloying and heat treatment, rather than on a high yield

strength, and that there is no substitute for actual long-term creep and exposure

testing. It was demonstrated that short-term tests and extrapolation techniques, can be

highly unreliable.

The progress made in the meantime in the COST Programme in the development of

improved ferritic creep resistant steels containing 9-12% Chromium indicates that

turbines with a maximum steam temperature of 600'C can already be designed without

cooling.

The potential of these steels in terms of strength and toughness has been

under investigation since 1983 [30].

An overview of the international research programmes on advanced coal-fired power

plants is shown in Fig. 8.

CRITICAL COMPONENTS IN ADVANCED STEAM POWER PLANT

Operating efficiency can be improved by an increase in the temperature and pressure

of the live and reheat steam and an increase in the temperature of the feedwater

entering the boiler to > 300'C (through incorporation of a further high temperature

steam extraction from the high pressure turbine). This leads to increased loading for

certain turbine and boiler components [23, 31]. Specifically, on the turbine side, higher

creep loading is experienced by:

• forged high pressure and intermediate pressure rotors and blades,

• cast high pressure and intermediate pressure inner casings and valve bodies

• bolts for securing flanged bodies such as turbine and valve casings

and, on the boiler side:

• main steam pipes and header sections

• waterwalls

• superheater tubes

- 7 -



The above-mentioned turbine components will be exposed to temperatures of up to

600'C and require specifically creep strength values at 600'C similar to those at 565'C

for the previous best m aterials in the 9 -1 2 % Cr steel class, along with at least the

same ease of fabrication. Main steam pipes and headers require a similar improvement

in creep properties although a certain advantage can be gained through the use of a

larger number of parallel pipes, through w hich the flow is distributed. The higher loading

of the furnace waterwall results from the increase in feedwater inlet temperature up to

300 'C or more. Consequently the outlet temperature may rise above 460 'C; a tempera

ture at which both creep strength and corrosion and oxidation resistance of the cur

rently employed low-alloy steels are no longer sufficient. H igher alloyed steels, such as

9 -12% Cr steels, are available but require, on account of their higher hardness in the

heat affected zone after welding, a post weld heat treatment, so that special precau

tions would have to be made for on-site assembly welds. Austenitic materials will

con

tinue to be used for thinner-wa lled superbeater tubes.

4. MATERIALS DEVELOPMENT

An overview of the historical development of the creep-resistant 9-12 %

CrMo(W)VNbN(B) steels is given in the upper part of Fig. 9. The lower part of the figure

shows the corresponding most recent values of the 100,000 hour creep strength at

600'C, extrapolated from long-term data. The steels X 22 CrMo(W)V 12 1, H 46, FV

448 and 56T5, developed in Europe (No. 1 ) and USA (No. 2) at the beginning of the fif

ties, have creep rupture strengths at 600'C and 100,000 hours of 60 to 64 MPa,

whereby only the Nb-free steel X 22 CrM o(W )

v

12 1 is suitable for thick-walled com po

nents.

The TAF steel (No. 4) developed in Japan for smaller components is a further

development of the E uropean Nb-containing steel (No. 3) (H 46, FV 448, 56T5) [32], as

also indicated in Fig. 10. In addition to an improved balance of the alloying elements, it

also has a boron content of 0.040%, which is too high to be achieved without major

segregation in large components. The creep strength, measured up to a time of about

30,000 hours at 600'C, lies at a very high value of about 200 MPa [33]. This indicates

the possible development potential of boron-containing steels. The rotor steel 11 %

CrMoVNbN steel (No. 5), patented in 1964 by GE, also represents a further develop

ment of the European Nb-containing steels [34]. In particular the Nb content was

greatly reduced in order to prevent harmful segregation in the rotor centre. Furthermore

the alloying elements were balanced in order to avoid the formation of delta ferrite. The

relatively high published creep strength of about 85 to 90 MPa [35] was extrapolated on

the basis of tests at 620'C up to times of 16,195 hours duration.

The steel often referred to in the literature as mod. 9Cr1 Mo or Ρ 91 (No. 6) already re

presents a steel of the newer generation. It was developed in the later 70's for the

manufacture of pipes and vessels in the American fast breeder programme. It is tough,

readily weldable and, as shown by creep tests at 593'C with durations of up to about

80,000 hours, has a high creep strength at 600'C and 100,000 hours of about 94 MPa

[36]. In comparison with earlier steels it is characterised, for example, by a lower C-

content of only about 0.10% and a Cr-content of 9%.

The Japanese rotor steels TR 1100 and TR 1200 (No. 7), developed in the 80's, were

based on the known properties of the steels 1 to 6 [37]. Compared with the GE rotor

steel in particular the C-content was reduced and the sum (C + N) was selected at

around 0.17% (Fig. 11

.

Based on the research work of Fujita, the Mo-content was

raised to 1.5% in TR 1100 whereas the W-content was raised to about 2% in TR 1150

and TR 1200, with a simultaneous decrease in Mo-content to 0.30% as shown in

Fig.

11 [38]. For the last steel creep results have only been published for times up to

about 10,000 hours at 600'C. For TR 1100 a creep strength of about 100 MPa is given

for 100,000 hours at 600'C, based on testing times of 30,000 hours. The rotor steels



No. 8 and No. 9 are primarily the result of the research work performed in the 80's

within the European co-operative programme COST501 [39]. Steel No. 8 ¡s a 9%

CrMoVNb steel alloyed additionally with about 0.01% boron. Creep test results, which

have so far reached a duration of about 42,000 hours, indicate a probable creep

strength at 100,000 hours and 600'C of about 120 MPa. The creep strength of the

steel X 12 CrMoWVNbN 10 11 (steel No. 9), containing about 0.8% W probably lies a

little lower. The longest testing time achieved so far is about 43,000 hours and the

creep strength obtained by extrapolation of these values to 100,000 hours at 600'C is

about 110 MPa [40]. Steel No. 10 is a 9%Cr pipe steel specifically alloyed with tungsten

and boron. It was developed in Japan in the second half of the 80's under the

designation NF 61 6. Based on the extrapolation of tests of upto 35,000 hours duration,

its creep strength for 100,000 hours is estimated to be about 132 MPa [41]. A similar

pipe steel (HCM12A) has also been developed in Japan, with a chromium content of

ca.

11 %

in order to improve oxidation resistance.

1 %

copper has been added to reduce

the tendency to δ ferrite formation [129].

An important characteristic of the successful variants of the newly-developed steels is

the continuous form of their creep rupture curves at 60 0'C , as shown in Fig. 12 in com-

parison with the DIN bolting steel X 19 CrMoVNbN 111, which is similar to the steel

56T5 developed in the 50's, and shows a sharp drop in creep strength at testing times

> 3,000 h.

An improvement in long-term creep strength is achieved by increasing the Mo^equiva-

lent (Mo% + 0.5W%) from 1.0 to 1.5%, as shown in Fig. 13 for the creep strength at

30,000 h. An overall consideration of the effect of composition on creep, ductility and

toughness properties of 9 - 1 2 % Cr steels is shown in Fig. 14, which shows the de-

pendence of these properties on the sum (C + N) and the Cr equivalent. A distinction is

made between four regions (A to D) differing in creep strength, ductility and toug hness.

The preferred region B, in which the steels currently under development are to be

found,

provides an optimum for the present applications, with acceptable creep

strength,

high ductility an d high toughness.

The difference in creep behaviour results from the different microstructures which are

determined by the chemical composition and heat treatment of the steels. Fig. 15 pro-

vides a summary of the characteristic strengthening mechanisms in the newly-deve-

loped steels in comparison with the steel X 22 CrMoV 12 1 traditionally used in steam

turbine manufacture. Basically the new steels exhibit more, smaller and more stable

chromium carbides of the type M23C6· In addition there are many small V/Nb carboni-

trides of the MX type and major solid-solution strengthening as a result of the higher

levels of molybdenum and tungsten. However, depending on the amounts of other

elements present, higher levels of tungsten may lead to the rapid precipitation of Laves

phases at grain and lath boundaries, greatly reducing ductility and creep strength.

Fig.

16 shows the appearance of Laves phases in the trial melt D1 which exhibited in-

ferior creep properties at 600 and 650'C. The stabilising effect of boron probably

comes from incorporation of boron within the carbides. This reduces the rate of carbide

coarsening hence improving long-term creep strength. The effect of microstructural

changes on creep strength is shown schematically in Fig. 17.

This representation of the development of the creep strength of the 9 to 12%CrMo

steels indicates a major improvement through a better balance of the alloying elements

and through addition of carbide formers and stabilising elements, along with an optimi-

sation of the heat treatment. Fig 18 shows in summary a series of steps taken during

the development of steels appropriate for use as small components, rotors, castings,

tubes and pipes, making use of known and supposed strengthening and stabilising

mechanisms. Whereas some of these steels are already in use, others are still under



development. The steps in alloy design, for an advanced tube/pipe steel H CM12A, cur

rently in the initial stage of development, are illustrated in Fig. 19 [129]. In this case

• corrosion resistance should be obtained using a high Cr content

• creep strength should be provided through a) grain refinement (Nb (CN))

b) solid solution strengthening (W, Mo)

c) fine carbon itride precipitation (V, Nb)

d) stable particles incorporating boron

e) limiting Nb content

• weldab ility is facilitated by the low C content

• toughness in ensured by limiting δ femte content (Cr equivalent by Cu addition)

• long-term embrittlement is suppressed by minimising Si content.

The following sections will deal specifically with the materials development and testing

performed within the COST501 programme, making comparisons with data from other

sources and for other materials w here appropriate.

5. NEW FERRITIC-MARTENS ITIC ROTOR STEELS

5.1 PROPERTY PROFILE

The steel properties aimed at were specified:

(A) For application temperatures approaching 600 'C:

100,000 h creep rupture strength at 600'C of about 100 MPa,

• good creep rupture ductility (> 10% elongation) and no notch sens itivity,

• through-hardening up to at least 1200 mm diameter,

• minimum yield strength of 600 or 700 MPa,

while other properties such as toughness and susceptibility to embrittlement should not

be worse than with conventional 12% CrMoV and

1%

CrMoV rotor steels.

(B) For conventional application tempe ratures:

• higher strength with good toughness or

• higher toughness without long-term embrittlement.

Especially for the first application range up to 600'C it was a main requirement to have

more stable microstructures obtained using higher tempering temperatures. A further

aim was the performance of very long-term creep tests in order to avoid the uncertainty

of extrapolation from short time tests. Full size rotors should be manufactured to get

experience in manufacturing these new steels as well as properties representative of

the inner and outer regions of a real component.

1 0 -



5.2 MATERIALS DEVELOPM ENT

Modem steam turbines are designed for a service life of over 250,000 hours. The ma

terials employed must have a microstructure which is a stable as possible so that their

properties are maintained throughout the turbine life. The design characteristics of ma

jor importance are tensile properties, resistance to creep deformation, creep rupture

strength and fracture toughness.

The development of materials to meet these requirements is a long process. Work is

still continuing to optimise the materials used in conventional steam turbines even

though the steam parameters and alloy types have not changed significantly for thirty

years.

The COST 501 activity on the development of new ferritic steels, which began in the

early 80's, has now progressed to the stage w here m aterials data on new ferritic steels

can be made available to the design eng ineer to utilise in the development of advanced

steam turbines [30, 42, 43] Further testing will be required to optimise these materials

and to provide an adequate materials data base. The general programme of this work

was:

(a) Alloy Selec tion. Potential alloys were identified after a critical review of existing

grades of 9- 12%Cr steel, steel making developments in Europe and develop

ment activities elsewhere in the world, particularly in Japan and America. Fig. 10

summarises these commercial and newly developed steel grades and the ex

pected creep rupture strength at 600'C after 10,000 and 100,000 h. The values in

brackets are extrapolated from tests performed at higher temperatures as well as

short-term tests [32 - 41]. Trial melts of the candidate steels were manufactured

and various heat treatments were applied.

(b) Trial Components. The most promising alloys were selected to gain experience in

the manufacture of full size components. To date three rotor forgings have been

produced.

These were subjected to detailed destructive examination to determine

the extent of segregation, variation in microstructure and properties throughout

the component and to determine long-term performance. The use of very long-

term creep tests to 100,000 hours is most important to avoid the uncertainties in

herent in the extrapolation of short-term test results, particularly for these very

complex high-alloyed steels [29].

(c) Microstructural Inves tigations. A detailed metallographic study of a large number

of test samples taken from long-term specimens is being undertaken to increase

our understanding of the effect of steel chemistry and heat treatment on the mi

crostructural stability and associated effects on material properties. The object of

these studies is to assist in the optimisation of the current alloys and the formula

tion of new 12%Cr ferritic steel alloy with improved properties.

(d) Data Base. In the longer term it will be necessary to conduct long-term creep

tests on a number of samples of the selected alloys in order to obtain an ade

quate da ta base to determine the extent of the m aterial data sca tter.

The current status of COS T programm e to develop new ferritic forged steels capable of

operating at steam temperatures up to 600'C is described.

Π



5.3 PRE-EVALUATION PROGRAM ME

Trial melts

After reviewing the existing grades of 9 - 1 2 % Cr steel [29, 44 - 52], steelmaking deve

lopments in Europe, the results of development in the first round of COST501 and de

velopment activities elsewhere in the world, five grades of 9 - 1 0 % C rMoVNb steel were

identified as candidates for development.

Steel Group A: Nitrogen Grad es

Recent developments in ESR steelmaking involving high nitrogen pressures have

enabled the production of steels with nitrogen levels as high as 0.3% [53]. Ten

high-

nitrogen melts with chromium levels of 9% and 12%, nitrogen levels between 0.09 and

0.22% and tungsten levels up to 0.96% w ere manufactured for characterisation.

Steel Group B: Boron Grades

The earlier COST501/I programme had already identified a 10%CrMoVBNb steel (Melt

B0),

with an optimised Β addition of 100 ppm, as having a major potential for high tem

perature applications [30]. Therefore, this steel was produced as a large scale forging

without preliminary characterisation of a trial melt.

Steel Group D: Tungsten Grades

Work in Japan has indicated that addition of tungsten , in partial replacement of Mo, at a

level approaching 2%W leads to significant increases in creep strength [49, 50, 54]

prompting the inclusion of such a steel in the current work. In total 3 test melts were

produced with carbon contents between 0.12 and 0.16% and chromium contents bet

ween 10.25 and 11.30%. In each case the Mo equivalent was 1.2%.

Steel Group E: Tungsten/molybdenum Grades

Three melts were selected to investigate the effect of more moderate tungsten levels

(0.5 -1.0%W) while retaining Mo levels at about

1%.

Steel Group F: Molybdenum Grades

Three test melts were selected with molybdenum levels between 1.14 and 1.89% and

C levels between 0.10 and 0.17%, based on the Japanese research on the TAF [32]

and TR1100 [37] steels. This steel was selected as a 10% CrMoVNbN steel with no

additions of W or B. However, an optimisation of C, Mo and Nb levels compared with

existing 12% CrMoVNbN alloys was attempted, influenced by work on such steels in

Europe [44, 55] and Japan [49 ].

Apart from steel B, all other candidate steels were produced as melts (50 - 1000 kg)

leading to selection of the steels to be used for full-scale forgings. The chemical ana

lyses of these trial melts is described in Fig. 20 (see also Fig. 23). Variants of all steels

were produced to investigate the analysis ranges normally required for melting and the

effect of any segregation which may occur in a large forging and in addition the influ

ence of carbon, chromium and other alloying elements on properties. At least 15 melts

were investigated.

12 -



Time-Temperature-Transformation (TTT) diagrams, Fig. 21, were derived for each steel

to establish its transformation behaviour and tests were also performed to investigate

the materials' hot workability.

The material which was formed to the shape of bars was heat-treated at different

austenitising temperatures as shown below for addition of:

A nitrogen at 1120°

Β boron at 1100° C

D tungsten at 1020,

E tungsten / molybdenum

\

1070 and

F molybdenum J 1120°C

The cooling rate on quenching in oil after austentising was 27min . and simulates the

centre of a rotor with a diameter of 1200 mm. A first tempering treatment at 570'C

sen/es to transform the remaining austenite into martensite, so that after the second

tempering treatment only tempered martensite is to be expected. On the one hand the

tempering temperature should be as far as possible above the operating temperature,

anticipated to be 600'C, however, on the other hand the requirement for a yield

strength of min. 600 or 700 MPa requires that the tempering temperature should ge

nerally lie at or above 7 0 0 Ό

Test programme

The test matrix in Fig. 22 shows the extent of the investigation programme, this being

further multiplied by the additional analysis variants of the basic compositions A to F.

Due to the different melts and heat treatments about 40 conditions were tested. The

following tests were performed on these heat treated bars:

• tensile tests at room temperature, 600 and 650'C ,

• impact tests to determine the FATT and room temperature and upper shelf

energies,

• isothermal rupture tests at 600 and 650 'C using plain and notched specimens at a

minimum of three stress levels giving durations in excess of 10,000 h,

• iso-stress rupture tests at 100 MPa and temperatures between 700'C and 6 2 0 Ό

Material was also aged at 480, 600 and 650'C for durations of up to 10,000 h. In order

to determine any effect on tensile and toughness properties, the following tests were

performed on these aged materials:

• tensile tests at room temperature, 600 and 650'C ,

• impact tests to determine the FATT and room tempe rature and upper shelf

energies.

In an attempt to model the effect of metallurgical changes in service and to provide a

better extrapolation of the creep strength, an overaging treatment (700'C, 200 h) was

applied to some lengths of bar. The selected heat treatment of 200 h at 700'C theoreti

cally corresponds (calculated with a time/ temperature parameter according to Larson-

Miller of 25) to a period of about 270,000 h at 600'C. In this way microstructural altera

tions and a consequent reduction in strength which would occur during service are

already simulated. The creep curves begin at lower initial strength values but are flatter,

- 1 3 -



so that extrapolation to longer times is more reliable [56]. The following tests were per

formed on this overaged material:

• tensile tests at room temperature, 600 and 65 0'C ,

• impact tests to determine the FATT and room tempe rature and upper shelf

energies,

• isothermal and iso-stress rupture tests as performed on unaged material.

5.4 RESULTS AND DISCUSSION

The symbols used to designate the different m elts on the subsequent figures are shown

in

Fig.

23.

All these steels were fully hardenable, showing 100% martensitic microstructures. Low

austenitising temperature resulted in grain sizes between ASTM 3 and 7, whereas

higher austenitising temperatures of up to 1120"C resulted in larger grain sizes of

ASTM 2 to 0, Fig. 24. In gene ral, increasing the austenitising tem perature also leads to

an increase in yield strength.

Strength and Toughness Properties

The required aims of yield strength > 600 MPa or > 700 MPa could generally be

attained with the required tempering temperatures. For the nitrogen steel considerably

higher yield strength values were measured and for the molybdenum steel the yield

strength > 600 MPa could only narrowly be attained w ith the highest tempering

temperature. The results shown in the upper part of Fig. 25 represent a selection of the

values after the heat treatment 1070'C (1100'C for boron and 1120'C for nitrogen

steel) + 570"C + approx. 720'C. Overageing significantly reduces the yield strength.

With increasing exposure temperature and time the boron and tungsten alloyed steels

remain relatively stable, whereas the molybdenum variants show a minor reduction and

the nitrogen variants a significant reduction of the yield strength with increasing

exposure temperature and time.

The long-term toughness behaviour, shown at the bottom of Fig. 25 determined through

exposure of the material up to 10,000 h, shows no alteration at 4 8 0 Ό At 600 and

650'C exposure temperature an increase in the FATT values by about 25'C can be ob

served.

The nitrogen steel shows no measurable alteration of FATT while the molyb

denum variant with a FATT < 0'C shows excellent behaviour even after long-term ex

posure.

In addition, for comparison, results of a commercial 12% CrMoV rotor steel manufac

tured according to the German standard SEW are included in Figs. 25 and 26. Material

from the central part of a rotor has been exposed at 480 and 530'C up to 10,000 h.

The comparison shows that the toughness of the trial melts E and F and rotor B is

much better after ageing at 600'C than for the conventional steel aged at 530'C. The

long-term ageing of Rotor E and F material confirms this behaviour (Fig. 27). Thus,

there is a real improvement of toughness with negligible long-term embrittlemen t.

14



Creep strength

The creep behaviour derived from tests at 600 to 650*C is shown as a function of the

Larson-Miller Parameter in Fig. 28 (for yield strength 600 - 650 MPa) and Fig. 29 (for

yield strength 700 - 750 M Pa). The results are from material austenitised at 1 070 'C, the

heat treatment selected for the tungsten and molybdenum modifications. Testing times

of up to 40,000 h have been achieved so far. In comparison with the mean values of

the conventional 12% CrMoV steel, according to SEW, the presently available creep

test results show a significant improvement for nearly all analysis variants inve stigated.

This applies both to the as-delivered condition and to the overaged condition (700'C /

200 h), as shown in Fig. 30 (YS of 600 - 650 MPa) and Fig. 31 (YS of 700 - 750 MPa).

However, the latter exhibit a flatter form of the creep curve thereby approaching the

values for the as-delivered condition at longer times. It is seen that the effect of over-

ageing on the creep results is less for the lower yield strength condition than for the

higher yield strength co ndition.

A comparison of all results for D, E and F of short-term creep rupture and ¡sostress

tests is shown in Fig. 32. These data were used along with the ongoing long-term creep

data to select variants for further investigation. Particularly at longer testing times a

marked reduction in the creep strength of steel D was noted.

Unfortunately steel A was shown to have a rupture strength lower than expected from

the results of earlier work [53]. This was attributed to the slow cooling rates applied

after austenitising to simulate the centre of a large forging. It must be concluded that

steel A is not suitable for application in such thick sections, although this conclusion

should not prejudice its potential future application in thinner sections such as discs,

blading and bolting.

When comparing these results with the Japanese results for TMK1 (TR 11Û0) [49] the

higher initial strength of the Japan ese s teel, as a result of the lower temp ering temp era

ture (680'C), should be taken into account. The results for the COST 501/11 programme

melts shown in the Figs. 28 and 29 show a flatter curve and confirm or exceed the

creep results obtained so far on the TMK1.

Higher tempering temperatures are favoured for the highest operating temperatures to

give the most stable microstructures and provide a greater margin between operating

and tempering temperatures. Nonetheless, it is recognised that these steels also have

advantages over current alloys used under conventional steam conditions and the short

term advantage of the high yield strength condition at temperatures above 600'C may

extend to the long-term at conventional temperatures around

540Ό Therefore, i t was

decided to investigate both 600 and 700 MPa yield strength conditions in the full-scale

forgings.

Consideration of the prom ising results from these trial melts led to a high level of confi

dence that the objectives of the programme can be met. Furthermore, the steels deve

loped for use at 600'C show advantages in rupture strength over current alloys for use

at conventional temperatures and the combination of good creep strength and very

good toughness suggests the alloys would be applicable to small single cylinder ma

chines which are used for industrial applications and in some smaller combined cycle

power plants.

SELECTION OF STEELS FOR FULL-SCALE FORGINGS

Careful analysis of all results permitted the identification of those steels and heat treat

ments most promising for the production of full-scale rotor forgings. Greatest weight

- 1 5 -



was given to the attainment of the creep rupture strength target as indicated by iso

thermal and iso-stress rupture testing in both the as heat treated and overaged condi

tions.

The attainment of good toughness, both in the as heat treated and aged condi

tions was also taken into account. Finally, the stability of these properties with respect

to variations in chemistry and heat treatment was considered .

This latter aspect was only possible due to the very large number of chemistries and

heat treatments tested, this in itself being possible only through the collaborative nature

of the programme.

In addition to the already manufactured first trial rotor B, two steels were identified as

most promising. Steels E and F also show ed an exce llent combination of creep rupture

strength (Figs. 28 and 29), iso-stress rupture times at σ =100 MPa (Fig. 33) and tough

ness.

It can be seen that ex trapolation of these resu lts indicates that the target strength

of 100 MPa at 600'C is comfortably exceeded both in the as heat treated and overaged

conditions. The austenitising temperature of 1070"C was identified as giving the opti

mum balance of properties for bo th steels E and F.

5.6 MANUFACTURE OF FULL SCALE ROTORS

Based on these data the second stage of the COST 501/11 programme was initiated in

December 1989, in which the manufacture of two further trial rotor forgings with dia

meters of up to 1200 mm and the analysis of E (tungsten/molybdenum) and F

(molybdenum) was agreed upon. In addition to the chemical compositions, the most

important data have been compiled in Fig. 34. The data for the rotor Β (boron) with a

diameter of 840 mm , previously manufactured in 1988, are also illustrated [57].

