Study of Modified 9Cr-lMo Welds Xiaotian Li1)，M. T. Cabrillat2)，Y. Lejeail2)

(1. Institute of Nuclear and New Energy Technology，Tsinghua University，

Beijing 100084，China 2. CEA Cadarache, DEN/DER/SESI, France)

Abstract ： Modified. 9Cr-lMo is the best candidate for higher service

temperature. It has attractive properties ： high creep strength with

good ductility, high resistance to cracking, high thermal conductivity

and low thermal expansion coefficient. This paper reviews the main

features concerning modified 9Cr-lMo welding. It is obtained that as

far as selecting optimum preheat temperature and suitable P W H T ,

controlling the chemical composition of weld metal, good material

properties of Mod. 9Cr-lMo weldment will be obtained.

Key words: Weld metal, Post weld heat treat, Heat affected zones,

Cracking

1. Introduction Since, modified 9Cr-lMo has high thermal conductivity, low thermal

expansion, high strength and resistance to corrosion at elevated tefnperature，it

has been considered as the best candidate for high temperature conditions.

Modified 9Cr-lMo steel is a relatively new structural alloy that was

originally developed for use as a steam generator material for advanced fast

breeder reactors in the United States. It is a ferritic steel and micro alloyed with

columbium and vanadium with a controlled nitrogen content. The alloy is

currently under active consideration for use as a steam generator alloy for

advanced breeder reactors in Japan； it is used world wide in the power

generation industry for superheaters and other applications requiring piping and

tubing applications for prolonged service to temperatures of about 600° C. It is

also used in the petrochemical industries for such high temperature applications

as distillation, cracking and hydro-treating units.

This study will point out the main features concerning modified 9Cr-lMo

welding by analysing many references.

64

2. Chemical composition

The improvement in propert ies of modif ied 9 C r - I M o are achieved by

control led addit ions of vanad ium，n iob ium，and n i t rogen，combined w i t h a

normalize-and -temper heat t reatment. The chemical composit ion of modif ied

9 C r - I M o is indicated in Tab. 1.

Tab. 1 Chemical composition mass fraction of Modified 9Cr-lMo

c Mn Si P S Cr Ni M o V Al Nb Cu N

0. 08- 0. 30- 0. 20- max max 8. 0- max 0. 85- 0. 18- max 、0. 06- max 0. 03-

0. 12 0. 60 0. 50 0. 015 0. 01 9. 5 0. 40 1. 05 0. 25 0. 04 0. 10 0. 1 0. 07

Th is steel is also called 9Cr - lMo~V，and is incorporated into A S M E as SA-

213-T91 for tubing，SA-387-Grade 91 for plates, SA-335-P91 and SA-369-FP91

for p ipe，and SA-182-F91 and SA-336-F91 for fo rg ings，or in Europe as

X 1 0 C r M o V N b 9 - l .

The high ( 9 % 〜 1 2 % ) chromium steels are either fu l l y austenitic or have a

duplex (austenite plus 8 —fer r i te ) s t ructure at austenit izing temperatures in the

range 850 to 1 200 °C. The austenite phase transforms to martensite dur ing air

cooling or rapid quenching to ambient temperature, and the steels are

subsequently tempered to obtain a good combinat ion of s t rength, duct i l i ty，and

toughness. The ferr i te phase inhib i ts austenite grain g r o w t h , but i t adversely

influences the st rength and toughness.

Modi f ied 9Cr_ lMo requires a normalize and temper heat t reatment to fu l l y

develop its st rength and toughness properties. Standard practice is to normalize

between 1 038 to 1 090 °C to avoid objectionable grain g r o w t h , and tempering at

760 to 788 °C which al lows carbides to precipitate homogeneously w i t h i n the

tempered fu l l y martensit ic s t ructure ( see Fig. 1 ) . Transmission electron

microscopy reveals a complex microstructure consisting of a h igh dislocation

density and subboundaries decorated w i t h carbides [ 1 ] . The subboundaries are

stabilized by the precipi tat ion of M C and M23 C6 precipitates which accounts for

the alloy，s increased strength.