Rotor B, with a heat treatment diameter of 840 mm, was manufactured from a 15 Mg

ESR ingot (1150 mm diameter). To get a constant boron content of 100 ppm all over

the ingot, about 0.7% boron oxide was added to the slag. This procedure avoided a

transfer of boron from the melted steel into the slag during the electro-slag remelting

process. The forging was heated to a maximum temperature of 1150'C and forged with

the following procedure: stretching

1.7:1,

3.6:1 upsetting and 3.7:1 stretching. The pre

liminary heat treatment was performed in the pearlite temperature range (750'C). To

perform different quality heat treatments, e.g. different austenitising temperatures, the

forging was cut into disks of 130 mm thickness. The disks w ere sepa rately heat treated

at an austenitising tem perature of 1100*C and with a cooling rate simulating the centre

part of a rotor with a diameter of 1200 mm .

Rotor E, with a heat treatment diameter of 1150 mm, was manufactured from a 42 Mg

ESR ingot (diameter 1300 mm). After double upsetting the ingot was forged to a dia

meter of 1250 mm with a stretching ratio of 4.9:1. The preliminary heat treatment was

performed in the pearlite temperature range at 700'C for 150 h and furnace cooled

down to 200'C. Preliminary tests have shown that rotors of similar chemical composi

tion and dimensions can be surface cracked after direct cooling from the forging tem

peratures to 100'C. Due to the pearlite transformation, this can be avoided. For quality

heat treatment the rotor was austenitised at 1070'C. After double tempering at 570'C

and 690'C and mechanical testing at both rotor body and coupling ends a radial core

(0 260 mm) was taken to obtain material from the centre part of the rotor to test the

higher yield strength level. For the second tempering procedure aimed at getting the

lower yield strength level, some near-surface m aterial of this radial core w as simulation

heat treated. Using these results the second tempering procedure of the whole rotor

was performed at 715*C and finally, after mechanical testing of tangentially oriented

material, an axial core with a diameter of 350 mm was rem oved.

16



Rotor F, with a heat treatment diameter of 1200 mm, was manufactured from a 45 Mg

ingot. The conventional melting process included a com bined ΑΙ/Si deoxidation w ith the

aluminium level being reduced to 0.006%. Segregation problems were kept to a

minimum by employing a steep taper ingot mould with an height-to-diameter ratio of

approx. 1.0 (the wide end diameter of the ingot was 1720 mm). For the forging process

the forging was heated to 1200'C following upsetting 2.15:1 and stretching in different

stages to a total reduction ratio 3.35:1 [58]. The preliminary heat treatment included an

isothermal transformation to fem'te-peariite at 690°C. The quality heat treatment was

performed with an austenitising temperature of 1070'C and double tempering at 570'C

and 680'C to obtain the higher yield strength level. After tangential testing two radial

cores ( 0 132 mm) were taken to obtain material from the centre part of the rotor at that

yield strength level. The second higher tempering procedure for 600 MPa y ield strength

was the next manufacturing step.

5.7 TEST PROGRAMME

According to the investigation p rogramme shown in Fig. 35, two tempering conditions of

rotors E and F are being tested. Fig. 36 shows the sample position plan.

An extensive test programme has been performed for steels B, E and F involving tests

on material from the centre, mid-radius and surface positions at axial locations corres

ponding to the bottom, middle and top of the original ingot. The characterisation is in

terms of tensile and impact toughness properties, fracture toughness, creep and creep

rupture strength and high and low cycle fatigue, the latter incorporating dwell times of

up to 30 minutes. In addition, the effect of long-term ageing at 480'C, 600'C and 650'C

on tensile strength, FATT and fracture toughness is determined. Since there are always

uncertainties in the extrapolation of creep data, whether from isothermal, iso-stress or

Larson-Miller plots, low stress-level tests have been started, which will only result in

creep failure at very long times (Fig. 37). Strain measurements in the secondary creep

range will provide an estimate of expended creep life.

S.8 TEST RESULTS

Rotor B

The chemical composition is very homogenous over the cross-section as well as over

the length of the forging. There is no boron segregation. The m icrostructure consists of

tempered martensite (no δ femte was seen at any position). There is no material simu

lating the outer part of the rotor due to separate heat treatment of each disk to simulate

the centre of a rotor of diameter 1.2 m. The yield and tensile strength of specimens

taken from ou ter and central positions of the disks shown in Fig. 38 is homogenous, but

there is a slight increase in FATT at the centre (Fig. 39), even though analysis and mi

crostructural investigations showed no segreg ation.

All other properties were investigated on disk 7 which had a slightly higher yield

strength.

The 0.2% yield strength is relatively high especially at 600*C and 650'C in

comparison to the other rotors. The impact energy A

v

at 20*C is lower in comparison

with rotors E and F. The long-term behaviour of the FATT is to be seen in Fig. 27. The

FATT of 45'C in the as-received condition is the highest in comparison with the other

alloys, but there is no change of this value after 10,000 h at 650'C. This indicates a

very stable microstructure.

17·



Rotor E

After the first tempering the yield strength was about 800 MPa tested at different near-

surface positions (Fig. 39). Obviously the body part of the rotor has a lower yield

strength of 750 MPa at the near-centre position. Due to the lower yield strength the

toughness increases and the FATT decreases from + 55'C to + 5'C near the centre

(Fig.

40). The FATT in axial and radial positions shows about the same value. This is

due to the ESR process. The impact energy temperature behaviour of specimens from

near the centre of the rotor confirms the results of the trial melts. The microstructure is

fully-tempered martensite, only near the centre were small amounts of δ femte found

( « 1 %). The austenite grain size in the centre of the rotor is significantly smaller than

in the near-surface region (Fig. 41).

After the second tempering procedure to obtain a yield strength of about 600 MPa, the

specimens tested at the near-surface position of the rotor show good homogeneity in

yield and tensile strength as w ell as in FATT (Fig. 42). The results are shown in Fig. 39.

The specimens taken from the axial core have slightly lower yield strength and even

better FATT values. The 0.2% yield strength and toughness values are also consistent

with the results of trial melts. It was surprising that the FATT in the near-centre position

is better than near the surface. Similar behaviour was also found in rotors manufac

tured in Japan [59, 60]. Fig. 39 shows excellent FATT and ambient temperature tough

ness values. The 0.2% yield strength at 600'C and 650'C is higher than for conven

tional rotor material (12% CrMoV steel according to G erman standard SEW 555).

Rotor F

After the first tempering procedure to obtain a yield strength of about 730 MPa, an in

crease from the bottom to the top of the ingot was measured at the near-surface posi

tion of the rotor body (Fig. 43). This has only a small influence on toughness level.

However there is a decrease of impact energy and increase of FATT at the near-centre

position,

Fig. 39. The microstructure of near-surface and near-centre positions of the

rotor is tempered martensite with small amounts (- 0.5%) of δ ferrite at the core and a

more uniform grain size than for Rotor E (Fig. 44).

After the second tempering procedure to ob tain about 600 MPa yield strength, the yield

strength and tensile strength show more homogeneity throughout the rotor (Fig. 45).

The impact energy at 20"C, FATT and upper shelf energy are much better than for

conven tional rotors.

5.9 CONCLUSION

The m anufacture of rotors Β, E and F was successfully performed. They show excellent

properties especially with good toughness even at high yield strength. There is a large

improvement in both yield strength behaviour at elevated temperature and creep

strength compared to conventional 12% CrMoV steels. Long-term embrittlement at

temperatures up to 600*C is negligible.

Yield strength values for specimens taken from the rotors of the new steels Β, E and F,

shown in Fig. 46, indicate superior values at 500 to 600'C, particularly for the B steel,

even for similar values of the yield strength at ambient temperature. Creep results for

steel B (YS = 670 M Pa), steel E (YS = 630 MPa and - 745 MPa) and steel F (YS =

610 MPa) are plotted in Figs. 47 to 50 and show a major improvement compared to the

conventional steel X20CrMoV 12 1.

- 1 8 -



NEW FERRITIC-MARTENSITIC CAST STE ELS


At the outset of the work and as a basis for evaluation of the literature, the desired

property profile of the cast steel to be identified and tested under the C OST programme

was defined as follows:

1. A 100,000 h creep rupture strength of ca. 100 MPa at 60 0'C .

2.

Good castability and weldability.

3. Through-hardening capability up to about 500 mm wall thickness

4. Properties such as fracture toughness , low-cycle fatigue strength and long-

term toughnes s corresponding at least to those of the ferritic cast steels cur

rently used up to 565*C.

6.2 LITERATURE SURVEY

Initial development work within the COST Programme was concentrated in the early

80's on increasing the creep strength of the traditional 12% CrMoV cast steel

(G-X 22 CrMoV 12 1) by the addition of 0.025% boron at carbon contents of 0.10 and

0.20% and niobium contents of 0.04% and 0.08%. The test melts revealed inadequate

toughness and weldability [61]. Inadequate toughness and strength properties were

also established under the screening programme of EPRI project RP 1403-15 when

adding 0.0075% and 0.014% boron to steels which had approximately the chemical

composition of the m odified 9 Cr1 Mo and the new TAF steels [62].

The best profile in the screening programme of the EPRI project was established for a

cast steel melt which in its chemical composition was largely equivalent to the modified

9 CrIMo pipe and forging steel (T91/P91/F91) developed in the USA. The 100,000 h

creep strength for this pipe steel, according to the latest publication by Oak Ridge Na

tional Laboratory [36], is roughly equivalent to a value of about 90 to 95 MPa at 600'C.

The high creep strength of this steel is basically attributable to the relatively stable

M23C6 carbides and to very small and finely distributed niobium-vanadium carbonitride

precipitation. Fig. 51 shows the 100,000 h creep strength for this steel versus the test

temperature in comparison with the traditional cast steels 1 % CrMoV (GS-17 CrMoV

5 11) and 12% CrMoV (G-X 22 CrMoV 12 1) used in Europe for temperatures up to

565'C.

Generally there is a clear superiority of the newly developed wrought and cast steels

across the full temperature range from roughly 450'C to 600"C. Under the EPRI project

an investigation was also made to establish the suitability of this cast steel for thick-

walled components based on a 5-ton stepped block with wall thickness of 150, 300 and

500 mm and a 5-ton high-pressure valve chest. The castability and weldability were

found to be roughly equivalent to that of the 1% CrMoV cast steel (GS-17 CrMoV 5 11)

frequently used in the manufacture of steam turbines. Through-hardening capability

under conditions of accelerated air cooling is guaranteed for a cross section of at least

500 mm and the fracture tough ness is distinctly higher than that of the 1 2% CrMoV cast

steel (G-X 22 CrMoV 12 1). Based on the results of creep tests at 600"C up to 40,000 h

it can be estimated that the creep strength of this cast steel is roughly comparable in

the long term with that of the pipe and forging steel mod. 9 Cr1

Mo.

These results have

also been confirmed by a joint UK programme [64], Comparable results have also been

established in tests on cast steel valves and boiler pipes [65, 66].

- 1 9 -



Cast steels d iffering som ewhat from the chemical composition of the mod. 9 CrMo steel

were tested in Japan in the 80's using laboratory melts, pilot and production castings

with unit weights up to abou t 19 tons [67, 68].

The chemical compositions of these melts and compared to those of the EPRI pro

gramme RP 1403-15 are shown at the top of Fig. 52 (compositions 1 to 3). Compared

to the mod. 9 CrIMo version (No. 1) the first Japanese version (No. 2) is distinguished

by higher C, Cr, Ni and Nb contents, whereas the second Japanese version (No. 3) is

mainly characterised by somewhat lower Mo and Nb contents. There is also a diffe

rence in the tempering temperature. While a tempering temperature of 738*C is speci

fied for the mod. 9 CrIMo, the Japanese versions are only tempered at 675'C and

710'C respectively in order to reduce the coarsening of the M23C6 carbides and to

obtain a higher initial strength.

Based on the experience in connection with the development of the mod. 9 Cr1 Mo pipe

steel, the influence of the chemical composition on the properties of the 9-10% CrMoV

steels is characterised according to the sum total of the elements C + N and a specific

Cr equivalent according to Fig. 53 [69].

A distinction is made between four regions (A to D) differing in creep strength, ductility

and toughness (FATT50). Region Β features the best property profile which is

achievable with a martensitic microstructure free of δ ferrite. It is characterised by

acceptable creep strength, high ductility and high toughness (low FATT50). Located in

the middle of this region are the mean value analyses of the mod. 9 Cr1 Mo pipe steel

(P91) and the cast steel melts investigated under EPRI programme RP 1403-15 (No. 1

of Fig. 52). The two cast steels (No. 2 and No. 3) developed in Japan have a distinctly

lower Cr equivalent, but are still in or on the border of the preferred region B.

6.3 DEVELOPMENT OF CAST STEELS IN COST

501

ROUND 2

6.3.1 Pre-Evaluation Programme

Based on the results of earlier tests and the evaluation of the literature the procedure

was to seek solutions which, while maintaining the good property profile of the mod.

9 CrIMo cast steel, still allow an increase in creep strength at 600'C. On the basis of

the chemical composition of the mod. 9 Cr1 Mo steel there were two optional proce

dures:

1.

Optimising the heat treatment;

2.

Adding tungsten to increase solid solution strengthening on the lines of the

results determined in creep tests on specimens of the TAF steel [70] and in

the rotor programme of the COST 501 Round 2 project [30].

Melts No. 4 and 5 of Fig. 52 were chosen for testing. Compared with the target values

of mod. 9 Cr1 Mo according to ASTM 21 3, Grade 91, the Cr content was increased from

8.0-9.5% to 10.0 - 10.5% to improve the solubility of nitrogen and to avoid surface

porosity. In addition, the nickel content of m ax. 0.40% was increased to about 0.85% to

obtain a ferrite-free martensitic microstructure and to further improve the through-

hardening capability of thick-walled cast components. For the addition of tungsten a

value of roughly

1%

was chosen (test melt No. 5). In the rotor programme of COST 501

this value had been found to be the optimum amount in respect to increasing the creep

strength and maintaining the fracture toughness and long-term toughness which, in

fact, is also to be expected on account of the position of the test melts (No. 4 and

No. 5) in Fig. 53. The selected chemical compositions feature a Cr equivalent and a

- 2 0 -



content of C+N such that the preferred range Β is achieved. The heat treatment ver

sions listed in Fig. 54 were chosen. For comparison the heat treatment instructions for

the piping steel mod. 9 CM Mo are also listed. For the new heat treatment versions the

main characteristics are:

• Austenitising temperature 110O'C instead of 1 0 4 0 Ό

• Pretempering 24 h at 550'C for stabilisation of the microstructure on the lines of

the procedure for the newly developed ferritic 10% CrMoVNbN rotor steels

• Tempering 730* C instead of min. 73 2'C

• Post-weld heat treatment 720'C instead of min. 732*C

Overaging for 200 h at 700'C to test the long-term stability of the materials and

obtain additional information for extrapolating the creep strength values (200 hours

at 700'C corresponds to 270,000 h at 600'C based on a C-parameter of 25 for the

Larson-Miller extrapolation).

• Long heat treatment periods to match the heat treatment instructions for large

thick-walled castings.

Plates of size 800 χ 400 χ 100 mm were cast to check these test param eters. After pre

liminary tests the heat treatment versions B, BO, C and C O, according to Fig. 54, were

chosen from the pre-evaluation tests in order to perform the long-term tests shown in

Fig. 55. The test programme included the development of a weld filler metal with 1 %

tungsten for manual welding, which in chemical composition was equivalent to the trial

melt no. 5 of Fig. 52. A suitable manual metal arc electrode for the tungsten-free ver

sion was already developed under EPRI Programme RP 1403-15 [62].

6.3.2 Results of Pre-Evaluation Programm e

Strength and Toughness Properties of Base Material

Fig.

56 provides an overview of the 0.2% proof strength at room temperature deter

mined in tensile tests. In contrast with the pipe steel mod. 9 CrIMo (Grade 91) for

which ASTM 213 specifies a 0.2% proof strength of min. 415 MPa, a value of at least

550 MPa for heat treatment conditions Β and C (Fig. 54) was the target for the COST

programme. This minimum requirement was satisfied by all test versions, i.e. also by

specimens in the overaged condition. The minimum requirements were also satisfied by

all other test results.

Fig.

57 shows the notch impact energy values determined for the Charpy V-notch

specimens at room temperature. For all heat treatment conditions the tungsten-free

version is characterised by a higher notch impact energy than the version alloyed with

1%

tungsten. However, the toughness of the tungsten-based version is also distinctly

higher than the minimum impact energy specified by D IN 17 240 for the

1%

CrMoV and

12%

CrMoV cast steels (GS-17 CrMoV 5 11 and G-X 22 CrMoV 12 1) traditionally used

in the manufacture of steam turbines. For the different heat treatment conditions the

specimens of the tungsten-free version tempered at 550'C showed better results.

There is no significant heat treatment influence for the tungsten-containing version. A

noteworthy fact is that the overaging treatment (200 h at 700'C) does not distinctly

affect the toughness.

21 -



The results of the exposure tests carried out on the Charpy V-notched specimens at

480, 600 and 650'C up to 10,000 h are shown for heat treatment condition Β in Fig. 58

which again plots the notch impact energy determined at room temperature. The

tungsten-free melt features a more pronounced reduction in toughness than the

tungsten-containing melt. For both materials the maximum reduction is at 600*C. Com

parable results were also established in long-term exposure tests of the mod. 9

Cr1

Mo

pipe steel [71 ] and the m od. 9 Cr1 Mo cast steel [62].

Mechanical Properties of Weld Metal

The mechanical properties of the weld metal of the manual metal arc electrode deve

loped for welding the tungsten-containing cast steel are shown in Fig. 59. After identify

ing a suitable chem ical composition, tests were m ade to determine the influence of

dif

ferent heat inputs, inter-run temperatures and post-weld heat treatments at 720 and

7 3 0 Ό After heat treatment at 72 0'C the impact energy of 33 Joule is relatively low.

After 12 hours of heat treatment at 730"C the values determined were roughly equiva

lent to those of the parent metal. The properties determined during the tensile test

dif

fered from those of the parent metal. As expected, a distinctly higher 0.2% proof

strength and tensile strength and, accordingly, also lower ductility values were esta

blished independent of the welding and heat treatment parameters be ing tested.

Creep Tests

At the end of 1991 the creep tests of the pre-evaluation tests at 600 and 650'C had

reached about 33,000 h. Fig. 60 provides general details on the creep rupture strength

(Larson - Miller diagram). The mean-value curve of the mod. 9 CrIMo pipe steel [63]

provides a basis for comparison. The tungsten-containing version (No. 5) shows a

higher creep strength over the full test period compared with the tungsten-free version

and in the long term shows a slight superiority over the mod. 9

Cr1

Mo pipe steel for all

four heat treatment conditions.

For the tungsten-free version (no. 4) the heat treatment condition C with the 5 50 'C pre

tempering treatment features the highest creep strength whereas heat treatment condi

tion Β features the lowest creep strength. For all heat treatment conditions the

tungsten-free version generally reflects a flatter pattern compared with the pipe steel

which has a similar chem ical composition. This is probably attributable to the longer

heat treatment period required for castings which generally results in increased and

coarser carbide precipitation.

6.4 COMPONENT PROGRAMME

The good behaviour of the tungsten-containing test melt (No. 5) in the pre-evaluation

programme resulted in this composition being chosen for the component programme

under which the components shown in Fig. 61 were selected. The best heat treatment

chosen was that designated "C" in Fig. 54, i.e. 1100'C for 12 h / forced air cooling /

550*C for 24 h / cooling in still air, and, after manufacturing welding, a further 12-hour

heat treatment at 730"C with furnace cooling. In the same way a 12-hour heat treat

ment at 730'C was chosen as straightforward post-weld heat treatment for the welding

procedure qualification tests. In addition, a completely new heat treatment according to

the above sequence C has also been carried out.

22·



The valve chest is undergoing tests according to the test matrix of Fig. 62. A total of 5

different specimen positions have been chosen - one rim and two core zones as well as

two manufacturing weld zones. For the 100 mm thick welded plates the test matrix

roughly agrees with that of Fig. 62. The valve chest and the welded plates were manu

factured in 1991 according to Fig. 61. Manufacture was not accompanied by any

diffi

culties neither in respect of the casting and welding process nor of the heat treatment.

A further positive result was obtained in the nondestructive tests. The num ber and size

of the detected flaws were similar to those found in 1 % CrMoV cast steel (GS-17

CrMoV 5 11) traditionally used in the manufacture of steam turbines. Fig. 63 provides

details of the mechanical properties established by testing cast-on coupons. In com

parison with the results of the heat treatment C of the pre-evaluation programme the

0.2% proof strength is somewhat lower, i.e. 573 MPa. Conversely, the impact energy

values lie between 44 and 60 Joule, i.e. correspondingly higher.

6.5 RESULTS OF THE COMPONENT PROGRAMME

The tests shown in Fig. 62 are continuing. The short-term tests have been completed.

In general they confirm the promising results from the preliminary programme. Fig. 64

provides a summary of the mechanical properties determined at positions A to E of the

valve body.

The specimen positions A and Β represent the thickest wall sections, whereas the

specimen position C shows the profile of properties in the thin-walled support. The

properties of m anufacturing welds are determined at the positions D and E.

The strength, ductility and toughness values for the positions Α, Β and C agree well

with the mechanical properties determined on the cast-on specimens after a post-weld

heat treatment of 12 hours at 730'C (Fig. 63). Only the impact energy, at 30 Joules,

lies somewhat lower. The impact-energy transition temperature (FATT50) lies between

45 and 63Ό

The mechanical properties of the manufacturing welds also satisfy the requirements,

even though the proof strength lies 100 MPa higher than for the preliminary tests

(compare Fig. 59).

The curtent status of the creep tests, with a testing time of 30,000 hours is shown in

Fig.

65. Specimens from positions A to E are being tested at 550, 600 and 65 0'C . The

results agree well with those of the preliminary programme, both for the base material

and for the manufacturing welds. Tests on welds in plate material of 100 mm in thick

ness provide concurring results, for test durations up to 1100 hours at 550 and 6 00 'C .

Considerable progress has also been made in the low cycle fatigue tests. Fig. 66

shows the results from the surface-near position A of the valve body. In the low cycle

range (< 2000 cycles) the results at 550 and 600'C are better than those at room

tem

perature. At greater numbers of cycles to failure the results at 55 0'C are the lowest and

are even inferior to those determined for the conventional 12% CrMoV steel, X 22

CrMoV 12 1, at 530"C [72]. It will be of particular interest to see whether this tempera

ture dependence is confirmed for specimens taken from the other positions, and also

for tests performed with hold times of 20 minutes in tension and/or com pression .

The long-term exposure tests have now attained 10,000 hours. Mechanical tests of

these specimens are still to be performed.

-23



6.6 CONCLUSIONS

The cast steel G-X 12 CrMoWVNbN 10 11 developed and tested within the COST 501

Round 2 P rogramme ap pears to have a creep strength which is at least as high as that

of the modified 9% CrMo steel P91. The castability, weldability and suitability for non

destructive testing are similar to that of the 1% CrMoV steel GS-17 CrMoV 5 11

con

ventionally used for turbine m anufacturing. The toug hness properties are also similar to

those of this cast steel. The long-term tests will be continued within the framework of

the third round of COST

501,

WP 11, in order to determine the service-relevant proper

ties for plant operating at 600'C.

7. CREEP-RESISTANT BOLTING MATERIAL


A major task of the

COST501

-2,

WP3 programme has also been to identify suitable

materials for the bolts required when adopting improved ferritic 9 -12% CrMoV steels

for the turbine casings, valve chests and piping exposed to temperature up to roughly

600'C.

Generally, the high-temperature bolt materials will have to satisfy the following re

quirements:

high stress relaxation resistance for intervals between overhauls of at least

50,000 h

reasonable agreement between the thermal expansion coefficient of the bolt,

flange, elastic sleeves and nut materials

no notch sensitivity under creep loading, i.e. goo d creep deform ation beh aviour

high fracture toughness at the temperatures encountered during assembly and

turbine operation

high yield strength to prevent permanent bolt elongation due to preloading and due

to steady state and transient thermal stresses

resistance to stress corrosion cracking over the full range between ambient and

steady-state operating temperature

a relatively stable microstructure to prevent any unacceptable reduction in the

properties of the material during sustained op eration.

Fig. 67 provides a quantitative appraisal of the compatibility of the materials of the

bolted joints with respect to their expansion coefficients [73]. The lowest thermal ex

pansion coefficient is shown by the 9 -12% Cr steels which, particularly for turbines

operating at 600"C, are suitable for manufacturing the casings and flanges [74, 75].

The value for the superalloy Nim 80A, frequently used in the past for the manufacture

of high-temperature bolts, agrees well with that of the 1% CrMoV steel.

Good long-term operating experience has been obtained, for instance, with the bolt

materials of DIN 17 240. Fig. 68 shows the residual stresses after 10,000 hours for the

temperature range between 400 and 65 0'C .

The ferritic steels

X 19 CrMoVNbN 11 1

X 22 CrMoV 12 1

21 CrMoV 5 7

- 2 4 -



material,

long-term rupture and relaxation strengths of the material had to be esta-

blished after the new treatment.

7.3 WORK PLAN OF BOLTING PROGRAMME COST 501

7.3.1 New high streng th ferritic steels

• Static 1,000 h stress-relaxation tests were carried out using plain bar samples, to

rank the candidate materials. These screening tests were performed at 600'C with

an initial applied strain of 0.15%. Similar tests on a X 19 CrMoVNbN 11 1 steel

were used as a b asis for com parison. Only those ferritic steels which exhibit better

relaxation properties than the X 19 CrMoVNbN 11 1 steel i.e. > 120 MPa after

1,000 h at 600"C were co nsidered for further evaluation.

• Long-term stress relaxation tests (up to 10,000 h) are being conducted at 550 and

600"

C on plain b ar samples from materials chosen in Phase 1. Complementary

tests are also being carried out using model bolted assemblies which assess the

behaviour of the complete bolt/flange combination. The material for this Phase 2

were taken from the rotor programme.

7.3.2 High Purity Nim 80A

Procurement of a "high purity" Nim 80A with a uniform fine grain size of ASTM

5 - 6 .

Standard 3 stage heat treatment and modified 3 stage heat treatment to coarsen

N¡3 (A l, Ti) p recipitates.

Stress relaxation tests on plain bar samples at 550 and 600'C with an initial strain

of 0.15% up to 10,000 h.

Stress relaxation tests with model bolted joints (bolt and nuts Nim 80A, flange

modified 9% CrMo cast steel up to 10,000 h at 540, 570 and 6 00 'C.

Long-term creep rupture tests with plain and notched specimens at 550 and 600 'C.

Constant strain rate stress corrosion cracking tests in 4% H2SO4 solution at

90' C

and a strain rate of 1.2 χ 10"

6

/sec.

constant load stress corrosion tests in 4% H2 SO4 solution up to 10,000 h at 90'C .

Long-term ageing and embrittlement tests at 450 to 700'C.

The chemical compositions of the 15 melts investigated in the 1000 h stress relaxation

screening tests at 600'C are given in Fig. 69. The newly developed steels mod. 9%

CrMo (T91), the boron-alloyed TAF steel introduced by T. Fujita, the experimental

"rotor" steels B1, D1, D3, E1, E2 and F1, and the nitrogen-containing test melts D135,

D191,

D93 and DE259 (see Chap. 5) were investigated, in comparison with the bolting

steel X 19 CrMoVNbN 11 1 (DIN 17 240) traditionally used at lower temperatures. The

mechanical properties of the test materials are given, along with the heat-treatment

data,

in Figs. 70 and 7 1 .

The chemical composition, heat-treatment data and mechanical properties of the

high-

purity Nim 80A alloy are shown in Figs. 72 and 73. The desired low levels of the trace

- 2 6 -



elements were achieved. Of particular importance is the fact that the phosphorus con

tent is below 20 ppm . In comparison with the standard heat treatment the mod ified pro

cedure results in a reduction in strength and an increase in ductility.

7.4 RELAXATION TESTS

7.4.1 1000 h Screening Tests at 60 0'C of High Creep Strength Ferritic Steels

7.4.1.1 Plain Bar Tests

Fig.

74 provides an overview of the test results obtained with a prestrain of 0.15%. For

a tempering temperature of 700'C the test melts E1 and F1 show the same relaxation

strength as the DIN steel X 19 CrMoVNbN 11 1. All other steels have a relaxation

strength which is as much as 3 5% lower. Clearly none of the test melts reach the target

relaxation strength of at least 120 MPa after 1000 h.

7.4.1.2 M odel Bolted Assembly Tests

Fig. 75 shows the model chosen for this investigation. The materials of the nuts match

those of the bolts. The flanges are generally made of the modified 9% CrMo steel, or

exceptionally of the TAF s teel. In accordance with DIN 17 240, the initial strain was s e

lected as 0.2%. The majority of the tests were performed at 600'C. Since the results lay

well below the residual stress target value of 120 MPa, additional tests w ere carried out

at 570 and 540'C. The results are compiled in Fig. 76. The values determined at 600'C

largely agree with those of the uniaxial stress relaxation tests (see Section 7.4.5.1.1 /

Fig.

69).

The nitrogen-alloyed version D 135 (similar to rotor steel A3) responded slightly better

than the DIN steel X19 CrMoVN bN 1 1 1 . The slight superiority was also found at 570

and 540'C. At 570'C the target value of 120 MPa is marginally exceeded by the nitro

gen-alloyed version D135 and DE259 and at 540'C by X 19 CrMoVNbN 11 1, the TAF

steel and the nitrogen-alloyed version DE259. For comparative purposes Fig. 76 also

shows results of the EPRI Programme 1403-15 (73B 1), with the material comb ination:

bolt - Nim 80A

nuts - Nim 80A

flange - Mod. 9% CrMo

However in this comparison it must be bome in mind that the prestrain of 0.2% applied

at room temperature is already reduced to a value of about 0 .1% at the relaxation tem

perature, as a result of the large difference in coefficient of thermal expa nsion between

the flange and bolting materials. Fig. 77 demonstrates the method of stress determina

tion using the bolt prestressing diagram, when different materials are used in a bolted

joint [73].

7.4.2 Long-term Relaxation Tests

The long-term tests up to 10 000 h have been concentrated on the investigation of Nim

80A, both as plain bar and model bolted assembly tests. The ferritic steels DE254

(nitrogen alloyed), TAF and B1 (boron alloyed), E1

(1%

W and 1% Mo alloyed) and

27·



mod. 9% CrMo (P91) were also subjected to long-term testing. A further test variant in

volved the use of a 0.25% prestrain of the Nim 80A model, in order to partially compen

sate for the difference in coefficients of thermal expansion between the flange and

bolting materials. The selection of a prestrain of 0.2% leads to calculated initial elastic

strains of ca. 0.15% at 600 'C, ca. 0.16% at 570 'C and ca. 0.17% at 540'C .

For these test temperatures Figs. 78 to 80 show the residual stresses determined as a

function of the testing time. Fig. 81 summarises the 10,000 h relaxation strength values

as a function of the testing temperature.

The results obtained can be summ arised as follows:

• The newly developed highly creep resistant steels do not exhibit a higher relaxation

strength than the currently em ployed DIN steel X 19 CrMoVNbN 1 1 1 .

• In the tests with ferritic bolting steels at long testing times the same relaxa tion

strength is obtained using plain bars or bolted joint models (Fig. 80).

• All tests show that the bolt material Nim 80A exhibits a higher relaxation strength

than the ferritic bolting steels.

• Even for similar initial strains, tests with the Nim 80A bolted joint models always

lead to lower relaxation strengths than are measured using plain bar uniaxial tests

in the relaxation test machine. The difference is relatively minor at 540 and 5 70 'C.

However at 600"C the difference is about 50% . The different behaviour can already

be observed in the first few hours of the test, and is a result of plastic deformation

in the thread and contact surfaces of the nuts and of primary creep processes and

stress redistribution in the more highly stressed regions of the bolted joint

• The relaxation streng th of the Nim 80A bolted joint model can be significantly im

proved by increasing the prestrain at room temperature from 0.2 to 0.25%.

• In comparison with the standard heat treatment, the modified heat treatment of Nim

80A always leads to a reduction of the relaxation strength in uniaxial relaxation

tests.

In the less highly stressed model tests the modified heat treatment only leads

to a reduced relaxation strength in the tests at 600'C, and only after about 4000 h.

It is most likely that the lower relaxation strength is a consequence of the coarser γ'

microstructure resulting from the furnace cooling from 1080 to 8 5 0 Ό

7.5 CREEP RUPTURE TESTS OF HIGH-PURITY NIM 80A

The Figs. 82 to 85 shows the results of the creep tests performed at 550 and 600'C

with smooth and notched (ak = 4.3) specimens, in comparison with the creep strength

scatterband according to DIN 17 240. The maximum testing time attained is about

18,000 h. In agreement with the uniaxial relaxation tests, the modified heat treatment

leads to a slight reduction in the creep strength, whereby the creep strength of the

standard heat-treated material lies somewhat above the DIN scatterband and the creep

strength of material with the modified heat treatment lies in the upper range of the DIN

scatterband. The ductility values of the two heat-treatment variants is also different at

550'C. The notched specimens with the standard heat treatment show a temporary

creep notch embrittlement (notch weakening) at times below 10,000 h, whereas for the

modified heat treatment the smooth and notched specimens show about the same

times to failure at 550'C. On the other hand the decrease of the creep fracture strain

28



and reduction of area values shows the opposite tendency. This occurs somewhat

earlier for the modified heat treatment. A significant temporary reduction of these

ductility values can also be seen at 6 00 'C for the modified heat treatment.

7.6 EMBRITTLEMENT TESTS FOR HIGH-PURITY NIM 80A

Fig.

86 gives a summary of the Charpy impact energy values determined for high-purity

Nim 80A after 10,000 h exposure at 600"C, both for the standard and modified heat-

treatment s chedules. Also shown within the scatterband are results of earlier tests per

formed by National Power (UK) [77] performed with melts having different impurity

levels and also results of an unsuccessful heat treatment, which resulted in a coarse

grain size (VRM ). For the modified heat treatment condition (V12) there is practically no

change in the impact energy (ca. 48 Joules) as a result of the exposure at 600'C.

However for the standard heat treatment condition (V12N) the impact energy is basi

cally lower (37 to 38 J) and long-term exposure reduces the values to about 30 Jo ules.

For the unsuccessful heat treatment the impact energy in the initial cond ition is very low

(10 Joules). However in long-term tests the impact energy rises to values which are

comparable with those of the modified heat treatment. In summary, the long-term expo

sure tests show that the em brittlement tenden cy, which was often noted in the past for

melts with normal impurity levels, can be prevented by the use of melts with reduced

impurity levels.

7.7 STRESS CORR OSION INVESTIGATIONS OF HIGH-PURITY NIM 80A AT 90 'C IN

4% H

2

S O

á

7.7.1 Constant Strain Rate Tests

Fig.

87 shows a summary of the results in terms of the variation in the reduction of area

at fracture as a function of the initial impact energy, based on tests performed by

National Power with melts of different impurity levels and with different heat treatments

[77]. The results confirm the earlier observation that for Nim 80A a higher co rrosion re

sistance is correlated with a higher value of the impact energy. Results for the earlier,

incorrect, heat treatment exhibit inferior toughness and corrosion resistance and lie at

the bottom of the scatterband for the previous tests.

7.7.2 Constant Load Corrosion Tests

The stress corrosion cracking behaviour under constant load was investigated both in

the standard heat-treatment condition and after exposure for 1,000 h at 600'C. The

specimens were loaded to 120% of the yield strength values. The testing times of from

4,000 to 12,000 h on a total of 5 specimens are summarised in Fig. 88. Microscopic in

vestigation after removal of the specimens showed no cracking in any case. This posi

tive observation agrees well with the results of the constant strain rate tests, in which

the high-purity melts were seen to posses a clearly improved resistance to attack by

stress corrosion cracking.

.8 CONCLUSIONS FROM THE BOLTING TEST PROGRAMME

• The newly-developed high creep strength 9 to 12%Cr steels exhibit the same or

poorer stress relaxation strength than the conventional ferritic material

- 2 9 -



Χ 19 CrMoVNbN 11 1 in the investigated temperature range of 540 to 60 0'C .

However the newly-developed steels have a higher toughness in the starting

con

dition and , as shown in the investigations of the rotor steels, a relatively high duc

tility in creep tests, bo th for smooth and notched specimens.

The residual stress of about 150 MPa, required to ensure long-term sealing of tur

bine components operating under internal pressure, cannot be ensured for the 9 to

12%Cr steels when operating at temperatures above 550'C. Hence for this range

the only solution lies in the use of highly creep resistance nickel-based a lloys, such

as Nim 80A, in conjunction with casings, valves and pipe flanges of ferritic steels.

The relaxation tests performed show that Nim 80A is a suitable material for bolts

along with ferritic flange steels. However a disadvantage lies in the higher coeffi

cient of thermal expansion of Nim 80A. This disadvantage can be largely elimi

nated by increasing the initial strain at room temperature by about 25%, i.e. from

0.2 to 0.25%.

A significant improvement in both the toughness values and resistance to stress

corrosion cracking can be ach ieved through reducing the level of impurities in Nim

80A.

In addition the use of a modified heat treatment leads to an improvement of the

stress corrosion properties and of the ductility of Nim 80A during creep testing. The

modified heat treatment requires furnace cooling from 1080 to 850"C rather than

air cooling from 1080Ό This results in coarser y-precipitates and a consequent

minor reduction in the creep and stress-relaxation resistance in comparison with

the standard heat treatment.

8. IMPROVED STEELS FOR STEAM PIPES AND HEADERS

8.1 INTRODUCTION

The increase in steam temperature and pressure requires the use of materials with

better creep strength than that of the well known ferritic-martensitic steel

X20CrMoV12 1 (referred to as X20) for main steam pipes and headers. This steel has

been widely used in Europe since its development in Germany in the 60's and has pro

vided good service. Whilst X20 possesses single-phase martensitic microstructure

which permits the specification of this steel also for use in thick pipe sections, the creep

strength reduces drastically at temperatures above 550'C. Additionally X20 has limited

weldability as a result of the high carbon content of 0.17 - 0.23%, so that after welding

thick sections cannot be cooled to ambient temperature before post weld heat treat

ment (PWHT) without risk of cracking. Nevertheless it continues to be specified by

European boiler makers for temperatures up to 565"C. Higher temperatures and pres

sures require greater wall thicknesses and thereby loss of flexibility in the pipework

system,

higher thermally-induced stresses during transient operation and difficulties

with PWHT on site.

Austenitic steels have considerably higher creep strength than X20 and they have been

used successfully in Germany for thick-walled components since the 50's. However

these power plant operated in baseload so that there were very few start-stop and load

cycles per year. The low yield strength and thermal conductivity and high coefficient of

thermal expansion of austenitic steels, in comparison with ferritic steels result in an in

creased susceptibility to low cycle fatigue in thick-walled components. Fatigue failures

of thick-walled austenitic components have occurred in the USA and Britain. Future

power plant will be required to operate with frequent stop-start and load-change cycles,

- 3 0 -



so that austenitic steels are inappropriate for heavy components. However they are

suitable for the manufacture of thinner-walled superheater tubes in high-temperature

plant, as a result of their superior corrosion resistance.

There is a need for ferritic steels with improved creep strength at temperatures up to

and beyond 60 0'C.

8.2 THE DEVELOPMENT OF 9% CHROMIUM STEELS

The basic 9%Cr 1 %Mo steel was developed in 1936, for the petrochemical industry, to

provide a steel quality with increased corrosion resistance compared with the 2.25%Cr

1%Mo steel [82]. The stable microstructure of this 9%Cr 1%Mo steel was noted to pro

vide satisfactory mechanical properties and creep rupture strengths for service tem

peratures up to 550'C and was adopted from the late 60's onwards for the British nu

clear power programme.

To improve the creep behaviour of the basic 9%Cr 1%Mo steel small amounts of

niobium and vanadium were added together with an increase in molybdenum to 2%.

This resulted in the development of steel grade EM12 [83]. This steel grade, which

possessed a duplex microstructure containing up to 40% δ ferrite, replaced the 300

series stainless steel grades in French fossil fuel power stations in superheater and re

heat superheaters for steel temperatures up to 6 2 0 Ό However, as a result of the high

delta ferrite content and the precipitation of Laves phases during service at 550"C this

steel grade became embrittled and displayed low creep rupture ductilities. The poor im

pact properties of steel EM12 further prevented the specification of this steel grade for

use in heavy, thick-walled com ponents

In 1974, the United States department of Energy established a task force to select

materials suitable for the Liquid Metal Fast Breeder Reactor programme and Clinch

River project. Following the recommendations of the task forced the Oak Ridge

National Laboratories, in conjunction with Combustion Engineering Inc., were commis

sioned to develop a modified 9%Cr 1 %Mo steel that did not suffer the detrimental em

brittlement processes experienced by EM12 whilst still retaining the attractive benefits

of a ferritic microstructure. The development programme, involving approximately 100

test heats employing various melting practices and the conversion of ingots to bar,

plate and tube sections permitted the cha racterisation and eva luation of each test hea t,

resulting in the selection of a com position now designated as Steel 91 [84].

Following the 1977 international conference on ferritic steels for fast breeder reactor

steam generators, at which the improved strength and toughness of the modified 9%Cr

steel were reported, the development of this steel grade was broadened to involve a

number of American industrial companies and the production of semi-commercial scale

(40 tons) heats. Ingot conversion to plate, pipe, tube and forgings allowed product

forms to be characterised and evaluated, mainly at the Oak Ridge National Labora

tories, to provide property values subsequently accepted by ASTM and ASME. During

1983, grade T91 (T = tube) was approved in ASTM standard A213 and became com

mercially available for pressure tube application. The ASTM approved grade P91

Ρ = pipe) in 1984 (standard A355) for piping and header applications.

The ASTM approval in 1983 of grade T91 was followed by national code recognition in

France (NF A49213 and NF A49219). In ISO/DIS 9392-2 "Seamless steel tubes for

pressure purposes - Technical delivery conditions" steel 91 is included as steel grade

X 10 CrMoVNb 9 1. It will also be standardised in the new European Standard which is

concurrentiy under preparation. In both cases no distinction shall be made between

pipe and tube. The creep strength values and permissible stresses are significantly

-31 -



higher than for X20 at temperatures above 550'C. The standardisation of P91 in the

USA led to the initiation of a number of research projects in Japan and Europe,

whereby the manufacturing, mechanical properties and application as thick-walled

components were primarily investigated. These programmes showed that thick-walled

pipes, pipe bends and welds could be produced in P91 more easily and with better

properties than X 20.

Fig.

89 shows the Time-Temperature-Transformation (TTT) diagrams for X20 and P91

[85]. There is a remarkable similarity between the two diagrams, but also some charac

teristic differences. The martensite start temperature lies about 100'K higher than for

X20 and the martensite hardness is about 150 Vickers units lower.

Both effects result from the difference in carbon content. The lower martensite hard

ness is of practical importance since it leads to simplifications in the manufacturing and

processing of the steel. For example, the danger of intercrystalline stress corrosion

cracking is greatly reduced in the hardened condition after hot bending or welding. The

danger of cold cracking after welding is also reduced so that components can be

cooled directly to room temperature.

After normalising, the tempering treatment at temperatures typically between 730 and

780'C leads to the precipitation of M23C6 chromium carbides at the boundaries of the

martensite laths. In addition fine V/Nb carbonitrides of type MX also appear. They serve

to stabilise the microstructure and further increase the strength. The optimum distribu

tion and size of these particles is controlled by the V/Nb ratio.

8.3 BRIEF COMPA RISON OF MECHA NICAL PROPE RTIES

The following figures indicate the typical relative mechanical properties possessed by

Steel 91 , grade EM12 and grade X20. Test results were obtained from samples ma

chined from 51 X 10 mm tube [86].

Fig.

90 presents a comparison of room temperature yield and tensile strengths in which

the direct influence of microstructure is readily apparen t. G rade EM12 with - 20 - 40%

ferrite displays the lowest strength whilst X20 (higher carbon) displays the highest

strength. The impact toughness properties shown in Fig. 91 reveal a significant

superiority of Steel 91 compared to both EM12 and X20. The Steel 91 displayed higher

absorbed energies and much lower FATT. Fig. 92 presents the elevated temperature

properties which, while reflecting the previously noted trend in room tempe rature tensile

properties, shows a diminishing variation with increasing temperature. The creep rup

ture strength at 100,000 hours as a function of temperature is shown for selected data

for each grade of steel in Fig. 93.

8.4 PROCESSING

As Steel 91 was initially deve loped for nuclear applications, considerable emphasis was

placed on melting practices, argon oxygen-decarburisation techniques and electroslag

refining processes to ensure compliance with strict composition limits. These steel

making techniques, coupled with careful scrap selection procedures to ensure very low

contaminant levels (e.g. Cu, Ρ, As, Sn) and the application of improved ingot teeming

practices, to restrict segregation effects, permitted the production of commercial

tonnage heats with satisfactory composition values as required in the standard codes.

32-



Hot

Working

Ingots should be reheated into the 1200 - 1250'C temperature range to ensure the

complete solution of all (niobium rich) carbo-nitrides, after which forging, rolling, con-

ventional pipe or tube making processes can be completed in the temperature range

1050- 1200"C to ensure adequate hot ductility and good recrystallisation. Higher hot

working temperatures are not recommended as laboratory testing has demonstrated a

significant ductility - dip for temperatures above 125 0'C.

Whilst thick sections should be permitted to cool slowly after hot working operations to

avoid hydrogen defects (flakes) such slow cooling con ditions may induce the precipita-

tion of coarse carbides. To avoid this possibility, a double-quenching heat treatment

(1200'C WQ + 1070'C WQ) has been successfully applied to thick, hot-rolled plate with

thicknesses above 100 mm [87].

Cold Working

Steel 91 possesses excellent cold formability allowing conventional cold rolling or

drawing of tubes (8 - 185 mm O.D.) and the cold bending of tubes for service require-

ments within the deformation ratios adapted to Steel 91. For cold-formed bends, a de-

crease in creep strength is observed due to the accelerated déstabilisation of the dis-

location structures in the cold-formed martensite [88].

Heat Treatment

For 9 - 12%Cr ferritic-martensitic steels, such as Steel 9 1 , the optimum combination of

metallurgical and mechanical properties is achieved by a normalising and tempering

heat treatment. The highest strength at room temperature and at elevated temperatures

is derived from the combination of a normalising treatment of 1100'C followed by a

tempering treatment at 750'C [89]. The optimum normalising temperature at 1100'C

results in hardness values, after air cooling, of 400 - 420HV indicating that most of the

carbon is in solution as this hardness corresponds to that of a 0 .1% carbon martensite.

Higher normalising temperatures are not presently recommended due to the austenite

grain coarsening observed above 1100'C and the appearance of δ ferrite (e.g. 5% fol-

lowing 1 hour normalising at 1200'C). Both these changes are generally considered

unacceptable for pressure vessel steels because of their effects on ductility and tough-

ness as well as an implied heterogeneity of properties.

The useful tempering range for steel 91 appears to be restricted to 750 - 800'C. The

lowest practical tempering temperature of 750'C provides a high strength level espe-

cially in conjunction with a normalising temperature of 1100'C. The maximum limit of

800'C is due to retransformation and the consequent re-hardening of a fully tempered

microstructure.

Weldability

Steel 91 has been reported by many research centres as possessing excellent weld-

ability properties, mainly as a result of the lower carbon content (0.1%) which provides

the Steel 91 with greater resistance to cold cracking (maximum hardness after welding

is approximately 450 HV for Steel 91 compared to 600 HV for Steel X20). In general,

experience has sho wn that Steel 91 is less critical for welding that the 12%Cr steel X20.

The lower HAZ hardness for Steel 91 weldments further permits a lower preheat

tem-

perature (150'C - 250'C depending upon thickness) and the welds may be allowed to

- 3 3 -



cool to room temperature after welding before applying the post weld heat treatment,

which should be in the range 730 - 760'C in order to obtain the optimum compromise

between toughness and creep properties. Lower post weld heat treatment tempera

tures resulted in a decrease in impact toughness (although no values were reported

below 50J). Higher temperatures resulted in improved toughness but were not con

sidered appropriate in terms of optimum creep strength. No reheat cracking was ob

served for the Steel 9 1 .

Apart from welding "standard" heat treated steel 91 base material the welding of "half

tempered" steel 91 tubes have been reported by both American and Japanese re

searchers as an effective means for avo iding the drop in creep strength in the intercriti-

cal, fine grained HAZ (Type IV zone). The "half tempering" treatment requires re-

austenitisation at 1050 - 1100'C followed by air cooling after which tempering is per

formed at 620 - 650'C instead of 730 - 780'C before welding. Whilst such a technique

cannot be applied to field weldments, it may offer opportunities for critical components

(e.g.

branch connections) welded in the workshop [90]. The beneficial effect of half

tempering is a result of incomplete precipitation of particles at the lower tempering tem

perature so that they cannot overage and coarsen during welding. Additionally, their

precipitation during the post weld heat treatment stabilises the matrix and improves the

creep strength of the m aterial.

The main conclusions from a recently reported weldability study [90] of Steel 91 tubes

and pipes stated:

• Steel 91 grade material is easily weldable with low preheat temperatures

150'C - 25 0'C depending upon thickness.

• Steel 91 weldments may be allowed to cool to room temperature prior to the appli

cation of post weld heat treatment techniques.

• Girth welds generally satisfied the ASME IX requirements concerning room

tem

perature properties.