65

The lower cr i t ical temperature A c l ( the temperature at wh ich the a ^ 7

t ransformation commences on heating) for modified 9 C r - l M o is determined to be

between 830 °C to 850 °C ； the upper crit ical Ac3 ( the temperature at wh ich the a

一 7 t ransformat ion is complete) exists between 900 to 940。C • The single phase

tempered martensite structure exhibits opt imum strength and toughness

characteristics compared to two phase ( tempered martensite and S fe r r i t e )

structure. For this reason i t is imperative that modif ied 9Cr_ lMo not be

subjected to temperature above A c l after heat treatment dur ing hot fo rm ing ,

bending, or any other heat treat ing operation. But heating above the cri t ical

temperature is impossible to avoid in weld regions.

F i g , 1 T y p i c a l t e m p e r e d m a r t e n s i t i c m i c r o s t r u c t u r e

o f m o d i f i e d 9 C r - l M o s tee l

The high-chromium martensit ic steels are generally regarded as being more

di f f icul t to weld than austenitic steels, because i t is often necessary to preheat

before welding to avoid cracking and i t is essential to post weld heat treat

( P W H T ) to temper the br i t t le martensit ic structures formed in the fusion zone

( F Z ) and heat affected zones ( H A Z ) .

The various zones in a h igh-chromium martensit ic steel fusion weld jo int are

shown schematically in Fig. 2.

FZ ( T 〉 T m ) : The f i rs t phase to fo rm during solidif ication of the mol ten

weld is 8 - ferr i te； the ferr i te to austenite t ransformat ion occurs on further

cooling，and austenite transforms to martensite on cooling below temperatureM s

66

( M s is martensi te start temperature, M f is martensi te f in ish temperature) . Some

8-ierrxte is usual ly retained in the FZ at ambient temperature, even when there is

no ferr i te present in the base and f i l le r w i re materials， as complete

t ransformat ion to austenite does not occur dur ing cooling at the fast rates typical

of the weld ing process [ 2 ] . Since the ^ - f e r r i t e can have detr imental effects on the

mechanical propert ies，part icular ly s t rength and fracture toughness，the content

should be control led at ^ 3 % by balancing the concentrations of the austenite

and fer r i te fo rming elements in the base metal and f i l ler wire.

H A Z - r e g i o n l ( T m 〉 T > ) : Th i s region consists of martensi te and

(Merr i te . The ferr i te is fo rmed along the pr ior austenite grain boundaries as the

region is heated into two phase f ie ld dur ing welding； some of the d - fe r r i te is

again retained at ambient temperature in a band typical ly 0. 3 to 0. 5 m m wide

adjacent to the fusion lines as a resul t of the rapid cooling after we ld ing [ 2 ] .

HAZ- reg ion2 ( TyS > T > A c3 ) ： The microst ructure is fu l l y martensit ic.

Th i s region is heated into the higher temperature part of the phase f ie ld dur ing

we ld ing , and the or ig inal carbides part icles are dissolved, resul t ing in coarse

pr ior austenite grain and martensi te la th structure.

HAZ - reg ion3 ( TrS > T > Ac3 ) ： The structure of this reg ion, wh ich is

heated into the lower temperature part of the / -phase f ie ld , is again martens i t ic ,

but i t is f iner grained than region 2，as some of the or ig inal carbides are not

dissolved and inh ib i t gra in g rowth .

HAZ- reg ion4 ( A c 3 〉 T 〉 A c l ) ： The st ructure consists of untempered and

overtempered martensite. Incomplete t ransformat ion to austenite and addit ional

tempering of the or ig inal tempered martensi te st ructure of the base metal occur

dur ing heating in this in tercr i t ica l zone, w i t h the austenite again t rans forming to

martensi te on cooling.

HAZ- reg ion5 ( A c l 〉 T 〉 T t ) : The or ig inal tempered martensi te in this

nar row zone is fur ther tempered dur ing we ld ing , but the microst ructure is

otherwise similar to that of the base steel.

A typical microst ructure of modif ied. 9 C r - l M o steel weld is i l lus t ra ted in

f ig3. Four regions are delineated - the We ld metal ( W ) corresponds to the FZ of

Fig. 2，the Transformed Zone ( T Z ) corresponds to regions 1 , 2 , 3 , 4 of Fig. 2，

the tempered zone ( T M P Z ) corresponds to region 5 of Fig. 2，and base metal

( B M ) is the base steel wh ich was unaffected dur ing the weld ing process.