• Toughne ss levels above 50J are easily obtainable with post weld heat treatments of

730 - 760'C for reasonable short times.

• To obtain good weld metal toughness and to avoid hot cracking, the welding heat

input should be stringently limited (at least with the European filler metals studied

during this investigation).

• As in other ferritic steel weldments, Steel 91 weldments contain a softer region in

the intercritica , fine-grained HAZ (type IV zone). T he relative crosswe id creep

strength loss of about 20% at 600'C is smaller than the loss observed for X20

weldments.

• "Half temp ering" provided an effective technique in limiting the creep strength loss

of Steel 91 weldments in the intercritical, fine grained HAZ.

8.5 SHORT-TERM PROPERTIES

Physical

The three most important physical properties of steels selected for power plant service

are the thermal expansion coefficient, thermal conductivity and modulus of elasticity.

Ferritic steels, such as Steel 91 possess more favourable values for these properties

- 3 4 -



compared to austenitic steels as these physical properties are dependent upon com

position and crystal structure only.

In the range of intended service (550 - 650*C) the coefficient of thermal expansion of

Steel 91 is approximately 30% less than that of comparable stainless steels. Within the

same temperature range, the thermal conductivity of Steel 91 is approximately 30%

greater and the decrease in the modulus of elasticity much sm aller compared to that for

austenitic steels. Fig. 94 to 96 present the variation of these properties with tempera

ture.

As a consequence of the above physical property comparisons, Steel 91 compo

nents will expand less during temperature increase to service conditions and be si

gnificantly less affected by changing load cond itions that may otherwise induce thermal

fatigue cracking.

Mechanical

Mechanical properties are sensitive to microstructure modifications and thus variations

in both normalising and tempering temperature and time have been reported to affect

the tensile properties'and creep strength of Steel 91 . Further, variations in composition,

within the specified ranges, may also affect the mechanical properties of Steel 91 .

Tensile

The influence of normalising and tempering treatments has been reported on the basis

of both laboratory testing and component manufacture. Whilst the strength of samples

normalised from 1050'C was found to be dependent on the tempering temperature,

after normalising at 1100"C there was no effect of tempering temperatures, above

750*0 However, the time at the tempering temperature was observed to influence

tensile properties, longer times providing lower strength values. This influence of time is

important for thick section components, for which relatively long treatment times may be

specified.

Impact

The impact property data published for a wide range of products, with section thickness

up to 300 mm, show Steel 91 to possess good toughness properties in the'normalised

and tempe red condition at both room temperature and - 2 0 Ό The FATT is typically

around 10'C for tubes and -20*C for pipes [91].

Ageing at temperatures of 480 to 600'C results in a significant loss in toughness and a

rapid increase in the transition temperature, the maximum effect being observed after

25,000 hours ageing. At higher ageing temperatures of 650 - 700"C there was little

change in the transition temperatures, but significant softening at room temperature,

and large increases in the upper-shelf ene rgy.

Whilst the appearance of Laves phase had previously been identified as being respon

sible for embrittlement after age ing standard 9Cr -1 Mo steel, the observation that frac

ture occurred via transgranular cleavage (rather than intergranular) indicates the Laves

phase in Steel 91 plays little or no role in the selection of the crack path. The sugges

tion that Laves phase particles aid crack initiation has not been supported by experi

mental evidence. Thus whilst the appearance of the Laves phase corresponds to the

observation of embrittlement, the exact mechanism for the embrittlement has not been

identified [92].

35



8.6 PROGRAMMES ON IMPROVED CREEP-RESISTANT STEELS FOR STEAM TUBES.

PIPES AND HEADERS

The COST programme was designed to determine the effect of fabrication steps, such

as cold bending a nd different welding techniques (both similar and dissimilar welds) on

the mechanical and specifically creep properties of Steel 91. Limited comparative

testing was also performed for X20 Material, was obtained from different suppliers and

with different wall thickness. Additional header-section components of both X20 and

P91 were ma nufactured by a powd er-metallurgical (PM) route and sectioned for testing,

in order to determined dimensional accuracy and the uniformity of mechanical and

creep properties, and compare them with the properties of conve ntionally manufactured

material.

8.6.1 Investigation of Tube and Pipe

For tubing, gas tungsten arc welds (GTAW) were made between

T91-T91,

T91-X20,

T91 2 4Cr1Mo an d T91-TP347H steels. Welds between T91 tubes were made using

modified as well as unmodified welding wire. All welds in these tubes were cooled to

ambient temperature prior to heat treatment. Tubes were also cold bent, using different

bending radii. Some of the bends were heat treated after cold bending.

For application in headers and steam piping shielded metal arc welds (SMAW) and

submerged arc welds (SAW) were made in thick-walled, large-diameter pipes.

In order to assess the degree of change in the properties due to welding and cold

forming, the programme also includes tests on the unprocessed base material. To get

some indication of the dispersion of the properties of T91/P91, a limited number of tests

was performed on base materials from 4 other suppliers.

For the prediction of the long-term behaviour of actual components fabricated from

T91/P91 on the basis of the test results, a comparison with the properties of a pre

viously existing high-temperature steel was considered useful. For this purpose, the

German 12%Cr steels X20 was chosen. Besides the creep tests, which formed the

major part of this research project, all base and processed materials were also sub

jected to extensive mechanical and microstructural investigations. The scope of the in

vestigation programme is summarised in Fig. 97. The chemical compositions of the in

vestigated base metals are presented in Fig. 98 and those of the welding consumables

in Fig. 99.

The assessment of the long-term creep behaviour was based on ¡sostress creep

testing, performed at 100 MPa and temperatures between 680 and 580Ό The tests

have been linearly extrapolated on a log time (h) versus temperature ('C) basis. Accor

ding to the figures derived from the ASME Code an average rupture time of 100,000 h

should be achieved at 98 MPa and 600'C. For tubes and pipes the creep specimens

were taken in the longitudinal direction. The specimens were uniaxially loaded at a

constant

load,

corresponding with a nominal stress of 100 MPa. The specimens for

testing the weldments were taken in the same direction. Thus, the specimens taken

from circumferential weldments were loaded perpendicular to the weld, so-called

"cross-weld specimens". The creep tests on bends were performed, using special ring-

shaped specimens cut from the bends which were loaded in the circumferential direc

tion Fig. 100. This means also that the intrados and the extrados are loaded equally,

since the cross sections were equal. The advantage of this method is that the weakest

zone of the bend can be detected (extrados or intrados). Furthermore, this facilitates a

comparison between test results and operational loading conditions (highest principal

stress is in circumferential direction). A disadvantages of this type of specimen is, how

ever, the presence of an additional bending stress which leads to higher creep

-36-



deformations on the inner surface, resulting in an overall lower creep performance with

respect to the results from axial specimens. The tests on ring-type specimens were only

performed to compare the results of bends with straight tubes and those results may

not be interpreted as absolute values. Numerical creep analysis performed by the

Netherlands Energy Research Foundation (ECN) showed a maximum deviation in rup

ture times by a factor of 2. The special test set-up for loading of the ring-sh aped sp eci

mens is shown in Fig. 100.

Base

metal

The results of the creep test on base materials from different suppliers are presented in

Fig.

101.

The data of most materials lie in a rather narrow scatter band. Rupture times

at 600'C and 100 MPa are about one fourth of the expected time derived from the

ASME Code. Only the material delivered by supplier C approaches the value derived

from ASME probably due to its higher carbon and nitrogen contents. The rupture time

of X20 base metal appears to be one fourth of the rupture time of

T91/P91.

In Fig. 102 the minimum creep rates for the base metals T91/P91 and X20 are pre

sented.

The creep rate of material C appears to be three times lower than that of the

other suppliers, and the creep rate of X20 is approximately six times higher. Since the

slope for X20 is steeper, the difference between X20 and T91/P91 decreases at lower

temperatures and disappears at about 540'C.

Weldments

The isostress creep tests on T91/P91 weldments have been restricted to the materials

of supplier A All welds were circumferential. The result of the creep tests are presented

in Fig. 103. The 10

5

h creep strength of the T91/P91 cross-weld specimens at 600'C is

approximately 15% lower than the creep strength of the base materials. This decrease

is more pronounced in GTAW weldments than in SAW and SMAW weldments. Since

both base metals have the same creep rupture strength, it is expected that the diffe

rence in creep strength between these weldments is caused by differences in heat in

put during welding. The SMAW weldment was made with the lowest heat input and the

GTAW weldment with the highest heat input (10

5

h creep strength at 600 'C, 77 and 73

MPa respectively). In all specimens rupture occurred in the fine-grained part of the HAZ

(near the base metal). This location corresponds with the weakest zone of the weld

ment.

Even the rupture location in the creep specimen of a T91 weldment made with unmodi

fied wire is in the HAZ, demonstrating that the shorter rupture time of the weldment in

relation to the base metal is not caused by the lower creep strength of the weld metal,

but by the properties of the base metal due to the heat input during welding. This phe

nomenon of rupture in the fine-grained HAZ is often referred to as Type-IV cracking.

For X20 the 10

5

h creep strength at 600"C for weldments is also considerably lower

than that of base metal. The limited number of tests suggests a decrease of

2 1 %

(from

62 to 49 M Pa).

The T91-TP347H dissimilar joint, welded with nickel base consumable, fractured in the

HAZ of T91 in approximately the same time as the T91-T91 weldments. The dissimilar

joint T91-X20 fractured in the fine-grained HAZ on the X20 side, the rupture time being

half the rupture time of T91-T91 weldments. The dissimilar joint T91-2V4Cr1Mo, w elded

with modified wire, fractured in the soft decarburised zone near the fusion line on the

2 Cr1

Mo side an d not in the Type-IV zone nea r the base metal.

- 3 7 -



Bends

The results of the creep tests on ring-sh aped specimens of cold-formed bends are pre

sented in Fig. 104. For comparison the results of the creep tests on the base metal of

the same tubes, using axial specimens as well as ring-type specimens, are also in

cluded.

The rupture times for the ring-type specimens of straight tube are somewhat shorter

than those for the axial specimens. All bends fractured in the extrados and the creep

rupture times were shorter than the rupture times of the axial and ring-type specimens

of the straight tube. The creep rupture time of the bends decreases with decreasing

bending radii (increasing plastic deformation), but this effect was not pronounced. A

post bend heat treatment (PBHT) of 740 'C/1h hardly improves the creep rupture time at

60 0'C . The slope of the curve is, however, steeper compared with that of bends without

PBHT. This suggests that PBHT may be useful for service temperatures < 5 8 0 Ό

The 105 h creep rupture strength at 600"C for the cold-formed bend w ith a radius of 60

mm (R/D = 1.3) is ca. 13% lower than that for the base metal (73 and 84 MPa respec

tively). After a PBHT of 740*C/1h the 10

5

creep rupture strength of the bend increases

on slightly (from 73 to 76 MPa).

The decrease of 12% (from 52 to 46 MPa) in 10

5

h rupture strength for the X20 bend is

comparable with that of the T91 bend. Such a drop is therefore not a typical T91 phe

nomenon.

Fig. 105 shows the results of the creep tests, extrapolated to 10

5

h [93].

The Larson-Miller values (with C = 30, a value which is realistic for the base metal) from

the extrapolated results for 600'C at 100 MPa are presented in the second column.

From these Larson-Miller values the creep rupture strengths for 10

5

h at 600*C are de

rived by extrapolation. These figures are presented in the third column. From these

figures the ratio of the creep rupture strength between the p rocessed material and base

metal has been derived, as presented in the last column.

The following conclusions may be drawn:

• The 10^ h creep strength at

600*

C of weldments performed on c ross-weld speci

mens appeared to be about 15% lower than that of the base metal. This drop is

caused by the behaviour of the fine-grained HAZ, being a typical phenomenon for

all ferritic and martensitic steels (type IV cracking).

• The

IO

5

*

h creep strength at 600'C of cold-formed tube bends with low R/D ratio

appeared to be about 10% lower than that of the base metal. The same holds for

cold-formed X20 bends.

• The workability of T91/P91 for welding and cold bending is very good (less critical

than X20). This aspect and the high creep strength make the use of this steel in

power generation components operating in the creep range very attractive. It is,

therefore, to be expected that T91/P91 will gradually take the place of X20 in those

applications.

A programme performed by Mannesmann (partially within the COST project), was di

rected towards an examination of the long-term creep properties of P91, measured by

different laboratories and for different sources of material, and the effect on the creep

properties of the pipe bending and w elding processes.

Fig.

106 shows a direct comparison between the creep strength of the steels X20 and

P91.

Since there is partial overlap between the two scatterbands only selected melts

have been included in the diagram. In order to maintain a reasonable basis for com-

- 3 8 -



parison four recent melts of the X20 steel were selected. Whereas at 550*C the values

coincide with the standard, those at 600 and 650'C lie somewhat above the mean va

lues, but still within the scatterband. The curves for the P91 data represent the mean

values for the overall evaluation which will be presented later.

In the temperature range considered the 10

5

h creep rupture values for the P91 lie

above those for X20. The relative difference increases with increasing temperature. In

short-term tests X20, with the higher hot yield strength, shows superior values, so that

there is an intersection of the isothermal creep curves at intermediate testing times

(between 100 and 1,000 h).

Properties of inductive hot bends and welds were investigated in the subse quent part of

the programme. As for X20, renewed quenching and tempering after bending is also

required for P91 in order to improve the creep properties. Fig. 107 shows that after

quenching and tempering the long-term creep results for specimens taken from an in

ductive bend in a P91 pipe are similar to those of untreated material. There is also no

difference between specimens taken from the inside and outside of the bend and no

dependence on the stressing direction (tangential or axial). The minor differences bet

ween undeformed pipe and pipe bend are due to differences in the heat treatment.

These differences are reduced at longer times and higher temperatures.

Earlier work [94] concentrated on the investigation of welds made in P91 using the

SMAW technique and concluded that creep strength was such that the material could

prove competitive with austenitic steels at 600 'C.

In the COST programme weldability was investigated for the SMAW, GTAW and SAW

techniques. Although welding consumables have not yet been fully optimised, there are

a number of candidates already commercially available. Toughness values of these

weld metals are generally higher than for X20. Hardness profiles of typical welds in X2 0

and P91 are illustrated in Fig. 108.

In both cases a hardness minimum occurs in the "intercritical zone". Tests on weld

simulation specimens show that this is also the zone with the lowest creep resistance

[95]. A change in failure location in cross-weld creep specimens is noted, as shown in

Fig. 109. Whereas short-term failures appear in the base material, the cracking location

is shifted at longer times into the intercritical part of the heat affected zone, with a

corresponding reduction of the creep strength in comparison with the base material.

The transition in failure location is both temperature a nd stress de pende nt. The effect is

particularly p ronounced at 600 'C and is qualitatively similar for the X20 and P91 steels.

A specific project in Denmark investigated the weldability of a thick-walled pipe (outer

diameter 353 mm, thickness 63 mm) of P91 from Sumitomo Metal Industries [96]. The

work included simulation of the HAZ microstructure in a Gleeble simulator with subse

quent microstructural examination, welding with the SMAW technique and ISO-stress

creep testing. It was co ncluded that:

• the improvement in creep strength of P91 results from the precipitation of V/Nb

carbonitrides

• weldability of P91 is similar to that of X20

• unavoidable softening of the HAZ locally reduces the creep strength

• the reduction in creep strength for cross-weld specimens is about 20 %

Investigation of long-term creep properties of P91 base material has been the subject

of a number of recent publications. Research programmes have tested samples

- 3 9 -



obtained from both laboratory test material and commercial products or components.

Studies concerning the effects of compositional variations heat treatment and micro-

structure, ageing and material thickness have been presented at several international

conferences and published widely in the technical journals concerned with high tem

perature steel properties.

The effect of variations in composition of P91, both within and outside the current

specified ranges was recently reported [97] to demonstrate the important influence of

the V/N ratio on creep rupture strength values measured at 6 0 0 *0 Fig. 110 reproduces

the results presented by Orr and Di Gianfrancesco [97] to indicate the optimum creep

rupture strength at 600'C to occur at a V/N ratio of about 4. Fig. 111 presents a

schematic representation of the compositional factors influencing the creep rupture

strength of

P91.

Significantly, these results were obtained from test material normalised

at 1040 -10 60 'C and tempered at 720, 750 or 760 'C.

The effect of composition variations, within the ASTM specification for P91, was also

reported [98] to demonstrate the importance of vanadium content during a study con

cerned with the properties of thick and thin-section P91 material. The important contri

bution of precipitation hardening was evidenced during this study when the slower

cooling experienced by thick sections resulted in coarsening of precipitates and a con

sequent reduction in high temperature strength. Whilst higher austenitising tempera

tures cause greater carbide dissolution and therefore would be expected to produce a

stronger solid solution / precipitation strengthening response, increased austenitising

temperatures (from 1040 to 1150*C) provided only a small increase in strength for thick

sections compared with a relatively large increase observed in thin sections. The large

increase in strength observed for thin section material was noted to be associated with

a large drop in ductility and the appearance of intergranular creep cavitation such that

higher austenitising temperature treatments are not recommended for thin section

material.

The influence of both austenitising and tempering temperature and time on the micro-

structure and mechanical properties of P91 was reported by Orr, Burton and Rasche

[89]. Results obtained from laboratory experiments and commercial heat treatments re

vealed the effect, of various tempering treatments on strength to also depend on the

initial normalising treatment. Fig. 112 describes the influence of various normalising and

tempering treatments observed during this study of the stress rupture strength of

samples taken from a 40 tonne cast. To optimise the tempered martensitic microstruc

ture, in which the small niobium and vanadium carbides/nitrides were noted to be

largely responsible for the high strength of Steel 91 at both ambient and elevated tem

peratures, a normalising temperature of 1100'C followed by tempering at 750'C for as

short a time as possibie was recommended.

Determination of Long-term Creep Data

Long-term values of the creep strength of P91 are required in order to guarantee ser

vice reliability. Although one specimen at ORNL has already exceeded 100,000 h at

538'C, the longest testing times in Europe still only exceed 50,000 h [85, 99]. A first

evaluation of the data was performed by ORNL [100] and served as a basis for the

ASME Code Case 1943. In 1991 a new analysis was performed by Mannesmann

Research Centre (MFI) including also results from ORNL, Sulzer Bros, and Vallourec

effectively somewhat reducing the creep strength values given by ORNL. A further

evaluation was performed by MFI in 1992 including the newest Japanese results on a

further 33 melts. This raised the mean values to a position intermediate between those

of the original high estimate by ORNL a nd the sub sequent lower results of the first MFI

evaluation (see Fig. 113).

- 4 0 -



High Temperature Corrosion Resistance

For successful long-term operation in power plant at elevated temperature materials re

quire not only a high creep strength but also a high resistance to corrosion both in

steam and in hot exhaust gas from the combustion process. Observations of the T91

tube thickness loss of 0.2 to 0.3 mm after 7 years of operation under severe conditions

are reported [101]. Steam oxidation may potentially have two negative effects:

• the growth of oxide on hea ted boiler tubes will reduce the conduction of heat

through the tube wall, thereby raising the tube temperature and reducing its creep

life

• growth and subseq uent spalling of oxide causes hard particles to pass through the

steam turbine, potentially causing erosion damage.

X20 is known to provide good long-term service at current turbine-inlet steam tempera

tures up to 565"C. Since the corrosion resistance normally decreases with decreasing

Chromium content [102], steam oxidation tests of X20, P91 and 2%Cr1Mo piping steels

have recently been carried out [103]. The results show that in steam at 550 to 650*C

creep resistant steels w ith 9 - 1 2 % Cr are almost immune to oxidation. The 12% Cr

steel (X20) exhibits anomalous behaviour in that the corrosion rate reduces with in

creasing temperature, due to the formation of more protective oxide layers. The 9%Cr

steels (P91) show normal behaviour, with a two-phase oxide layer, comprising an inner

layer of iron-chromium oxide and an outer layer of magnetite. Extrapolation to 10

5

h

predicts that even at 650"C wall thickness will not be reduced by more than 0.3 mm, a

value which should be technically acceptable since no spalling is expected below this

thickness. This p rediction is in reasonable agreement with the observed oxide thickness

of 0.09 mm for a P91 pipe operated in steam for 30,000 h at a metal temperature of

610"C [85 ]. The same pipe showed a scale thickness of 0.06 mm on the fireside.

Practical Application of P91 (pipes and tubes)

P91 is now in use throughout the world. American utilities and the CEGB have been

using P91 since 1980 at temperatures of 593 to 6 2 0 Ό It has often been used to re

place P22 (2

1

4Cr1Mo, 10C rMo9 10) for thick-walled headers. Pipes in Drakelow C (GB)

were replaced by P91 in 1991 and the new power plant Kawagoe, in Japan, built in

1989, made complete use of P91 for the main steam piping.

In Europe test sections of P91 are incorporated in the steam lines of Esbjerg 2 (DK)

and a Preussen Elektra plant (D). A recent case in which P91 has been installed in the

live steam line of a German power plant is particularly well documented, including de

tails of m anufacturing and qualification weld testing [104].

Fig.

114 shows that replacement of X20 by P91 can lead to a major weight (and cost)

reduction. In this case a T-piece of P91 for operation at 585'C is over 60% lighter than

the corresponding part manufactured form X20, as a result of the difference in creep

strength.

The most advanced steam power plant currently under construction in Japan (Matsuura

No.

2, 1,000 MW, 593 / 593'C ) uses P91 main steam pipes.

¡.6.2 Pow der-M etallurgically Manufacture d Header Sections

The objectives of this part of the CO ST programme were to

- 4 1 -



• Asse ss the suitability of the powder metallurgy (PM) man ufacturing route with re

spect to bo th production and sen/ice requirements of full-sized boiler heade rs.

• Manufacture a section of a live steam header using net shape pow der metallurgy

(PM), including the nipples and nozzles required for the connection of superheater

tubes and pipework.

• Determine the mechanical properties of PM-produced X20 and P91 materials.

Basically components with complex geometries can be manufactured with conventional

techniques, such as machining from solid material or welding together sub-assemblies

or alternatively with the help of powder-metallurgy techniques. The latter results in very

uniform properties through avoiding segregation in thick-walled components and

through avoiding welds at highly loaded points. For complex parts the PM technique

can result in lowered m anufacturing costs. The technique comprises:

• me lting and ladle metallurgy, whereb y melting is earned out by inductive heating of

the ladle and the alloying elements are added subsequently. The chemical compo

sition of the steel is adjusted by vacuum treatment, stirring and temp erature control

in the ladle.

• spray ing of the melt horizontally into inert gas. The melt enters the spray chamber

directly through an opening in the base of the ladle and is powderised by the inert

gas. The powder is of high purity with a low oxygen content.

• encapsu lation of the powder in sheet steel capsules with a form as similar as

possible to the shape of the final component. The capsules are filled with the

powder and subsequently evacuated and sealed.

• hot isostatic pressing (HIP) comprises subjecting the capsules to high pressure and

temperature, resulting in a 100% densified workpiece with the required mechanical

properties.

• heat treatment and final machining are performed subsequen tly.

The steam header, which collects the steam from the individual superheater tubes, is

one of the most highly loaded components in the steam power plant. Headers are nor

mally manufactured from thick-walled pipe and T-pieces, by drilling holes into the pipe

and w elding on connecting pieces for later welding to the smaller diameter su perheater

and steam pipes and tubes on site. A critical point in such a welded component is the

welded joint between the header and the connecting pipe. This location experiences

high thermal and mechanical loads and bending moments exerted by the pipework

system.

The aim of the COST programme was to seek an alternative procedure to the welded

design mentioned above. Boilermakers, steelmakers and research institutes partici

pated [105 - 108]. A PM-header manufactured during the programme is shown in

Fig. 115 (outer diameter 295 mm, wall thickness 50 mm, axial length 680 mm). The

steps in this programme included:

• manufacture of three PM header sections

• complete nondestructive inspection of the test material

• mecha nical and metallurgical investigation of the material from the headers

• compa rison of the properties of PM and conventional material.

42-



8.7 NEWEST DEVELOPM ENTS FOR TUBE AND PIPE STEELS

A comprehensive review of the newly developed high temperature ferritic-martensitic

steels for tube and piping applications from USA, Japan and Europe was presented re

cently at the VGB Conference, in Kolding [111]. Fig. 121 illustrates the development of

9 -12% Cr steels with an indication of the actual or expected approximate 10^ h creep

strength at 600*0 The chemical compositions of these steels are shown in Fig. 122,

and their estimated 10

5

h creep rupture strengths in Fig. 123. Creep values for the W-

alloyed steels are given only as an approximate range.

Steel Development

A major difficulty for the development of high creep strength 9 - 12% Cr steels is the δ-

ferrite balance. High content of the ferritie formers Cr, Mo, W, V and Nb is beneficial for

creep strength, and low content of the austenite formers C and Ν is necessary for

weldability and toughness However following these trends will introduce high amounts

of δ-ferrite in the steels. Balancing additions of e.g. Ni and Mn, to ensure fully martensi

tic microstructure, lower the austenite transformation temperature, and makes it difficult

or impossible to temper the steels to acceptable strength and toughness levels. This is

of special importance after welding.

X20CrMoV12 1

Although manufacturing, welding and fabrication require special precautions, the steel

X20 CrMoV 12 1 has been intensively used in many large power station boilers in

Europe and worldwide. 30 years of excellent service experience have demonstrated the

applicability of the material [112, 113].

Having a creep rupture strength of 128 MPa at 550"C and 59 MPa at 600'C for

100,000 h according to DIN 17.175, X20 CrMoV12 1 was a considerable improvement

compared with 2VSCr1Mo steel. The X20 material allowed construction of the first coal-

fired power plants with supercritical steam parameters.

P/T91 and HCM9M

In the 70's a number of new ferritic-martensitic steels were developed in Japan and

USA. The basis for these developments was the well known 9% Cr1%Mo steel T9,

mainly used as hydrogen resistant steel in chemical plants and refineries. Its creep

strength is similar to 2%Cr1Mo. Improvements of the high temperature creep strength

were achieved by adding V, Nb and Ν or doubling the concentration of Mo.

Steels resulting from these developments were, among others, Sumitomo 9%Cr 2%Mo

HCM9M and the 9% Cr 1%MoVNbN P/T91 developed by Oak Ridge National Labora

tories,

ORNL, in co-operation with Combustion Engineering. While HCM 9M had a creep

strength similar to X20, P91 possessed a long-term creep strength ca. 50% higher than

X20 at 600'C. In 1983 and 1985 steel grade 91 was approved by ASTM/ASME as

material for power plant superheater tubes (T91), pipes (P91) and forgings (F91). With

the commissioning in 1989 and 1990 of two 700 MW units at power plant Kawagoe,

with steam data 310 bar and 566'C, the Japanese demonstrated for the first time the

use of P91 in new constructed plants with elevated steam parameters [114]. Present

construction of two 400 MW coal-fired power plants by Elsam has shown that with P91

- 4 4 -



used for superheater headers and steam lines it is possible to design for steam data of

290barand580*C[115 ] .

New W-alloyed steels

After the development of P91 Japanese developments concentrated on the effect of W-

additions to 9 - 12% Cr steel [116 - 118]. W is added at levels up to 2% to the steels

NF616, HCM12A and TB12M producing a clear improvement of creep strength com

pared with P91. W acts by solid solution strengthening of the matrix. However, some

uncertainty still exists as to whether the strengthening effect of W can be sustained

during prolonged service exposure. Investigations have shown that part of the W will

precipitate as intermetallic Laves phase in these steels after relatively short service

times [119]. Whether this has a significant effect on strength or ductility, has to be clari

fied by long-term testing. The two new ferritic-martensitic high temperature tube steels

TB9 and TB12 were introduced as a result of nearly 35 years of development work.

Sumitomo produced a 1 2% Cr tube steel HCM12 [120].

TB9 (later called NF616) is a 9% Cr <4%Mo2%WVNbN steel. The decrease of Mo and

addition of W is expected to double the creep strength compared with X20 or to exceed

by 40% the creep strength of P9 1. TB91 is a modified version of TB 9, based on a

12%Cr matrix, whereas HCM12 is a 12%Cr1Mo1%WVNbN tube steel containing app.

30 %

δ-ferrite. The development is based on principles rather similar to those for TB9

and TB12. Creep rupture tests with durations up to 80,000 h demonstrate a creep

strength slightly better than T91 [121].

Modifications of these three new 9 -12% Cr steels named NF616 (TB9), TB12M (TB12)

and HCM12A (HCM12) have been under development since 1989 for thick-section

application, in the EPRI RP1403-50 project, where Japanese, European and US steel

makers, boiler manufacturers and utilities are participating [122]. Thick-section pipes of

all three steels have been produced and a large testing programme is under way.

In the current round of COST 501, a new version of a W-alloyed pipe steel, designated

E911,

will be tested. Trial melts have been produced and tubes and pipes manufac

tured by European steelmakers.

It is expected that the new W-alloyed steels will have creep rupture strengths above

120 MPa at 600'C and 10

5

h [123]. Preliminary studies by Elsam indicate that this will

allow the construction of thick-section components, for example, for a 400 MW unit with

steam data 350 bar and 60 0'C.

Materials for high-temperature thick-section boiler componen ts and steam lines require:

• High creep strength at elevated temperatures

• High yield and tensile strength

• Good ductility and high impact strength

• Weldability and fabricability

The present status of documentation of these properties is given below, specifically for

the new W-alloyed steels.

Since NF616 was considered to be an appropriate material for HP piping, Elsam or

dered in 1988 from Nippon Steel a test piece of a steam pipe. The steam pipe material

was used for a project in which material investigations and welding trials should pro-

- 4 5 -



duce sufficient data for a plant test of the material in the HP steam line of the supercriti

cal 250 bar, 560"C, 385 MW power plant Vestkraft unit 3. The pipe had nominal di

mensions: D outer diameter 352 mm, wall thickness 56 mm.

The investigations verified the microstructural and mechanical properties, based on

short-term tests. The weldability was good and the microstructural and mechanical

properties of the weld were ac ceptable [124 ]. On this basis the Danish National Labour

Inspections Authorities accepted the plant test. In 1992 Vestkraft unit 3 was com

missioned with the NF616 steam pipe test piece installed in the X20 high pressure

steam line. The wall thickness of the test piece at the transition weld NF616-X20 was

the minimum wall thickness for X20, but in the middle of the test piece the wall thick

ness was reduced by 30% in order to increase the service stresses. A butt weld was

placed in this area. P osition and dimension of the test piece are show n in Fig. 124.

In 1989 National Power (UK) and EPRI launched the international R&D project

RP1403-50 with the aim of developing ad vanced 9 -1 2 % Cr steels for thick-section

components of fossil-fired pow er plants. The p articipants are steelmakers: Nippon Steel

(J),

Sumitomo (J) and Forgemasters Engineering Ltd. (UK); boilermakers: ABB Com

bustion Engineering (USA), Mitsubishi Heavy Industries (J) and NEI I RD (UK) and utili

ties: Elsam/Elkraft (DK), EPRI (USA) and National Power (UK). The project is based

upon the existing Japanese 9 - 12% Cr steel developments NF616, HCM12 and TB12

and modifies the last two steels as thick-section versions named HCM12A and TB12M.

Phase two of the project, scheduled to start in late in 1993, will include full size compo

nent fabrication and plant trials.

As the development work done on the NF616 steel since 1985 has produced the

largest amount of test data and the longest testing times among the new W-alloyed

steels, a more detailed discussion is given of the properties of this ste el. N F616 m ay be

regarded as an example of the whole group of new W -alloyed steels.

Physical properties

Modulus of elasticity and coefficient of thermal expansion for NF616 are similar to the

values for P91 and X20. Thermal conductivity for NF616 is comparable with that for

P91, i.e. NF616 has the same advantageous physical properties w ith respect to use for

thick-section high temperature components as P91.

Mechanical properties

• Yield and ultimate tensile streng th, ductility and impact data

For the test pipe ordered by Elsam yield strength at RT and 600'C are about 500

MPa and 300 MPa respectively. Tensile properties are plotted against testing

temperature in Fig. 125 in comparison with P91 [1 25]. Impact data for NF616 tube

and pipe material in the as-tempered condition and after ageing for up to 3,000 h

are shown in Fig. 126. In the as-tempered condition the NF616 pipe data are

similar to those for

P91.

The drop in impact energy with increased ageing time is

a well known phenomenon for 9 -1 2 % Cr steels and the minimum value at room

temperature of about 40J/cm

2

expected for NF616 pipe material after long-term

ageing is acceptable for high temperature thick-section boiler component use .

• 4 6 ·



Creep rupture properties

Creep rupture data at 600'C, 650"C and 700'C for three different heats of NF616

tube steel with test durations up to 36,000 h and at 550, 600 and 700'C for one

heat of NF616 pipe steel with test durations up to 21,000 h are shown in

Figs. 127 and 128 [119 ]. Based on this da ta, extrapolation methods give

estimates of creep rupture strength at 600'C and 10

5

h in the range 120-140

MPa. This data must still be considered with c aution. Test durations are relatively

short, and the metallurgy of the new W-alloyed steels with respect to long-term

stability has not yet been fully demonstrated. However, the present data produce

strong evidence that the NF616 steel will have a creep rupture strength at 600'C

and 10

5

h which is clearly better than P9 1.

Test programmes by Nippon Steel [126] and by Elsam/Elkraft [124] have demon-

strated the we ldability of thick section NF616. In the Nippon Steel test programm e

GTAW and SAW was investigated, whereas the Danish programme investigated

only GTAW.

The filler material for both test programmes was produced by Nippon Steel. The

filler material composition matches the base metal with the exception of lower

carbon content and addition of Ni and/or Mn to minimise the formation of δ-fem'te.

Using a preheat temperature of about 250*C neither SAW nor GTAW showed any

problems with hot cracking or restraint cracking susceptibility. After welding and

full martensitic transformation of the microstructure during intercooling the welds

were post weld heat treated at between 740* C and 770* C for m in. 2 h. A typical

hardness distribution in a GTAW weld before and after PWHT is given in F ig. 129.

Yield and tensile strength of NF616 welded joints are similar to the base material

since the rupture position is in the base metal. Impact tests on samples with the

notch located in HAZ and in weld metal produced values at 0'C of 50J/cm

2

and

25J/cm

2

respectively [126 ]. These results are similar to those foun d for thick-sec-

tion P91 welds [127]. Short-term crossweid isostress creep rupture tests indicated

an app. 20% reduction compared with the base material creep strength. The

typical type IV cracking of ferritic steels was also observed in NF616. Long-term

crossweid creep testing is in progress. At present, these tests have reached

run-

ning times of about 15,000 h. Recent Japanese publications [128, 129] deal

specifically with the weldability and creep properties of welded joints in the W-

containing steels Nf616, HCM12 and HCM12A. The reduction in creep strength is

modest.

A comprehensive review of the properties of NF616 has recently been published

[130],

including a summary of physical and mechanical properties, with creep

testing duration of over 40,000 h at 600'C. Field tests exceed 50,000 h for tubes

and 10,000 hours for pipe. An application has been made for inclusion in the

ASTM/ASME standards during 1994 under the designation T92/P92. Based on

the available creep data and the design criteria of the ASME code, the allowable

stresses have been calculated for different temperature. These values are

illustrated in Fig. 130 forT /P2 2 (2%Cr1Mo steel), X20 , T/P91 an d T/P92.

Besides the potential of raising the thermal e fficiency of new pow er plants close to

50%,

the development of the 9 - 12% Cr ferritic-martensitic steels also offers

benefits to existing or more conventional. These new and stronger materials

make it possible to reduce thickness and weight, as shown, for example, in

Fig.

131.

■47 ·



CONCLUSIONS AND FUTURE TRENDS

A major requirement for successful high temperature service of new ferritic-martensitic

steels is that the microstructure should remain essentially stable. Long-term creep and

exposure tests were therefore an essential part of the programm e.

Careful analysis of all results obtained on trial melts permitted the identification of those

steels and heat treatments most prom ising for the production of full scale rotor forgings

and a cast valve body. Greatest weight was given to the attainment of the creep rupture

strength target as indicated by isothermal and ISO-stress rupture testing. However, the

attainment of good toughness was also taken into account. For both cast and forged

steels the creep strength advantages are 35 to 45* C.

The newly-developed forged steels do not offer sufficient improvement in relaxation

strength for bolting applications at the highest steam temperature. In combination with

the 9% CrMo steel used here as flange material, Nim 80A bolts possess a sufficient re

sidual stress at temperatures upto 600*C, provided that the initial prestrain is raised to

0.25% in order to counteract the difference in coefficient of thermal expansion between

bolting and flange material. Nim 80A was produced and tested with a modified com

position (reduced impurities) and modified 3-stage heat treatment to reduce the

ten

dency to embrittlement after long-term service and stress corrosion cracking. The pro

gramme was successful in that all targets could be met with the modified Nim 80A.

Extensive testing of T/P91 tubes and pipes produced by different manufacturers

con

siderably extended the da tabase for this steel, including also the effects of cold and hot

bending and of welding on key properties such as the creep strength. Testing of a

header section of P91 manufactured by a power-metallurgical route showed highly iso

tropic mechanical properties and a creep strength at least as high as that for conven

tionally manufactured material. The PM route therefore offers an alternative for geo

metrically complex parts.

An important achievement in 1992 has been the formulation of a programme for

Round 3. The new aims concern both an extension of the current work and also a new

orientation of the programme to more advanced steam cycles in the light of information

which has been generated during the course of the COST programme, both inside and

outside the CO ST activity.

The basic approach in the current round is:

A) consolidation:

Aimed at providing a more extensive data base for design at a live steam

temperature of up to 600 'C.

B) optimisation:

To get the most possible out of ferritic steels with the objective of

vali

dating their use at temperatures up to about 6 2 0 Ό

C) further developm ent:

Higher steam temperatures through increased use of improved austenitic

steels (£ 65 0'C)

The programme was approved by the COST501 Management Committee and began

on 1st January 1993.

48·



9. "Hekinan No. 3 Unit 700 MW: The World's largest Steam Turbine with 593'C

Steam Conditions", H. Oh-hara, et al., JSME-ASME Power Engineering Conf.

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10.

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11 .

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50-



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30.

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- 5 1 -



37 . T h e P robability of a New 12% Cr Rotor Steel Applicable for Steam Temperature

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- 5 2 -



53. "Nitrogen Alloyed Steels for High Strength and High Temperature Applications in

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5 3 -



66 . "Creep Properties of Cast T91 from a Comm ercial Casting", P. Schepp, et al.,

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67 .

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19-21,

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1978, pages 115-124.

71. "Modified 9Cr-1Mo Steel, Technical Program and Data Package for· use in

ASME Sections I and Vili", J.E. Cunningham et al., ORNL Report.

72 . "Der Einfiuss des Gefügezustandes und des Umgebungsm ediums auf das

Langzeitdehnungswechselverhalten warmfester Werkstoffe", D. Obst, FVV-For-

schungsbericht 395/1987.

73 .

"High Tempe rature Bolting for 1100*F Coal-Fired Power Plants", K.H. Mayer and

H. Koenig, International Symposium on Improved Technology for Fossil Power

Plants - New Retrofit Application, March

1

- 3, 1993, Wash ington, DC, USA.

74 . "An assessment of Alloy 80A as a high temperature bolting material for ad

vanced steam conditions", S.M. Beech, S.R. Holdworth, H.G. Mellor, D.A. Miller

and B. Nath, The international conference on advances in materials technology

for fossil power plants, 1 - 3 Sept. 1987, Chicago, IL, USA.

75 . "High temperature bolting of steam turbine for improved coal-fired power plants",

K.H. Mayer and H. König, Second international conference on improved coal-

fired power plants, 2 - 4 Nov. 1988, Palo Alto. Ca ., USA.

76. "Heat-resistant and highly heat-resistant materials for bolts and nuts", DIN 17

420,

July 1976 edition.

77. "An Assessme nt of Ni-based Alloys for High Tempe rature Bolting Applications

(an overview of COST 505 projects)", National Power Research Report,

ESTB/L/0/80/R90 June 1990.

78 . "Behaviour of Nimonic 80A Fasteners in Turbines", B. Nath, COST 505 Project

UK21.

79 . "Damage Assessment of Service Stressed Nimonic 80A Steam Turbine Bolts",

K.H. Mayer, COST 505 Project D29.

- 5 4 -



80. "Behaviour of Nickel-Base Alloys for Steam Turbine Bolting", H.H. Loser and

H. Granacher, COST 505 Project D11.

81 . "COST 501/11, WP3: High temperature bolting ma terial", B. Nath, Progress report

on UK consortium activities, National Power TEC/L/MWC/M90, 05.10.1990.

82. "Properties of the 9% Chromium Steel", H.D. Newell, Metal Progress, Feb. 1936,

p. 51.

83 . "Properties and Industrial Applications of a Superhe ater Tube Steel Containing

9%Cr2%Mo V Nb". M. Caubo, J. Mathonel, J. Revue de Metallurgy 1969, [66]

p. 345.

84. "Modified 9Cr - 1Mo Steel - Technical Programme and Data Package" ORNL

Technology Transfer Meeting Knoxville - TN , April 1992.

85. "Eigenschaften der 9 - 1 2 % Chromstähle und ihr Verhalten unter Zeitstand

beanspruchung", W. Bendick, K. Haarmann, G. Wellnitz and M. Zschau, VGB

Kraftwerkstechnik 73 (1993) 1, p. 77.

86. "Experience in Data Collection and Assessmen t for Material Standards"

J.

Orr, D. Burton, ESCS Information Day, Düsseldorf, November 1992.

87 . "Manufacturing Experience of Thick Plates for Pressure Vesse ls" P. Bocquet,

Ph. Bou rges, J. Buriat, A. Cheviet, ibid.

88.

"Workability and Long Term Properties of Modified 9% Chromium Steels"

F. Arav. OF. Etienne, H.J.M. Lentferink, J C. van Wortel, ibid.

89.

T h e Sensitivity of Microstructure and Mechanical Properties of Steel 91 to Initial

Heat Treatments" J. Orr, D. Burton, C. Rasche, ECSC Information Day, Düssel

dorf, November 1992.

90.

"Weldability and High Temperature Behaviour of the Modified 9%Cr Steel Grade

91 Tube and Pipe Base Materials and Weldments" C. Coussement, M. De Witte,

T. De Backer,

ibid.

91.

"Comparative Evaluation of Tubular Products in T91 and P91 with

X20CrMoV 1 21 Grades". J. P elabon, F. Pe llicani, ASM/EPR I Intnl Conf.

Chicago

1981,

p. 243.

92 .

T h e Effect of Long - Term Aging on the Impact Properties of Modified 9Cr - 1 Mo

Steel", D.J. Alexander, P. J. Maziasz, C. R. Brinkman, ASM Conf. October 1992.

93.

"Effect of Fabrication Process on the Creep Behaviour of 9 - 12% Chromium

Steels", F. Arav, H.J.M. Lentferink, OF. Etienne, J.C. van Wortel, Proceedings

5th Int. Conf. on Creep , Orlando, Fa, 17 - 21 May 1992.

94.

"Verhalten des 9% Chromstahles X10CrMoV Nb91 im Kurz- und Langzeit

versuch.

Teil 2: Schweissverbindung", F. Brühl, H. Cerjak, H. Müsch,

K. Niederhoff und M. Zschau, VGB Kraftwerkstechnik 69 (1989), p. 1214 - 1231 .

- 5 5 -



95. "Erweichungsverhalten der Wärm eeinflusszone des hochwärm efesten Chrom-

stahles X 10 C rM oV Nb 91 ", in Schweissen und Schneiden 41 (1990),

p. 515-5 20.

96 .

"Untersuchungen von Grundwerkstoff und Schwe issverbindung der

neu-

entwickelten Stähle NF616 und P91", R. Blum, J. Hald und E. Lund, VGB Conf.

"Werkstoffe u nd Sc hweisstechnik im Kraftwerk 1991", 9-10 Jan. 1991, Essen.

97 . "The Effect of Com positional Variations on the Properties of Steel 91 " J. Orr,

A. Di Gianfrancesco, ECSC Information Day, Düsseldorf, N ovember 1992.

98 . "Properties of thick and thin section grade 91 steel for use in conventional and

advanced coal-fired power plant." D. J. Bartow, O J. Middleton, E. Metcalfe I

Mech.

E. 1990, p. 267

99 .

"Evaluation of Design Values for Steel

91 "

W. Bendick, K. Harmann, G. Wellnitz,

ECSC Information Day, Düsseldorf, November 1992.

100.

"Modified 9 C r -1 Mo Steel for Advanced Steam Generator Applications,"

OR. Brinkmann, D.J. Alexander and P.J. Maziasz, Proc. Jt. ASME/IEEE, Power

Generation Conf., Boston, 2 1 - 2 5 Oct. 1990

101. "The long-term Field Test Results of Modifed 9Cr1Mo Stee l", T. Iwasaki,

I. Kajigaya, T. Tosuki and M. Nakashiro, JSME-ASMW COnf. Tokyio, Oct, 1993.

102. "Korrosion un dKorrosionsschutz von Stählen", Verlag Chemie, Weinheim

(1977), Α. Rahmel und W. Schwenk

103.

"Hochtempe raturkorrosionsbeständigkeit des wann festen Stahles

XIO CrM oV Nb 91 ", K.Haarmann, W. Schwenk, J. Venkateswariu und

M. Zscha u, VGB Kraftwerkstechnik 73 (1993) 9, p. 837

104.

"Werkstoff P 91 ", E. Ambs, E. Tolksdorf, K.E: Leich, R. Schlieben und

H. Schwarzwalder, VGB Kraftwerkstechnik 73 (1993) 7, p. 634

105.

"Improved Coal-Fired Power Plants", R.B. Scarlin, P. Sche pp, State of European

COST-Activities. Second EPRI International Conf., Nov. 2 - 4, 1988, Palo Alto,

USA

106. "Improved Coa l-Fired Power Plants", R.B. Scarlin, C. Franklin, Third EPRI Inter-

national Conf., April 2 - 4

1991,

San Francisco, USA

107.

"High Performance Steel by Hot Isostatic Pressing. Processing, Applications and

Perspectives", K. Torsell, Proc. of Conf. High Temperature Materials for Power

Engineering, Liège, Belgium, 1990

108.

"Pulvermetallurgische Herstellung von Dampfleitungs-Kompo nenten mit kom-

plizierter Geometrie", U. Heisel, T. Hollstein, P. Schepp, K. Torsell,

G. Schuhmacher, VDI Berichte Nr. 7978, 1990

- 5 6 -



109. "The effect of man ufacturing route on the creep streng th and microstructure of

12%

Cr

steel",

R. Wu, P.J. Henderson, R. Sandström, B.G. Invarsson, Sheffield,

June 1991

110.

"Auslegung dickwandiger Bauteile auf der Basis neuer Stähle", F. Pietzonka,

C. Henry and C. Torsell, VGB Conf. on Fossil-Fired power P lants with Advan ced

Design Paramters, June 1993, Kolding (Dk)

111.

"Newly deve loped High Tem perature ferritic-martensitic Stee ls", R. Blum,

J.

Hald, W. Bendick, A. Rosselet and J.C. Vaillant, VGB Conf., 1 6 - 1 8 June

1993, in Kolding (Dk)

112. "Eigenschaften und B ewährung des Stahles X 20 CrMoV 12 1 im Kraftwerk",

H. Jesper, H.R. Kautz, VGB-Konferenz "Werkstoffe und Schw eisstechnik im

Kraftwerk 1985", Essen, Feb. 1985, S. 274 - 316

113.

"Das Rohr im modem en Kraftwerksbau", G. Kalwa, K. Haarmann, Stahl und

Eisen 102 (1982) Hf. 17, S. 829 - 832

114. "Planung und Betrieb überkritischer Dampferzeuger mit 311 bar im Kraftwerk

Kawagoe", T. Kawamura, T. Toyoda, I. Kurihara, H. Haneda, VGB-Kraftwerk-

technik71 (1991), H. 7, S. 6 37 -6 43

115.

"Kohlenstaubbefeuerte Kraftwerksblöcke mit fortgeschrittenem W as se r/ Dampf-

prozess", S. Kjaer, VG B Kraftwerkstechnik 70 (1990) H. 3, S. 201 - 208

116. "Advanced High-Chromium Ferritic Steels for High Tem peratures", T. Fujita,

Metal Progress, August 1986

117. "Development of a 9% CrMoV Steel", H. Matsumoto, M. Sakakibara,

T. Takahashi, H. Sakurai, T. Fujita, First International Conference on: Improved

Coal-Fired Power Plants, November 19 -

21,

1986, Palo Alto, California

118. "Current Progress in Advanc ed High Cr Ferritic Steels for High-temperature

Applications", T. Fujita, ISIJ International, vol. 32 (1992), no. 2, p. 1 7 5 -1 8 1

119. "Developm ent of Tubes and Pipes for Ultra-Supercritical Power Plant Boilers",

H. Naoi, H. Mimura, M. Ohgami, M. Sakakibari, S. Araki, Y. Sogoh, T. Ogawa,

H. Sakura i, T. Fujita, Shinnittestsu G iho, No. 347, November 1992, P. 27 - 31

120.

"The super 12%Cr Boiler Tubing", F. Masuyama, T. Daikoku, H. Haneda,

K. Yoshikawa, A. Iseda, H. Yuzawa, COST-EPRI W orkshop, Creep-Resistant

9 - 12%Cr Steels, Schaffhausen, Switzerland, October 13 -14, 1986

121.

"Development of New 12%Cr Steel Tubing (HCM12) for Boiler Application",

A. Iseda, Y. S awarag i, H. Teranish, M. Kubota, Y. H ayase, The Sumitomo

Search, No. 40. November 1989, p. 41 - 56

122. "Advanced 9Cr/12Cr Steels for Thick Section Compo nents of Fossil-Fired Power

Stations", E. Metcalfe, 123rd ISIJ Annual Meeting, Chiba April 1, 1992, Vol. 5,

p. 513

- 5 7 -



123.

"Creep Life Prediction of a 9Cr-0,5Mo -1 ,8 W Steel by Modified Theta Projection

Method", H. Mimura, M. Ohgami, H. Naoi, K. Maruyama, VGB-Konferenz

"Restlebensdauer 1992" 6 - 7. Juli 1992, Mannheim

124.

"Untersuchungen von Grundstoff und Schweissverbindung der neuentwickelten

Stähle NF616 und P91 (9%Cr- Stähle), R. Blum, J. Hald, E. Lund, VGB-Kon

ferenz "Werkstoffe und Schwelsstechnik im Kraftwerk 1991" 9 - 10 Jan. 1991

Essen

125.

"Properties of a Large thick wall 9Cr 0,5Mo 1,8W Steel Pipe", H. Mimura,

M. Ohgam i, H. Naoi, Y. Sogo, T. Ogawa, T. Fujita, 3rd International Conf. on

LO O P . San Francisco (C.A.) April 1991

126. "Development of a 9Cr - 0.5Mo - 1.8W-V-Nb for Boiler Tube and Pipe",

H. Nimura, M. Ohgami, H. Naoi and T. Fujita, COST 501 and 505 Conference:

High Temperature Materials for Power Engineering 1990, Liège, Belgium,

2 4 - 1 7 Sept. 1990, p. 485 - 494.

127. "Application and Properties of Modified 9Cr-1Mo Steel Tubes and Pipe for

Fossil-Fired Power Plants", A. Iseda, M. Kubito, Y. Hayase, S. Yamamoto,

K. Yoshikawa, The Sumitomo Search No. 36, May 1988. Sumotomo Metal In

dustries, Ltd. Osaka & Tokyo, Japan.

128.

"Evaluation of High Temperature Strength Properties of New Tubes for Ultra

Super Critical Boilers", T. Sato, Y. Fukuda, K. Mitsuhata and H. Yamanouchi,

JSME-ASME Conf., Tokyo, Oct. 1993, p. 243.

129.

"Development of a High Strength 11Cr-2W-0.4Mo-1Cu Steel for Large Diameter

and Thick Wall Pipe", A Iseda et al., ibid, p. 59.

130. "Physical and Mechanical Properties of newly deve loped 9Cr 1.8W ferritic steel",

H. Nimura, M. Ohgami, H. Naoi and T. Fujita, TMS Fall Meeting 17-21 Oct.

1993. Pittsburgh

58



WNet

M

c

D

E

ω

>

o

Ι -

Ο .

E

Φ

to

ra

Q)

Χ

cni·

*

M

■ : · : ■ : ■

·.·.·.·

n

CD

co

■ : ■ : ■ : ■ :

w i s

. · . · . · . ■

ι

ijSf^

A-h

:·:·:· i o

f í íN.

í * i

i "j

CM

(1)

0 0

T f

CM

ai

*

)

•ν

ö

'-y-

z_

535 °C 560 °C 60 0 °C 650 °C

LS/RH Temperature (°C)

EÜH Single reheat

Double reheat

Fig. 1 Process Improvements with Single and Double Reheat Cycles

59



(Coal Fired Units except Kawagoe # 1,2)

(Base) 0 -

o

ra

1—

c

Φ

E

Φ

>

o

i—

O .

E

Φ

m

ra

Φ

2 -

Matsuura # 1 1.000 MW (1990)

Hekinan # 1 700 MW (1991)

Hekinan # 2 700 MW (1992)

Noshiro # 1 600 MW (1993)

Sohma # 1 1.000 MW (1994)

Sohma »2 1.000 MW (1995)

Hekinan # 3 700 MW (1993)

Tsunjga # 1 500 MW (1991)

*1

Helhoku # 1 700 MW (1995) J

Noshiro # 2 600 MW (19 94 )"

Nanao-ohta #

1

500 MW (1994)

Haranomachi #

1

1.000 MW

(1997)

Matsuura # 2 1.000 MW (1997)

(LNG) Kawagoe # 1 700 MW

Kawagoe # 2 700 MW

246 316

Main steam pressure (Kg / cm *)

Fig.

2 USO Steam Units

60



Components

Superheater tubes

Live steam pipes,

header, valve casing

High tempera ture rotors

High temperature casing

and

valve casing

Conventional

plants

CrMo steels

9% Cr steel

2 1/4 CrMo steel

1% CrMoV steel

CrMo steels

max. 593 °C

Stage I

Austenitic steels

or mod ified

9%

CrMo steel

9% Cr steels

Improved

12 %

Cr steels

Improved

9-12% Cr steels

max. 649 °C

Stage II

High strength

austenitic steels

Ausi steels

Aust. steels

(A286)

Aust steels

Fig.

3 Materials for Conventional and Improved Japanese Coal-Fired Power Plants

AIMS OF THE RESEARCH PROGRAM

Evaluat ion of s team turb ine design and mater ia ls

for coa l - f i red power p lant w i th doub le reheat

Phase 0: 3 1 0 bar

Phase 1: 3 1 0 bar

56 5 °C - 56 5 °C - 56 5 °C

59 3 °C - 59 3 °C - 59 3 °C

Ou tpu t r ange: 400 M W to 900 M W

Research pe r iod: 198 6 - 19 90

Fig.

4 Development of More Economic Coal-Pired Power Plant in the USA

(EPRI Project 1403-15)

- 6 1 -



ü

o

s

ö

I I

n

C Φ

g

Ä Q.

S

ra

S

S

o (β w

ra

co

ra

E E E

700 η

600

500 Η

400

300

200

100

0

♦ Siemens / AEG

max. steam temperature

ABB

(BBC + Escher Wyss)

max. output

50

— ι —

52

τ » f / l å f i r [

ι—

60

Fig. 5

54 56 58 60 62

Year of commissioning or construction

Turbine Data of European Plants

64

Components

Rotors, discs

Valve casing

Turbine casing

Blades

Inlet sections

Bolts

max. s team inlet 1

600 °C

X 2 0 C r M o V N b 1 2 1

X 2 2 C r M o V 1 2 1

X 4 0 C r N i C o M o W N b 1 3 1 3 1 0

G-X22CrMoV121

X8C r N i M oN b1616

AISI 347

G-X22CrMoV121

AISI 347

X 1 2 C r N i W T i 1 8 1 0

X 4 0 C r N i M o W N b 1 3 1 3 1 0

X8C r N i N b1613

X 2 0 C r M o V N b 1 2 1

X 2 2 C r M o V 1 2 1

X8C r N i M oB N b1616

X40Cr fMiCoMoNb131310

emp eratures in °C

> 6 00 - 65 0 °C

(24CrMoV 55)

X 2 2 C r M o V 1 2 1

X 8 C r N i M o V N b 1 6 1 3

X 8 C r N i M o B N b 1 6 1 6

X 1 2 C r N i W T i 1 7 1 3

G-X22CrMoV121

G-X8CrNiNb1613

X 8 C r N i M o - N b 1 6 1 6 - T i 1 6 1 6

X 8 C r N i M o V N b 1 6 1 3

G-X22CrMoV121

G-X12CrfMiWTi1713

AISI 347

X 2 2 C r M o V 1 2 1

X 7 C r N i M o N N b 1 6 1 3


X 8 C r N i M o V N b 1 6 1 3

AIS I 316

AIS I 347

X8C r N i N b1613

X 8 C r N i M o V N b 1 6 1 3

X 2 2 C r M o V 1 2 1


X15Cr fMiWNNb1912

X 8 C r N i M o V N b 1 6 1 3

Fig.

6 Steels in European Turbine Power Plant with S team Inlet Temperatures up to

650°C

6 2 -



Net efficiency

100

%

90

8 0 -

70

60

50

40

30

2

+

+

1

3

y Vi

£ £

4

. e «

ces5

5

+

6

1

Esberg

3

2 advanced

steam process (1995)

3

advanced

steam process (2000)

4 pressurized

fluidised bed (today)

5 coal gasification

(today)

6

combined gas and

steam process

after

S,

Kjaer

400 600 80 0 10 00 °C

Process temperature

Fig. 7

Coal-Fired Power Plant with Advanced Water/ Steam Processes

(Elsam Project A/S Denmark)

Intemat onalProjectsOKÄdvanceaiEoweriPIantsÄ

i

Japan

Ψ

USA

Ψ

EUROPA

Í " R & P : E P D C j

Turbine : MHl - Toshiba - Hitachi

ι

1981 -1991

I

•

3iebarS66/56e/566*C

- · 314

t* f

583/593/593*0

- · 343 bar 649/593/593*0

Π

1989'

1991 ■

1990

1993

R & D : EPRI

GE - Westinghouse

ι

Studie 1978-1980

I

•

310 bar 566/566/566*0

- · 310bar593/S93/593 C

•

345 t w 649/649/649-C

R & D :

1986-1993

■ H

EPR I- RP 1403-15

300 -

9O0

MW

¡ Steam Power Plants ¡

-

700 MW

:

246 bar/538 C/593*C com. 1993

-

500 MW

:

248 bar/5e8*C/5S3*C com. 1SS4

- 600 MW : 248 bar/56e*C/593*C com. 1894

-1000 MW : 248 bar/593*C/593*C com. 1967

-

600 MW

:

248 bar/5ee*C/S93*C com. 1998

•

1050 MW

:

255bar/eoO*C/610'C planned

i_

ι ι

ABB

-

GECA

-

MAN

-

Siemens

-

NB

-

QF

-

VOEST

-

SV

-

FEL - Súber

-

VSQ

-

Bonier

-

ENEL

-

NP

I

1983-1995

I

•

3C0 bar 600/600/600*C

•

3C0 bar 600/620'C

•

300 bar 620/650*0

[ Steam Power Plants l

•

400 MW

:

280 bar/580*C/580*C/580*C ordered 93

•

700 MW

:

275 bar/58CTC/e00*C planned

- 440 MW : 266 bar/5B0*C/800*C planned

-

900 MW

:

259 bar/680 C/582*C planned

•

750 MW

:

250 bar/575*C/595 C planned

Fig.

8

International research projects on improved coal-fired power plant

- 6 3



US A Germ. :1

UK: 2

France : 3

Steel Development

Χ 22 CrMo(W)V 1 2 1 / rotors, casinos, bous, Wades, pipes

H 46; FV 448 / bo js, blades, oas turbine discs

56 Τ 5 ; bolts, blades.

Japan : 4 TAF / biade«, disca, aman rotors

1) —tog arm i«a - is

S mcHÉngCOSTeoi'I

Development tor

tastbreeder

US A

:

5 1 1 % CrMoVNbN / rotors (GE)

USA : 6, | X 10CrMoVNbN8 1 (P it) /pi p« «, oreas««vassets, cesino, 1)

Japan: 7 I TR 1100;TR 12 00/rotors

COST 501 : 8 | X 18 CrMoVNDB 9 1 / rotors

COST 501, EPRI: 8 i X 12 ClMoWVNbN 10 11 /rotors 2)

J « p « n : 1 0 , 11 N F - 1 6 / H C M 1 2 A / P C « «

MPa

120

too 000 h creep strength at 600 "C

\

2.3

1

5

4

> |

6

. - - '

0, 11

s

8

MPs

120

Fig. 9 Development of Heat Resistant 9 - 12%CrMo(W )V(Nb)N(B)-Steels for

Improved Power Plants

Country

USA

BRD

UK

France

USA

UK

France

USA

Japan

USA

Japan

COST 601-2

Steel

C h e m i c a l C o m p o s i t i o n In %

C | Cr | Mo | Ni

Basic Steels (service exDerience U D to

565*

CI :

T 9 (X 12 CrMo 9 1)

X 22 CrMoV 12 1

Η 46 (X 16 CrMoVNbN 11 1)

EM 12 (X 10 CrMoVNbN 9 2)

A IS I 422 (X2 2Cr Mo WV 12 1)

FV 488 (X 13 CrMoVNbN 10 1)

56 Τ 5 (X 19 CrMoVNbN 11 1)

11 % CrMoVNbN (GE)

TAF (12 % CrMoVNbB)

0.12

0.20

0.16

0.10

0.23

0.13

0.19

0.18

0.18

9.0

12.0

11.5

9.0

12.5

10.5

11.0

10.5

10.5

1.00

LOO

0.65

2.00

1.00

0.75

0.60

1.00

1.50

0.50

0 7 0

0.75

0.70

0.40

0.70

0.05

N e w l v d e v e l O D e d S t e e l s ( s e r v i c e t e r r t D e r a t u r e s U D t o

Τ 91 (X 10 CrMoVNbN 9 1)

T R 1 1 0 0

T R 1 1 5 0

N F 6 1 6

H C M 1 2 A

X 1 2 C r M o W V N b N 1 0 1 1

X 18 CrMoVNbB 9 1

0.10

0.14

0.13

0.07

0.10

0.12

0.18

9.0

10.2

10.3

9.0

11.0

10.3

9.5

1.00

1.50

0.30

0.50

0.40

1.00

1.50

< 0 . 4 0

0.60

0.50

0.06

< 0 . 4 0

1.0 Cu

0.80

0.05

W

(0.5)

1.00

600*

C

2.00

1.80

2.0

0.80

V I Nb

0.30

0.30

0.20

0.25

0.15

0.20

0.20

0.20

0.300

0.450

0.450

0.450

0.085

0.150

JU

0.22

0.17

0.17

0.20

0.22

0.18

0.25

0.080

0.055

0.050

0.050

0.060

0.050

0.050

Ν I Β

Creep

Strength

MPa at 600 ' C

10*4 h

0.050

0.050

0.050

0.060

0.010

0.040

59

103

118

120

130

139

144

165

216

0.050

0.040

0.050

0.060

0.060

0.055

0.010

0.004

0.003

0.010

124

170

185

160

156

165

170

10"5h

34

59

62

82

60

64

64

(85)

(150)

94

(100)

(120)

(132)

(127)

(107)

(122)

Fig. 10

Chemical Composition and Creep Strength of Commercial and Newly

Developed Ferritic Heat Resistant 9 - 12%CrMo-Steels

■ 6 4 -



Steel

X22CrMoVÏ21

Newly-developed steels

X10CrMoVNbN91

(P91)

X12CrMoWVNbN1011

X18CrMoVNbB91

Component

Pipes, Bolts, Rotore

Casings

Pipes,

Forgings

Casings

Rotors, Casings

Rotors, Casings

Strengthenin

Precipitation Strengthening

Chromium carbides of

M23C6 type

Many, finely-distributed

\ stable carbides:

\ · Chrom ium-carbides

) of M23C 6 type

/ · V/Nb-carbonitr ides

' of MX type

Many, finely-distributed

stable carbides:

• Chromium-carbides of

M23C6/M23(C,B)6 type

• V/Nb-carbonitr ides

of MX type

3 Mechanisms

Other Mechanisms

Solid solution strength

ening by Cr and Mo


ening by Cr and Mo


ening by Cr, Mo and W


ening by Cr and Mo

Boron within the

carbides reduces

diffusion and coarse

ning rates

Fig.

15 Microstructures of Creep-Resistant Ferritic 9 - 1 2 % C rMo(W)V(B)(Nb)N -Steels

As received Aged at 650°C for 3,000 hours

Fig. 16 Appearance of Laves phases in a W-alloyed steel after long-term ageing

- 6 7 -



Creep

strength

Creep curve

Carbide reactions

M

2 3

C

6

>-MX > M

6

C, etc.

Basis strengthening mechanisms.

Time

Fig.

17 Effect of Microstructural Changes on Creep Strength

Fig. 18 Schematic concept of new Generation of 9 -1 2 % Cr Steels

68



Corrosion

resistance

Creep

strength

VN, stable & effective

Nb (C,N), grain refinement

> W, stable

Weldability

Stable

long-term strength

1

> | T W O

phase

r->(

L

ow

' ι toughness j

Toughness

Martensite

δ -f er rit e ^ 5 %

Heat embrittlement

^K

11 mass % Cr

V, Nb and Ν

addition

high W & Mo

Β addition

ι

ι :

Low Ni

ΣΓ

Low C

Cu addition

■^ Cr eq.

¿

9 mass %

Low Si

0.1 C -1 1 Cr - 2 W - 0.4 Mo - 1 Cu - 0.2 V - 0.05 Nb - B - 0.06 Ν

Fig. 19 Development concept of the Japanese piping steel HCM 12A

X 10 CrWMoCuVNbNB 11 2)

69



type of steel C SI Μη

Ρ

S

ΑΙ Cr Mo NI V W Nb B

Ν

trial melts for forgings

A

Β

D

E

F

♦nitrogen

♦boron

♦tungsten

♦tungsten/

molybdenum

♦molybdenum

0 05-0 06

0.15-0.18

0 10-0 16

0.10-0 18

0.10-0 1B

0.1-0.3

- 0.1

- 0 . 1

- 0 1

- 0.1

0.2-0.7

- 0.1

- 0.5

- 0.5

- 0.45

<.01

< 0 1

< 0 1

< 0 1

<005

<.005

< 005

< 005

<005

< 0 2

< 0 1

< 01

<.01

9-12.5

9-10.5

10-12

10-12

9 5-12

1-2

-

1.5

< 0.5

- 1

1-2

0.5-1.0

- 0 . 1

0.5-1

0

0.5-1.0

0.5-1.0

- 0 . 2

- 0 . 3

- 0 . 2

- 0 . 2

- 0 . 2

- 2

- 1

- 0 6

- .06

- .06

- 06

- .06

005-01

0.1-0.3

<07

<07

<.07

components rotors

Β

E

F

0.17

0.12

0.11

0.07

0 10

0.03

0 0 6

0 4 5

0.52

.007

006

.010

001

.002

.005

.012

.008

.006

9.34

10.39

10.22

1

58

t 06

1 42

0 12

0.74

0 5 8

0.27

0.18

0.18

0.81

0 0 5 9

0.045

0 0 5

.0080

0002

.0012

.015

052

.056

COST

501/11

P rogram, Chem ica l C om pos i t ion

(%)

9 - 1 2 % Cr -S tee l s

Μ93906

.

Fig. 20 Chem ical Com position (%) for trial forging melts

70-



1000

600

10 IO

2

O

3

IO* IO

5

»'me s

Fig. 21 Time -Tem perature-Tra nsform ation diagrams for the new steels

Task

Heat treatment-

Austen, temp. °C 1020 10 70 (11 00 ) 1120

Temper, temp. °C ~ 70 0 - 72 0 - 70 0 - 72 0 ~ 70 0 - 7 20

1.

Microstructure

2. Strength

2 0

°C,

6 0 0

°C,

65 0 °C

3. Toughness A

v

-T, FATT

4. Creep behaviour

α -, σ

2

, σ

3

a t 600 °C , 650 °C

smooth, notched

5. ISO-Stress-behaviour

a

- 1 0 0 M P a a t 6 2 0 °C ,

6 4 0 °C , 6 6 0 °C , 680 °C,

700 °C

6. Long time toughness behav

iour at 48 0 °C. 60 0

°C ,

650 °C

for 1,3, 10 χ 1 0

3

h

then tasks 1—3

7. Overaging at

700 °C / 200 h

then tasks 1—5

Melts

D,E

D.F (B)D ,EF (B)D,E,F A.E.D A.E.D

Α-N-Stee ls

B-B

D - W

E-W/Mo

F - M o

R g .

22 Investigation Progra mm e of Trial Melts

71



ASTM-Grain S ize

5 -

-1 -

-3

<700°C >700°C

<>

T T :

< 700 "C > 700° C

T T :

< 700° C > 700° C

1020°C

1070°C(1050be i B )

Aus ten i t i s ing Tempera ture

1 1 2 0 ° C ( 1 1 5 0 b e l B )

F ig . 24 A us t en i t e g r a i n s i z e f o r t he t es t m e l t s

• 7 3 -



800

700

600-

500.

X N

H B

♦ W

OW/Mo

ώ Μο

RpO.2 (MPa) RT

Aim «no

rotor Β

+

χ

χ

5 β^Α * ο

Α

« 5

ca

· *

^ Γ - ^ Γ

ώ

Χ^Α

tor

•

χ

Ο

Α

Χ

<Α

χ

t

¥ ·

χ

Χ χ

a a *

5

*<* ^ Α

-ν.

* ·

χ ♦ χ*

5

ν

Α

χ, Ι

0

A A

¿

FATT (°C)

r-100

50

+

1 10

Aging t ime

1 3 10 30 1

-50

3 1 0 *1 0

3

h

As rece ived ' As rece ived ' 480°C

+ 7 20° C / 200h

î 600° C / 27 0 .000h

600°C

* at 530°C

650°C

Fig.

25 Yield and FATT values after long-term ageing of the forged steels

-74



100 -jFATT (°C)

80 -

aging temp.

rotor Β

'triai melt E

I trial

melt

F

480 530 600 650°C

y.

. ( . χ

Λ A *

z m ♦

conv.Rotor

2000 4000 6000 8000

aging t ime (h)

10000

Fig. 26 Effect of long-te rm ag eing on tough nes s of trail melts D, E and F

7 5 -



100 -

80 -

60 -

40 -

20 -

0 -

-20 -

-An .

FATT (°C)

A

♦

A; ;

-41 ) * t

C 2000

aging temperature : 600°C

specimen position : near center

orientation : radial

530*C

Rp0 .2(M Pa)i >600 >700

+ B rotor B

+ W +Mo rotor E

+ Mo rotor F

♦

A

conv. rotor ;

♦

A

O

Fig. 2 7

1 1 1 1 1 1 Γ

4000 6000 8000 10000

aging time (h)

Dependence o f FATT on age ing t ime and temperature

Stress (MPa)

1000 q

*

^

100 :

10

A s r e c e i v e d

■*

JC-SN&

6 0 0 ° C

1 0 0

000 h

— ι

*

A + a

Δ<0> □

-ώ-Ο-Β*

-F Β

-FW

(B)

(D)

-F W -F Mo (E)

•F Mo (F)

M e a n v a l u e s S E W 5 5 5

6 0 0 ° C

1 0 0 - 1 0 0 0 0 0 h


6 5 0

e

C 1 0 0 - 2 0 0 0 0 h

Ι ι ι ι ι j

•e-

23.00

24.00

25.00 26.00 27.00

LMP = T ( 25 + log(t) ) /1 00 0

Fig. 28 Cre ep behaviou r of steels with yield strength of 600 - 650 M Pa

7 6 -



Stress (MPa)

1000

■ As rece ived

100

10

Χ + Ν

♦ ■ -FW

O 3

-F

W

-F

MO

(A)

(D)

(E)

600 °C 100 000 h

— ι

M e a n v a lu e s S E W 5 5 5

6 0 0 ° C 1 0 0 - 1 0 0 0 0 0 h

■ ι « ■ ■ ■ ι

M e a n v a lu e s S E W 5 5 5

6 5 0 °C 1 0 0 - 2 0 0 0 0 h

ι 1

t-

23.00 24.00

25.00 26.00 27.00

LMP = T (2 5 + l og ( t ) ) /100 0

Fig.

29 Cree p beha viour of steels with yie ld strength of 700 - 75 0 MP a

Stress (MPa)

1000

100

10


6 0 0 ° C 1 0 0 - 1 0 0 0 0 0 h

Ι Ε0ΕΟΔΔ

▲ ♦ ■ + W (D)

Δ

ό

□ + W

F

MO (E)

■ -F Mo (F)

I I I I


6 5 0 ° C 1 0 0 - 2 0 0 0 0 h

23.00

24.00

25.00 26.00 27.00

LMP = T ( 2 5 + l og ( t ) ) /1000

Fig. 30 Creep behaviour o f s tee ls wi th y ie ld s t rength o f 600 - 650 MPa after

ove rage ing

■77 ·



Stress (MPa)

1000

100 :

10

♦ ■ +w

ζ > -F W -F Mo

I I I I I I I I I

Mean va lues SEW555

600°C 10 0 - 10 0 000 h

(D)

(E)

Mean values SEW555

650°C 1 0 0 - 2 0 000 h

23.00 24.00

ι ι ι ι Ι I I I I

25.00 26.00 27.00

LMP = T ( 2 5 + l og ( t ) ) / 1000

Fig.

31 Creep behaviour of steels with yield strength of 700 - 750 MPa after

overageing

78



¡¡Stet

A 1

A 2

B 2

D l

0 2

D 3

E l

E2

E3

F1

F 2

F3

H e i l

T f M t m t n t

1 1 2 0 / 6 9 0

1 1 2 0 / 7 2 0

1 1 2 0 / 6 9 0

1 1 2 0 / 7 2 0

1 1 0 0 / 7 0 0

1 0 2 0 / 6 8 5

1 0 2 0 / 7 2 0

1 0 7 0 / 7 3 0

1 1 2 0 / 7 3 0

1 0 2 0 / 6 * 5

1 0 7 0 / 6 9 5

1070 /720

1 0 2 0 / 6 9 5

1 0 7 0 / 6 9 5

1 0 7 0 / 7 3 0

1 1 2 0 / 6 9 5

1 0 2 0 / 7 2 0

1 0 7 0 / 7 2 0

1 0 2 0 / 7 1 0

1 0 7 0 / 7 0 0

1 0 7 0 / 7 2 0

1 1 2 0 / 7 1 0

1 0 2 0 / 7 1 0

1 0 7 0 / 7 0 0

1 0 7 0 / 7 2 0

1 0 7 0 / 7 2 0

1 0 7 0 / 7 2 0

1 0 7 0 / 7 2 0

Extra,

p o l a t i o r R u p t u r e O t t a 6 0 0 * C

( T M K 1 1 0 0 0 0 h -

1 6 5 M P 1 .

3 0 0 0 0 h - 1 3 3 M P a )

1969

A R O A

151

151

160

139

156

160

182

146

146

165

154

144 100

150

144

139

10000h

1990

A H

90

115

100

OA

152

130

136

14«

140

115

135

146

154

145

165

140

134

165

4 7

137

130

148

140

135

152

135

130

140

112

128

138

135

132

100

130

132

128

135

1991

AR OA

90

110

115

100 130

170 135

145 146

130 133

135 130

131 136

140

115 112

135 128

146

150

137 133

161

140 135

139 135

145

165 122

147

145

130 100

148

140 130

140 132

128

141

3 0 0 0 0 h

1991

AR OA

91

145

102

112 11»

116

109

106

113 113

94

102

111 108

108

120 120

120 122

130

133 112

125

113

103

109

112 115

122 121

E x t r a p o l a t i o n R u p t u r · D a t a 6 5 0 * 0

f T M K l 1 0 0 0 0 h - S O M P a )

1989

AR OA

65

85

SO

80 80

7 2

70

74

90

85

60

83

88

88

9 0

112

130

looooh

1990

A R O A

5 8

88

62 83

64 60

65 66

67 67

65

75 78

73 70

64

70

67 74

76

82 81

85 82

78

80

74

58 58

73

73 73

83 81

6 1

8 7

1991

A R O A

50

50

58

55 60

91 75

62 68

64 60

65 68

70 70

65

75 78

73 70

64

70

71 74

76

82 81

85 82

82

78 62

78 78

74

58 5«

73

73 73

86 82

68

82

3 0 0 0 0 h

1991

AR OA

46

64 66

67 69

6 8

70 72

57

»

E x t r a p o l a t k i n I S O - S i r · « *

l O O M P i / 6 0 0 * 0 ( x 1 0

A

5 h )

1989

| AR

O A

2.5

< 1

2

2

4

< 1

2

< 3

< 3

3

3

5

2

2.5

1.5

1.5

< 1

3

4

5

1.5

< 1

3

0 . 6 1 '

0.1

2

1.5

3

1.5

2

0.6

1

2

2

1990

AR OA

0.42

0 .01

0 .22

0 .23

7 .16 0 .36

0.57 1.55

0.70 0.82

0.76 1.01

1.50 1.38

0 .84

1.31 0.11

0 .82 0 .88

2 .52

2 .50

3 .00

3.34

1.15 0.9B

1.08 1.48

3.84

2 .34 5 .90

2 .02 2 .33

1.49

2,3 1.29

0 .61

0 .74 0 .81

1.42 1.26

2 .02

1991

AR

0.42

0 .01

0 .22

0 .23

7 .16

0 .57

0 .70

0 .76

0 .61

0.84

1.31

0 .82

2 .52

2.50

3 .00

3.34

1.22

1.14

3.64

2 .34

2.02

1.49

2,3

0.7B

0 .73

1.56

OA

0.36

1.55

0.82

1.01

1.38

0 .33

0 .88

1.03

1.61

5 .90

2 .33

1.29

1.04

1.39

8 .14

Fig.

32 Short-term creep rupture and iso-stress data for trial melts D, E and F

■

7 9 -



Task

Specimen posi t ion

Bottom and Top of Ingot

Surface Center

Middle of Ingot

Surface Half Radius Center

1. Macro- /Mlcrostructure

2.

Strength at 20-650°C

3. Toug hness Av -T, FATT, Kic

4. Creep behaviour

at 550, 600, 650°C up to 100.000h

then task 1

5. ISO-stress behaviour

0 = 1 0 0 M P a

at (700), 680, 660, 650, 640, (625)°C

6. Long t ime toughness

at 480, 600 , 650°C

then tasks 1,2, 3

7. Overaging at 700°C/200h

then task 5

8. Low cycle fatigue at 20, 550, 600°C

9. Cycl ic crack grow th, Δ κ

0

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

F ig . 35 Invest igat ion Prog ram me of Rotors B, E and F

T e st m a t e r ia l a f te r t e m p e r i n g f o r y i e ld s t r e n g t h R p 0 . 2 ^ 70 0 M P a ^ \ \ X \

Rad ia l Core

E : 0 350 / F : 0 206

Test ma te r ia l a f te r tem pe r ing fo r y ie ld s t re ng th Rp0 .2 ^ 60 0

M P a

/ / / / ,

Fig. 36 Sam ple Posi t ion P lan of Rotors E and F

8 1 -



Stress (MPa)

10

expected creep rupture curve at 600°C

stress

MPa

180

160

135

120

100

90

80

expected time

h

- 1 0 0 0

~ 3000

- 1 0 0 0 0

~ 30000

- 1 0 0 0 0 0

>100000

>100000

[ I l i l i l í

100 1000

10000 100000

Time (h)

Fig. 37 Selected Stress Levels for Long-term Creep

Tests,

Rotors B, E and F

Specimen orientation : radial

Rp0.2 MPa

Rm MPa

Av 20°C J

FATT °C

6 42

801

2 2

+ 4 5

670

813

33

+45

627

799

21

651

813

33

+60

Bottom

Disk No. 6,7 separate heat treatment similar to center of diameter 12 00 mm

QHT : 1100 °C / 2h / +59 0 °C / 8 h / +7 00 °C / 16h

Fig. 38 Mechanical Properties at different locations in Trial Rotor B

82-



Rotor

Β

E

F

Rim

(top tangential)

RpO.2 Rm

MPa

642 801

801 914

647 783

770 892

600 737

Av 20°C

J

22

27

76

55

99

FATT

°C

+45

+55

+20

+30

-2

Center

(midsection radial)

Rp0.2 Rm

MPa

651 813

744 875

631 774

729 855

609 764

Av 20°C

J

33

86

146

42

75

FATT

°C

+60

+ 5

-10

+40

+5

Hg.

39 Tensile and impact Properties of Rotors B, E and F

Specimen orientation

Near surface - tang

Center - axial

x

radiai"

4

"

Bottom

0350

749+

874+

56 +

+ 15 +

744 x

875 Χ

86 Χ

+ 5 x

RpO.2 MPa

Rm MPa

Av20°C J

FATT °C

Fig. 40 Mech anical Prope rties at different locations in Trial Rotor E

(Yield strength c a. 750 MPa)

83



Bottom

Martensite

Ferrite

~ 100%

< < 1 %

> 99%

< 1 %

> 99%

< 1 %

~ 100

<< 1

~

100%

< < 1 %

100%

< 1 %

100%

<

1 %

Austenite

grain size gene ral

parts

1 - 3

3 . 5 - 5

1 - 3

3 . 5 - 5

5 - 7

3.5 - 4.5

5 - 7

3.5 - 4.5

6 . 5 - 8

4 - 5

3 - 5

1.5 - 2.5

3 - 5

2 - 2 . 5

QH T : 1070° C / 1 7 h / Oil + 570° C / 25 h + 690° C / 24 h + 715° C / 22 h

Fig.

41

MicrostructuraJ investigation of Trial Rotor E

Specimen or ientat ion

Near surface - tang

Cen ter - axial

x

B o t t o m

618X

748 x

143X

5 x

630X

770X

157X

-5 x

6 3 6 X

7 6 4 *

1 4 6 X

- 1 5 *

Rp0.2 MPa

Rm MPa

Av20°C J

FATT °C

\

627

764

130

10

\

655

786

105

15

647

783

76

20

639

773

135

20

Fig.

42 Me cha nical Prop erties at different locations in Trial Rotor E

(Yield strength ca. 630 MPa)

- 8 4 -



Specimen orientation

Near surface - tang

Center - axial *

radial

+

607X

746*

115*

+ 0 *

80 +

+ 4 +

609 x

7 6 4 *

87 x

- 1 2 *

1

75 +

+ 4 +

621 x

765*

108*

_ 4 x

67 +

±0 +

Rp0.2 MPa

Rm MPa

Av20°C J

FATT °C

\

626

772

81

+ 1

\

642

788

92

- 5

' /

600

737

99

- 2

/

610

751

65

+ 5

Fig.

45 Me chan ical Prope rties at different locations in Trial Rotor F

(Yield strength ca . 610 MPa)

700 -

600 -

500

400

300

200 -

100 -

0

p0.2

(MPa)

specime n position : near center

orientation : radial

Rp0.2 (MPa)

Rotor B

Rotor E

Rotor F

conv. rotor

>600

X

Δ

0

>700

♦

A

0 100 200 300 400 500 600

Temperature (°C)

Fig.

46 Yield strength at elevated tempe rature for specimens from the trial rotors

compa red wi th steel X

21

CrMoV 12 1

86



Δ

Stress (MPa)

1000

RA

1

1

RpO.2 -670 MPa

Trial melt BO O

Rotor B2 half radius long. D tang. O radial Δ

Rotor B2 center

long.

■ tan g. ·

=F

23.00 24.00

25.00 26.00 27.00

LMP = T ( c + log ( t ) ) / 1 000

Fig.

47 Creep Rupture Strength for

Trial

Melts

B0

and Rotor B2

Stress (MPa)

1000

1

R p 0 . 2 ~ 6 3 0 MPa

Trial melt

E1

O

Trial melt E2 O

Rotor E surfac e D center

600°C 100 000

h

23 00 24 00 25 00 26 00 27 00

LMP = T ( c + l o g ( t ) ) / 1 0 0 0

Fig. 48 Creep Rupture Strength for Trial Melts E1 and E2 and Rotor E

8



Stress (MPa)

1000

1

1

23.00 24.00 25.00 26.00 27.00

LM P = T ( c + l o g ( t ) ) / 1 0 0 0

Fig.

49 Creep Strength for Trial Melt E2 and Rotor E

Stress (MPa)

1000

100

10

Π

,□

* L - *

Mean values X21CrMoV121

■* %£̂ j

R p 0 . 2 ~ 6 1 0 M P a

Trial m elt F1 o

Rotor F surface D center

23.00 24.00

25.00 26.00 27.00

LMP = T ( c + l o g ( t ) ) / 1 0 0 0

Fig.

50 Creep Rupture Strength for Trial Melt F1 and Rotor F



3 0 0

1 2 %

CrMoV Cast Steal

'(G-X 22 CrMoV 121) "

o.

E

2 0 0

100

mod.

9 Cr 1 Mo Piping Steel / ORNL

(X 10 CrMoVNbN 9 1)

Gain of Temperature by

ualng mod. 9 Cr 1 M o

1)

O

17 245

4 5 0

5 0 0 5 5 0

T e s t T e m p e r a t u r e l n " C

6 0 0

Fig.

51 100,000 h Creep Rupture Strength of ferritic Cast and Wrought Steels

No.

1

2

3

4

5

6

Cast S teel

m o d . 9 C r

1

M o

1 0 % C r M o V N b N

M J C

1 0 % C r M o V N b N

(trial mert)

1 0 % C r M o W V N b N

(trial melt)

1 0 % C r M o W V N b N

(valve body)

R e s e a r c h

P r o g r a m

E P R I 1 403 - 1 5

( M A N / G F )

J S W / T O S

M H I

C O S T 5 0 1 - 2

C O S T 5 0 1 - 2

C O S T 5 0 1 - 2

C

. 0 9 -

.12

. 1 3 -

.15

.12

.13

.13

.12

M n

.41 -

.51

. 5 4 -

.60

.68

.57

.54

.62

C h e m . C o m p o s i t i o n ( W e i g h t

S I

. 2 9 -

. 49

. 2 4 -

.31

.49

.39

.33

.29

C r

8 . 9 3 -

9 .36

9 . 6 7 -

1 0 . 6 1

9.7

1 0 . 5

1 0 . 0

1 0 . 5

M o

. 9 2 -

.99

.81 -

.93

.78

1.01

1.02

.99

W

.

-

-

-

1.01

.99

N I

. 1 2 -

. 27

. 4 9 -

.55

.51

.86

.85

.93

):

V

. 1 9 -

. 22

.21 -

.23

.15

.21

. 2 2

. 2 2

N b

. 0 6 8 -

. 07 8

. 0 9 -

.11

.05

.07

.07

.07

N

. 0 4 1 -

. 05 3

. 0 3 -

.05

.04

.05

.05

.05

Fig.

52 Chemical compositions of newly developed 9 - 1 0 % CrMoVNbN cast steels

- 8 9 -



c

m

σι

o

.■S

ζ

15

10

.05

Region Β: Acceptable

^ Δ 4 Δ 5

Reflion A;

- High Creep Strength

- High Ductility

• High FATT (50)

1

Creep Strength, High Ductility, Low FATT (SO)

™-l£'l/Ñ{ //lii

s.

Region B:

y '

——ι 1 1

/ . Q ^ - ' ' R e g i o n C :

"**^ Acceptable Creep Strength

- High Ductility

- High FATTJSfl)

R e g i o n D :

- Low Creep S trength

- High Ductility

- High FATT (50)

—ι ι

8

Cr

10

11

1 2

13

Equivalent = Cr + 6 SI + 4 Mo + 1.5 W + 11 V + 5 Cb + 8 "Π + 12 Al

- 40 C - 30 Ν - 4 Ni - 2 Mn - Cu - 2 Co In wt - %

14

Fig.

53 Influence of the chemical composition on properties of 9 -1 0 % CrMoVNbN

steels

V a r i a t i o n A u s t e n i t iz i n g P r e - T e m p e r i n g T e m p e r in g P o s t w e l d O v e r - a g e i n g

H e a t T r e a t m e n t

mod. 9Cr1Mo 1040°C 1)

A 1 1 0 0 ° C 1 2 h

Β 1 1 0 0 ° C 1 2 h

B O 1 1 0 0 ° C 1 2 h

C 1 1 0 0 ° C 1 2 h

C O 1 1 0 0 ° C 1 2 h

5 5 0 ° C 2 4 h

5 5 0 ° C 2 4 h

min. 732 C 1)

7 6 0 ° C 1 2 h 7 5 0 ° C 9 h

7 3 0 ° C 1 2 h 7 2 0 ° C 9 h

7 3 0 ° C 1 2 h 7 2 0 ° C 9 h

7 3 0 ° C 1 2 h 7 2 0 ° C 9 h

730

< >

C 1 2 h 7 2 0 ° C 9 h

700

C

200 h

700 C 200 h

i) 1 h hold lime per 25 mm Ih lckntis

Fig.

54 Investigated heat treatment variations of trail melts

9 0 -



TASK

Material Pre-Evaluation:

- Manufactur ing of

Test Pieces and Welds

- Mechanical Tests

- Creep Tests

- ISO - Stress T ests

- Long Term Embri tt lement Tests

Component Program:

- Manufactur ing of

Tr ial Casting

- Creep Tests

- Low Cycle Fatigue Tests

1988

-

1989 1990

YEAR

1991 1992 1993

>

Fig. 55 Test schedu le o f cast com ponent programm e

0,2 - Limit at RT

600 -

MPa

400 -

200

Target for Heat Treatment Β and C

£ 550

mum,

S 415

tiiliin

Gr. 91

Heat Treatment Β Bo C Co

without Tungsten

B Bo C Co

with Tungsten

Fig. 56 0,2% Proof Streng th Valu es for test plates as a func tion of ana lysis and heat

t reatment

- 9 1



Weld metaMO M

AS TM A 213 Gr. 91:

ζ

1 0

1 4

0 8 -

1 2

Μ η

.80

1 .20

.30 -

.60

S I

.20

.50

. 2 0 -

.50

C r

1 0 . 0

1 1 . 0

8 . 0 -

9.5

M o

.95

1.05

. 8 5 -

1 .05

w

.95

1 .05

-

.

NI

.70

1 .00

m a x .

.40

V

.20

.25

. 1 8 -

.25

N b

.06

.10

.06 -

.10

Ν

. 0 3 0

. 0 5 5

. 0 3 0 -

. 0 7 0

Heat Inpu t

kJ/cm

15 1)

77 2 )

17 1)

WELDING DATA

I n t e r p a s s

T e m p .

c

1 5 0 - 2 0 0

1 6 0 - 1 9 0

1 0 0 - 1 3 0

P W H T

C/h

7 2 0 / 8

7 3 0 / 1 2

7 3 0 / 1 2

0 2-Umlt

MPa

6 3 7

6 6 2

6 3 7

MECHANICAL PROPERTIES

T e n s i l e

S t r e n g t h

MPa

8 0 7

8 0 0

8 0 2

E l o n g a t .

A 5

%

1 8

1 7

1 7

R e d u c i ,

o f A r e a

%

5 6

5 4

5 3

I m p a c t

E n e r g y

Joule

3 3

5 3

5 0

1) string bead 2) Weav e bead

Fig. 59 Chem ica l com posi t ion, weld ing data and proper t ies o f we ld meta l 10 M

6 0 0

4 0 0

a 3 0 0

o.

ε

c 2 0 0

σι

Ρ 100

6 0

4 0

3 0

2 0

1 0

2 3

mod .

9 Cr 1 Mo / ORNL

5 5 0 ' C / 1 0 ' h

Heat

Tratmerrt

Β

BO

C

CO

Casi No

4

O

θ

φ

θ

5

•

β

©

ο

<

$1tr-4

24 25

LMP = T( C + log t ) /1 00 0

Q

¥^j

-xLs

eoo*c/io'h

26

Fig. 60 Creep Rup tu re S t rength o f 10% CrMoV NbN and 10% CrM oW VN bN - cas t s tee l

(Cast N o 4 and 5)

-93-



\ X

^ 5 4 ^

100x300x800 mm

S t r e s s

Relief ed

R ehea t

T r e a t e d

Va lve Ches t

(about 61)

W e ld P r o c e d u r e Q u a l i f i c a t i o n

T es t P l a t es

Fig.

61 Trial castings 10% CrMoWVN bN steel

Task

C h e m i c a l A n a l y s i s

M a c r o S t r u c t u r e

Micro St ructu re

T i m e T e m p . T r a n s i t i o n - D i a g r a m

T e n s i l e P r o p e r t i e s

H a r d n e s s

(cross top and t h i c k n e s s )

T o u g h n e s s P r o p e r t i e s

C r e e p ( s m o o t h , n o t c h e s )

Low - Cyc le F at igue

L o n g t i m e e x p o s u r e

A/rlm

X

X

X

-

X

-

X

X

X

-

S p

B/core

X

X

X

-

X

-

X

X

X

X

e c l m e n

P o s i

C/core

χ

χ

χ

χ

χ

-

χ

χ

-

χ

Ion

Repair

Shallow D

Χ

Χ

χ

-

χ

χ

χ

χ

χ

•

welds

Through wall E

Χ

χ

χ

-

χ

χ

χ

χ

χ

χ

Fig. 62 Test programme of valve body

- 9 4



Heat Treatment

Condit ion

Quenched and tempered:

1100°C 12h/ forced air.

+ 550 °C 24 h/s t i l l air

+ 730°C 12h/s t i l l air

Plus st ress rel ieving:

730°C 12h/ fu rnace

0,2-Llmit

MPa

637

641

π

y

573

573

Tensile

Strength

MPa

C a s t e d -

772

771

741

740

Elong.

Reduction

A 5 of Area

% %

on Test Bars :

18 50

19 46

22 58

20 48

Impact

Energy 1)

Joule

37-43

34

π

y

64-50

44

1) Charpy-V-notch specimen

Fig.

63 Mechanical properties of pilot valve chest at room temperature

Repair welds

hea t t r ea tment :

1100*C /12h / fo r ceda i r

+ 550*C/24h/st l l l

air

♦ 730*C /12Wst i l l air

+ 730*C /12h f fumace

specimen pos

A

Β

C

D transverse

weld

E transverse

weld

ition Rp0.2

MPa

571

567

586

-

-

-

744

Rm

MPa

734

719

742

758

-

616 2)

854

A

%

18.7

20.2

18.2

-

-

-

15.4

Ζ

%

47

43

45

51 D

-

8 1)2)

51

Av

Joule

31

30

30

-

27

-

31

FATT

•c

+63

+60

+45

-

-

-

+63

1) f racture base metal

2) micro shr inkage

Fig.

64 Mechanical properties of pilot valve G-X 12 CrMoW VNbN 10 11

95



Test temperature 600°C

1 000 1000 0

Time to fracture in h

100 000

Fig.

65 Creep rupture strength of pilot valve G-X 12 CrMoWVNbN 10 11

2.00

1.00

0.80

0.60

0.20

L

100

/ 530°C (MPA Stuttgart)

1000

1

000

Cycles to crack

100 000

Fig. 66 Low-Cycle fatigue strength of casting steel G-X 12 CrMoW VNbN 10 11 - Pilot

Valve B ody / Position "A"

-96



I

O

c

o

'S

o

ü

c

o

w

c

co

Q .

Χ

LU

"cã

E

k .

α>

-C

r -

10

ι—

O)

a »

200 300 400 500

Temperature in °C

Fig 67 Therm al expansion coefficients

- 9 7 -



HATEItlAl

I 1 9 C m o V N b N l l 1

(AUBERI t D U V A L )

M O D I F I E D

9

C r M o

( B S C )

S T E E L Ol

S T E E L El

S T E E L El

S T E E L

Fl

S T E E L Fl

T A F l

H E A T T R E A T M E N T

1 1 5 0 / 6 9 0 8 h r s .

1 0 5 0 / 7 1 0

t

h r · .

1 0 5 0 / 5 7 0 6 h r · .

7 5 0 i hrs.

1 0 5 0 / 5 7 0 6 hrs.

7 5 0 6 hrs.

1 0 5 0 / 5 7 0

6 hrs.

7 00

6

hrs.

1 0 5 0 / 5 7 0 6

hrs.

7 4 0

6

hrs.

1 0 5 0 / 5 7 0 6

hrs.

7 0 0

6

hrs.

1 1 3 0 / 7 2 0 4

h r s .

0, 2 \ P S

HP.

797

593

590

577

773

551

748

794

U T S

NPt

967

735

755

743

90S

724

889

980

2 0

*C

EL

16

21

23

21

19

20

16

17

RinA

56

71

M

68

63

69

64

60

6 00

*C

0, 2 s PS

NP.

383

2 í

273

309

395

284

393

416

UTS

Κ Ρ ι

503

386

352

405

475

439

478

529

EL

21

25

2a

24

22

25

22

23

»i HA

75

90

83

86

84

87

83

SO

Heat Treatment Conditions and Medianica. Propertjes of

the Ferritic Steels (plain bar screening relaxation test)

Fig 70 Heat treatment conditions and mechanical properties of the ferritic steels for

plain bar relaxation tests

Steel

U 9 C r M o V H b N U l

T91

TAF

Bl

1)

2)

D 3 / U 4

E2/A2

D135

D191

D193

093

D

f.259

Heit treatment

1 1 3 0 - 1 1 5 0 C / 0 1 1 . 680-70 0 *C/Air

lh 1 0 5 0 * C / A l r . 4ti 70 0 C/Air

2fc 1 1 5 0 *C/0(1 » 2h 7 0 0 * C / A l r

2h 1 1 3 0 * C / A l r . 4h 7 0 0 *C

2h 1 1 3 0 ' C / Ai r . 4h 7 2 0 *C

4h 1 0 7 0 * C / A l r . 8h 5 7 0 *C/Alr . 2 4 h 7 3 0 *C

4h 1 0 7 0 ' C / A I r . 8h 5 7 0 ' C / A I r . 1 6 h 7 0 0 *C

lh

1 1 5 0

* C / A l r

. 2h 7 1 0 *C

lh

1 1 5 0

C / A I r

. 2h 7 1 0 *C

lh

1 1 5 0

* C / A l r . 2h

7 1 0

*C

lh

1 1 5 0

* C / A l r

. 2h

7 1 0

*C

lh 1 1 5 0 * C / A l r . 2h 7 1 0 * C / A ( r ♦ 2h 6 8 0 ' C / A l r

R

p 0 , 2

N/s-.'

809

790

S7B

726

666

622

750

798

D

D

548

735

R

.

Ν / ne '

960

897

1 0 3 , 9

900

840

790

874

914

D

D

767

870

A5

t

14,7

17.0

14,6

18

IS

19

17

17

D

D

1 8 , 0

1 8 , 6

Ζ

t

56

70

56

62

62

65

52

58

D

D

56

64

»

Joule

> 16

4S

36

49

96

84 - 80

48

D

D

D

D

10 0

( 6 0 «

cristallin

1) no resulti a vtlUble

Heat Treatment Conditions and Mechanical Properties

ot the Ferritic Steele (Model sroenlng relaxation tests)

Fig 71 Heat treatment conditions and mechan ical properties of the ferritic steels for

model relaxation tests

- 9 9 -



M e l t

N o .

E 5662

C Cr Al T i Mn Ir S i Pe Co Mo Hb V H

%

. 0 7

1 9 . 6 1 . 3 2 2 . 52 O . O E . 0 4

. 28

. 0 1 . 0 7

. 0 2 . 0 1 . 0 1 . 0 1 . 0 3

B Mg P S Cu Pb Su Sn S b T i

PP

2 0 -

30

1 6 -

12

20

5 200

0 . 3

0 . 1

5

Ag < 0 . 1 ppet ; Te < 0 . 5 p p a; Zn < 2 ppei

Fig 72 Chemical compositions of super clean Nim 80A

H e a t T r e a t m e n t

St an da rd i 1080 *C Bh/AC + B50 C 24/AC

+ 700 *C 16 h/AC

H o d i f i o d i 1 0 80 C B h / f u r n a c a c o o l e d a t

2 C /a in u t e t o 8 50 c , ho ld

a t 8 50 *C 24 h/AC + 700 'C

16 h/AC

0 , 2 - L i a i t

MPa

707

637

T e n s i l e S t r e n g t h

MPa

1198

1175

E l o n g a t i o n

%

2 2 . 9

2 5 . 2

R e d u c t i o n

of Area

%

35

37

l a p a c t

Energy

J o u l e

44

46

Fig 73 Heat treatment and mechanical properties of Nim 80A

a

α

S

Ol

s

a

m

s

a

a

rr

120

100

80

60

40

20 ■

Alm =

T

m

ï?

¡ι

o ♦

t - α

υ o

ι— *—*

χ

Ο G

ο s

Ol "

il

O

O

u

o

IO

o

Q

120 MPa of screening test

u

o

IN

l_l

I *.

IX.

<

1-

L i

S

0

O

u t

o

<**■

o

o

o

o

m

o

*t—

LU

8

O

O

U_

O

o

~n

o

i—

LU

Fig 74 Results of 1000 h uniaxial screen ing relaxation tests of ferritic steels

1 0 0 -



140

120

1

S'

n

20 /

♦ ♦ u

¡MI

015

i^i

20

JQ -

35

« 3 6 -

-M2D-

T

23

1U

Fig 75 Bolt mo del

PA

200

100

D Initial strain 0,20 %

I

600

e

C

^^ . Aim = 120 MPa of screen ing test

o

I N

r»

Z

Q

Ci

Χ

σ*.

r-

>

IN

<

ν

αι

I N

m

m

Ο

ÍN

<

(Ν

LU

■ <

α_

Μ

«Ι

tr-

en

Ο

σ-

Ο

«Ι F lange mater ia l : TAF

Nimonic 80A / 9 % CrMo

yj

Ο

σ-

Ο

C r l í l V

/

570

e

C

o

•w

r·»

Z

q

eri

χ

Τ

>

tv»

<

Q.

M

ru

t-

m

m

«—i

D

<■

ort

i n

CJ

UJ

O

o

■β-

IN

Γ-

Ζ

O

οι

Χ

>

U t

Γ*

U-

<

p-

( Γ

ρ-

<

G .

Μ

•LI

U

i n

ΡΠ

•Ή

Ο

■

ι η

LU

α

Rotor steels

of COST 501-2

N-al loyed

9-12 ·/. Cr-steels

R g 76 Screen ing re laxat ion tests on bo l t mod el (F lange mater ia l T91 mod.)

1 1



Standard Heat Treatment

Temperature: 550*C

1000

ε

E

100

a

ι

ο ί

1 0

O smooth specimen

A notened specimen

(DIN50 118/tumed)

Ο Δ «m

• À Core

aoc DIN 17 240

0.1

10 100 1000

Time to Rupture in Hours

10 000 100 000

Fig. 82 Creep rupture tests of high purity Nim 80A at 550°C (standard heat treatment)

- 105



Mater ia l :

N imon ic 80A

2.

Modified Heat Treatment Tem perat ure:

550*C

1000

E

ε

ζ

100

10

: - ; ■ · ; ; ■ ; ■ ' :

:

: ■ . ■ ■ : ■

O smo oth spe<

ι ¿\ notched spe

(DIN 50 118/

j Ο Δ «m

• À Core

ace.

DI

amen

ci men

(umed)

^ 17 240

1

:

:..-'.

. . . . . · . .

. .:.;.:. :.:

: ; : i ' i i i

y

:

. :

;

' ' '.

0.1

10 100 1000


10 000 100 000

à°

100

80

~ S«

S ç

< c 60

S c

4 0

o °

3 LU

"g 20

CC

· ■

:

; - ■ - ;

. . -I

I I Red uction of Area

O Elongation

O D «m

• ■ Core

Fig. 84 Creep rupture tests of high purity Nim 80A at 550°C (mo dified heat treatment)

-

107



Mater ia l :

N i m on i c 80A

2. Modified Heat T rea tment

Tem perature: 600°C

1000

E

E

100

10

>■

;4rH·:;

: . :

·.-....'.. . .

>-rl-fH-H

O smooth specimen

A notched specimen

(DIN 50 118/

Ο Δ «m

• A Core

ace or

turned)

Ν 17 240

.1

. - . · ■

· · '

; ί · . ; . . : . . ; . ;

:

Δ

:

■

■■■:■-■:

yiìm

.. i..;.·..;.;

; - : · : ■ : ■ ; ■ : ; ■ :

iv ■

_ * : :

T T S S

0.1

10 100 100 0


10 000 100 000

100

80

- o

cu

rr

6 0

20

Γ

J _ — . , -

.. .....; ..;.;.:.;

I..L.U1B

.1

I I Reduction of Area

O Elongation

O D «m

• ■ Core

Fig. 85 Creep rupture tests of high purity Nim 80A at 600°C (modified heat treatment)

108



100 300 1.000 3,000

A g e i n g t i m e , h o u rs

V 1 2 :

Standard

Heal Treatm.

V 12 N: Modified

Heat Treatm.

V 12 MMncorrect ·

modllled

Heat Treatm.

with very

coarse grain

size

10 000 30 000

Fig.

86 Impact energy of Nim 80A after exposure at 600°C (high purity and

conventional)

15

20 25 30 35 40

C h a rp y i m p a c t e n e rg y , J o u l e s

45

V 1 2 : Standard

Heat Treatm.

V 12 M:lncorroct

modllled

Heat Treatm.

with very

coarse grain

size

50

Fig.

87 Stress corrosion cracking resistance of Nim 80A me asured in cons tant

extension rate tests (high purity and con ventional)

109-



Rp 0.2 (MPa)

Rm (MPa)

A(%)

Ζ (%)

Stress

(% Rp 0.2)

729

1200

28.5

38

140

120

100

80

60

40

Medium : sa turated Na 2S 03 So ln . , 90

C

C, stagn.

Heat treatm ent : 1080°C, 8h air, +850 °C, 24h air,

+700°C, 16hair

fflw

10 100 1000 10000 100000

Time (h)

□ Standard heat treatment ■ + 600°C, 1000 h, Air

Fig. 88 Stress corrosion cracking resistance of high purity and conve ntional Nim 80A

measured in constant load tests

110



a) X20 C rM oV 12 1

Chemical

composition

i n%

C

0.2

Si

0.24

Mn

0.47

Ρ

0 026

S

0 009

Ν

0 0323

Al

0 002

Cr

11.59

Ni

0.39

Mo

0.98

V

0.28

Nb

-

1000

Tempe-

9 0

°

rature 800 -

(°C)

7 0 0 -

6 0 0 -

5 0 0 -

4 0 0 -

300

200 H

100

A

c 1 b

= 820°C

Austenitising temperature 1050°C

Holding time 8 minutes

Grain size ASTM 4 - 5

I Ι ΤΤΤΠ

- ι 1—I

I T IT

M

■ T 1—Τ 1 T IF I 1 T i l l

ΪΤΤ

1 10 100 1000

Cooling time between 800 and 500°C

10000 100000

Seconds

b ) X 1 0 C r M o V N b 9 1

Chemical

composition

i n%

C

0.10

Si

0.36

Mn

0.42

Ρ

0.017

S

0.004

Ν

0.058

Al

0.024

Cr

8.75

Ni

0.13

Mo

0.96

V

0.2

Nb

0.07

1000

900

Tempe

rature 800 -

(

°

c )

700

6 0 0 -

5 0 0 -

400

300

200

100

A

c1 b

=810°C

A + K

Austenitising temperature 1040°C

Holding time 20 minutes

Grain size ASTM 10

M

s

M

1 10 100 1000

Cooling time between 800 and 500°C

F + K

10000 100000

Seconds

Fig.

89 Time-Temperature-Transformation diagrams for X20 and P91

111



900

800

Stress

700

MPa)

600

500

400

300

200

100

0

Steel 91

EM12

X20

^Tensile strength MY\e\d strength

Fig.

90 Room tempe rature tensile properties for Steel 91 , EM12 and X20 tubes

112-



Absorbed

energy

(Joules)

250

200

15

1 0 0 -

• T9 1

O E M 1 2

■ X2 0 C r M o V 12 1

100

Brittle

fracture

8 0

( ) 60

H

40

20

0

Temperature (°C)

Fig. 91 Charpy impact test results for Steel, E M12 and X20 tubes

113



fi 1

Mean coefficient of linear thermal expansion (10 K )

(reference temperature: 20°C)

20

1 0 -

T P 3 1 6 L N

ι 1 1 1 1

1 2 3 4 5 6 7

Temperature (°C)

Fig.

94 Tem perature-d epen denc e of the coefficient of l inear thermal expansion

Thermal conductivity (Wm"

1

/ Κ " ' )

42

3 8 -

3 4 -

3 0 -

26

22 Η

18

14

P 22

Ρ 91

Temperature (°C)

Fig.

95 Temp erature-dependence of thermal conduct ivi ty

115



Material

T91

T91

T91

P91

P91

P91

P91

T91/P91

12CrMoV12 1

Dimensions

4 4 . 5 x 7 . 1

25.4 χ 2.8

45 χ 6.4

3 8 0 x 5 0

3 6 8 x 7 3

3 4 3 x 7 4

2 1 9 x 1 8 . 3

Supplier

A

Β

G

A

C

D

E

Required (ASME)

230 χ 22 F

Required (DIN 17175)

C

0.096

0.084

0.100

0.097

0.110

0.090

0.090

0.08-

0.12

0.18

0.17·

0.23

Si

0.35

0.41

0.33

0.40

0.46

0.46

0.20-

0.50

0.23

0.23

max.

Mn

0.36

0.45

0.37

0.44

0.35

0.47

0.35

0.30-

0.60

0.52

1.0

max.

Ρ

0.020

0.014

0.002

0.016

0.010

0.020

0.014

0.02

max

0.016

0.03

max.

S

0.003

0.004

0.004

0.003

0.001

0.002

0.005

0.01

max.

< 0.005

0.03

max.

Al

0.0102

0.024

0.01

<0.O2

<

0.005

0.008

0.015

0.040

max

Cr

8.36

8.11

8.52

8.33

8.68

8.95

8.38

8.00-

9.50

11.9

10 -

12.5

Ni

0.11

0.08

0.18

0.13

0.12

0.12

0.13

0.4

max

0.66

0.3-

0.8

Mo

0.928

1.03

0.93

0.92

0.93

0.95

0.95

0.85-

1.05

0.92

0.8-

1.2

V

0.20

0.20

0.22

0.24

0.21

0.18

0.22

0.18-

0.25

0.26

0.25-

0.35

Nb

0.07

0.072

0.08

0.08

0.09

0.074

0.10

0.06-

0.10

Ν

0.0584

0.053

0.034

0.039

0.068

0.043

0.030

0.030-

0.070

Fig.

98 Chem ical comp osition of the investigated materials (weight %)

Weld Process

GTAW

GTAW

S M A W

S M A W

SAW

GTAW

GTAW

SMAW

Type

9Cr unmod.

9Crmod.

9Crmod.

9Crmod.

9Cr mod.

12Cr

Nickel base

12Cr

Diameter

(mm)

2.0

1.6

3.2

4.0

3.2

2.0

2.0

2 . 5 - 4 . 0

C

0.063

0.08

0.06

0.OB

0.08

0.21

< 0 . 0 3

0.18

SI

0.62

0.20

0.33

0.36

0.16

0.40

< 0 . 2 0

0.3

Mn

0.54

0.95

1.52

1.56

1.67

0.60

3.0

0.8

Ρ

0.008

0.ΟΌ6

0.005

0.005

0.007

—

< 0 . 0 2

—

S

0.011

0.006

0.003

0.003

0.005

—

<o.oi

—

Cr

6.85

9.02

9.25

9.40

8.93

11.3

1 9 - 2 2

11.0

Ni

—

0.69

0.92

0.90

0.60

—

> 6 7

0.5

Mo

0.99

0.90

1.07

1.07

0.89

1.0

—

0.9

V

_

0.18

0.17

0.17

0.27

0.3C

—

0.3

NO

—

0.04

0.03

0.03

0.06

—

—

—

Cu

—

0.15

—

_

0.02

—

<02

—

W

—

—

—

—

—

0.45

—

0.5

Nb/Ta

—

—

—

—

—

—

2 - 3

-

Fig.

99 Chem ical comp osition of welding consum able (weight %)

__LJ

Fig.

100 Location and loading rig for creep testing ring-type specim ens

117



100000 j

Creep rupture

time

(h)

ASME (P=T (30+log t) = 30.49)

• T91-«uppl.A

± P91-euppl.A

■ P91-$uppl. E

♦ T91-»uppL Β

Ο P91-«JppLC

A P91-«uppl. D

* X20-euppl. F

10000 :

1000

100 :

10

580 600 620 640 660 680 . 700

Temperature (°C)

Fig. 101

Results of iso-stress creep tests

on

base materials (100

MPa,

axial specimens)

1 r

minimum

creep rate 0.1

(% / h)

0.01

0.001 =

0.0001

0.00001

0.000001

£

~

• ■ J ^

r^Ltz

• T91-euppl.A

A P91-suppl.A

■ P91-«uppl. E

♦ T91-suppl. Β

O P91-euppl. C

Δ P91-8uppl. D

* X20-suppl. F

τ

580 600 620 640

660 680 700

Temp erature (°C)

Fig.

102 Minimum creep rates for base metal T91/P91 and X20 (100 MPa)

118.



100000 3-v

creep rupture

t ime

(h)

10000 :

1000

100 ;

560 580 600 620 640 660 680 700

Temperature (°C)

Fig.

103 Results of creep tests on weldments

(100 MPa, cross-weld specimens, T91/P91 suppl. A, suppl. F).

100000

10000

Creep rupture

time

h) 1000

100 =■

\

\

\

v

.

\

\

\

■ T91 -unbent axial specimen

♦ T91-unbent ring specinen

A T91-ring specimen r - 60, no PB HT

Τ T91-ring specimen r- 60,

+1

h/ 740 °C

■ T91-ring specimen r ■ 70, no PBHT

O T91-ring specimen r - 80, no PBHT

+ T91-ring specimen r - 80, + 1h / 740 °C

* X20-unbent ring specimen

D X20-ring specimen r - 6 0, no PBHT

560 580 600 620 640 660 680 700

Temperature (°C)

Fig.

104 Results of iso-stress creep tests on cold formed bends

(100 MPa, supplier A, 44.5 χ 7.1)

119·



T91/P91

X20

ASME

BASE METAL (axial specimen)

WELDMENTS (cross-weld specimen )

GTAW (44.5x7.1)

SMAW (380 χ 50)

SAW (380 χ 50)

BASE METAL

(44.5 χ 7 .1 .

nng specimen)

BENDS (cold lormed, 44.5 χ 7 .1 , ring specimen)

R

=

60,

no

PBHT

R = 60,-i- 740'C/lh

DIN 17175

BASE METAL (230x22. axial specimen)

WELDMENTS

SMAW

(230 χ

22, cross-weld specimens)

BASE METAL

(44.5 χ 7 1, nng

specimen)

BENDS (cold lormed.

44.5 χ

7.1,

nng

specimen)

R

=

60.

no

PBHT

Rupture Time

lrV6O0V100MPa

(n)

85.000

30.000

3.000

6.700

5.000

20.000

2.000

3.500

5.000

7.000

1,910

2,650

850

P L M ( ' )

30,49

av.

30.11

29.22

29.50

2942

29.95

29.07

29.30

29.40

29.55

29.05

29.18

28.75

Rupture suengtn

6O0V1o5h(MPa)

98

87

73

77

75

84

73

76

59

62

49

52

46

Ratio

(%) Ol

rupture

strength: processed

metal· ' base metal

100

84

89

36

100

37

90

100

79

100

38

f

1

)

P

LM - T | 30* log

t ) 1 0

-3

c»lcul«l»i3 Ironi lp / 600 'C / 100 MPi

Fig . 105 Ex t rapo la ted resu l t s of iso-s t ress creep tes ts on b a s e m e t a l , w e l d m e n t s and

b e n d s

4 0 0 -

St ress

(MPa)

1 0

H

^o■

J

Time in h

Fig. 106 Comparison of creep strength values of P91 and X20

- 120



Stress

(MPa)

200

100

60

40

20

■

550 °C

X

•

A

«à

X

i· '

M"~~

1 ,.

i

600 °C

+

o

♦

1 1—ι ι ι

ι

ι i l 1

1 '

Initiai pipe

Beno-pressure zone

Bend-tension zone

Scatterband 1991

1 , ι .1

-£ *

J ^

1 T » * * * - - .

1 *

10

100 1000 10000 100000

Time (h)

Fig. 107 Creep strength of a 90" inductive bend in P91 with R/D = 2.5

(Dimensions 380 χ 50 mm)

X 20 C rM o V 12 1

X 1 0 C r M o V N b 9 1

200

Position

Position

O : Base ma terial A ■ HAZ · : Weld metal

Fig.

108 Hardness profiles of weld in X20 and P91

- 121



200

Stress

(MPa)

1 0 0

60

40

20

550°C·

6 00 °C — - , _

650°C — — ^

Base material (BM)

• A Failure in BM

O D Failure in HAZ

π

•

D

τ — Γ Τ Τ Τ Τ Τ

o -

-^

D

α o>

D->

Δ *

O ^ N

10 100 1000

Time (h)

10000 100000

Fig.

109 Creep strength of the weldment P91

150

Stress

(MPa)

100

X 10'OOOh

at 600°C

O 100'OOOh

at 600°C

6 8 10 12 14

Vanadium / available nitrogen ratio

Fig.

110 Effect of V:N ratio on Stress Rupture Strength

122



10*-' h

creep

strength

MPa)

300

200

1

^

S

* N - >

^ v ^

—: MFI-Evaluation 92

- -: M FI-Evaluation 91

—: ORNL-Parameter 90

I

5 * 5 ^

- >■*-·

- ^ * ^ w

- s * · ^ ^

^ * * > ^^

I

500 550 600

Temperature

°C

650

Fig. 113

Creep strength

of P91 X

10 CrMoVNb

91)

Fig. 114

Comparison

of

dimensions

of

a T-piece

of

X20

or P91, for

operation

at

585°C

and 300 bar

124



T91 RM. Specimen Header Section

Length

O.D.

I.D.

608 mm

300 mm

212 mm

Conn ect ions 34 mm and 100 mm

Spacing 20° and 48 m m

* »

Fig.

115 Powder metallurgical^/ manufactured header section

Sleet

X 2 0 C r M o V 1 2 1

(DIN 17175)

Header 1

P M - X 2 0 C r M o V 1 2 1

Header 2

T91 mod.

(ASME SA 213 )

PM-T91 Header

X I O C r M o V N b N

C

0.17

0.23

0.19

0.24

0.08

0.12

0.12

Mn

—

1.00

0.56

0.61

0.30

0.60

0.46

Ρ

—

0.030

0.017

0.02Z

—

0.020

0.017

S

—

0.030

0.013

0.006

—

0.010

0.010

Si

—

0.50

0.26

0.28

0.20

0.50

0.37

Cr

10.00

12.50

11.65

12.10

8.00

9.50

9.00

M o

0.80

1.20

1.05

1.10

0.85

1.05

1.00

Ni

0.30

0.80

0.42

0.40

—

0.40

0.24

V

0.25

0.35

0.29

0.30

0.18

0.25

0.25

N b

—

—

0.14

0.06

0.10

0.12

Ν

_

—

0.046

0.030

0.070

0.063

Heat t reatment

1 0 2 0 - 1 0 7 0 ' C

730 - 780 'C

1 0 5 0

-

C / 5 h

1 0 2 0 ' C / 4 h

770"C / 5h

> 1 0 3 7 X

> 732" C

1 O 6 0 " C / 1 . 5 h

760'C / 3 h

Fig.

116 Com positions of PM header sections

125



100

Total strain

Etot = 0.001s-

1

; R = -1

range

(%)

10

:

1

:

0.1

:

-

-

-

-

* f c *

: φ * * * * * * *

* * *

♦

:

.

X20 CrMoV 12 1 (tube)

Δ

25*C

A 550°C

Φ

°'aå

J

fa*

■ft

-

Y

550* 0X2 0 CrMoV 121 , Bendbk

(19Í

-

Q

55 0* CP 91 Bendick (1991)

. X 530 *0X 22 CrMoV

121,

Obst (191 8)

+ 90 *C X2 2 CrMoV 12 1. Lachmann (1987)

* 25 *0 X2 0 CrMoV 12

1,

Maile (1987)

1

1—ι—ι ι ι 111 1 1—ι—1

l i n i

1 1—ι—Γ

Ί

γ

M I M —

X20 CrMoV 12 1 (welded)

»Nipple

P/M X20 CrMoV

12

1 (Header)

o25"C

• 550*C

O Nipple

T91 (Header)

D 25"C T91

■ 550 *C T91

<*

9 +

X

▼ Û

10 100 1000 10000 100000 1000000

Cycles to fai lure

Fig.

120 Fatigue strength of PM and conventional steels (X20 and P91)

35

MPa

| 2 1 / 4 P - 1 M Õ 1

P/T22

| 9 Cr-1 Mo "~r^-

T 9

C

Mo

6 0

- 80 MPa

I 9 Cr-2 Mo |

MCM9M

V

Nb

8 0 - 1 0 0

MPa

>j 9 Cr-2 MoVNb |

EM12

■

V

Optimized

N b ,

.

V,

Nb . ,+ W

— H 9 Cr-1 MoVNb | —*\ 9 Cr-1 MoVNb \-

F9

P/T91

1 2 0 - 1 4 0

MPa

(Expected)

—H 10Cr-1Mo-1WVNb |

E911

Mo

>J9Cr-0.5Mo-1.8WVNb|

NF616

■Ni

12Cr | *

M o

> Π 2 ^ Ό Ι Μ Ο ~ 1

410

12O-0.5M0-1,8WVNbN j

TB12M

- C

-Mo

♦ W +W

I 1

+ N b

J 1

+ C u

I 1

I 12 Cr-1 MoV | >| 12Cr-1Mo-1 WVNb | Η 12Cr-0.5Mo-2WVNbCu |

X20CrMoV121 MCM 12 HCM12A

Fig.

121 Developm ent Progress of 9 - 1 2 % Cr Steels (after Mayuama)

- 128



Charpy impact

energy at 20°C

(J/cm2)

300

2 0 0 -

1 0 0 -

P91 (After Iseda et al)

Aged at 600°C

10

100 1000 10000

Aging time (h)

100000

Fig. 126 Change in Charpy Impact Energy with Aging Time for NF616 pipe and tube

Stress

(MPa)

100 1000

Time to rupture (h)

10000 100000

Fig. 127 Creep properties of NF616 Tubes

-



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European Comm ission

EU R I 6858 — N ew s teels and manufacturing processes for critical components in advanced

steam pow er plants

K. H Mayer C Berger ft B Scariin

Luxembourg: Office for Official Publications of the E uropean C omm unities

1996

—

133

pp. — 17.6 x 25.0 cm

Physical sciences se ries

ISBN 92-827-6578 -4



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