67

=/①m鋼r Fusion Zone ( F Z ) ： T〉'Tm

H e a t - A f f e c t e d - Z o n e ( H A Z ) [as-welded]：

Region 1 Tm>T>Tr3 Martensite+61

Region 2 Tyd>T>Tc3 Coarse grained Martensite

Region 3 Tyd>T>Tc3 Fine grained Martensite+厶

Region 4 A c 3 > T> T c i Martenslte+ Overtempered Martensite

Region 5 T c i > T > T r Overtempered Martensite

where T = temperature achieved during welding

Tm = melting point of steel

TrS = temperature at which y^d transtormation is complete on heating

Tt = original tempering temperature of steel

A c i = temperature at which transformation starts on heating

A c 3 = temperature at which transformation is complete on heating

F i g . 2 S c h e m a t i c d i a g r a m o f t h e h e a t - a f f e c t e d zone r e g i o n s

i n a f u s i o n w e l d o f h i g h - c h r o m i u m m a r t e n s i t i c s tee l

F i g . 3 T y p i c a l m i c r o s t r u c t u r e o f a m o d i f i e d 9 C r - l M o s tee l w e l d

i l l u s t r a t i n g t h e w e l d ( W ) a n d t w o r e g i o n s o f t h e H A Z ( t r a n s f o r m e d zone ( T Z )

a n d t e m p e r e d zone ( T M P Z ) )

The microhardness profi les across G T A W (Gas Tungsten Arc Weld ing)

welds in modif ied 9Cr - IMo [3 ] is shown in Fig. 4. The data in Fig. 4 show that

the various regions of H A Z can be differentiated by their hardness. The

presence of the softer 8 - fe r r i te phase in the untempered martensite ma t r i x is

responsible for the reduct ion of hardness in H A Z region 1 (of Fig. 2 ) . Since the

hardness of martensi te increases w i t h increasing carbon content , the d issolut ion

of the carbide particles in the austenite in H A Z region 2 has resulted in the

format ion of h igh carbon martensi te w i t h max imum hardness. The carbon

content of the martensi te in H A Z region 3 is reduced relat ive to H A Z region 2

due to the incomplete d issolut ion of the carbides at the lower austenit iz ing

temperature and hence the hardness is lower. The hardness reduct ion in H A Z

regions 4 and 5 resul t f r om overtempering of the or ig inal microst ructure at

temperatures between Ac3 and A c l and below A c l， respect ively.

A f t e r P W H T at 732 °C for l h , the hardness of the FZ and H A Z are reduced

as a resul t of the temper ing react ions, wh ich include a reduct ion in dislocat ion

densi ty , la th break up and development of a polygonized s t ruc ture , reduct ion in

solid so lut ion strengthening due to prec ip i ta t ion, and coarsening of precipitates.

I t is impor tan t to note that the hardness of the H A Z approaches that of the base

steel after P W H T for t imes up to 80 h at 732。C. See Fig. 5 [ 3 ] .

4. Weldablity for thick section

I n general, the h igh al loy content makes i t more hardenable in the we ld

region, and therefore more susceptible to we ld cracking. Mod i f ied 9 C r - I M o has

proven to be h igh ly weldable, provided careful a t tent ion is paid to preheat,

post-heat, and consumable control . Many tests are per formed, weld ing

modif ied. 9 C r - I M o plates w i t h thickness 25 m m , 51mm，203 m m , modi f ied

9 C r - l M o pipes w i t h 於76X13，於 273X45，於 350X45，於 2 6 0 X 6 0 and so on. I t can

be obtained satisfactory mater ia l propert ies by suitable weld ing process.

4 .1 Welding method

A l l of the t rad i t ional arc weld ing methods of jo in ing steels except f lux-cored

weld ing are appropriate for the jo in ing of modif ied 9 C r - l M o material . I t

includes, Shielded Me ta l A r c Weld ing ( S M A W ) , Gas Tungsten A r c Weld ing

( G T A W ) , Submerged A r c Weld ing ( S A W )， M a n u a l Meta l A r c Weld ing

( M M A ) .

Advanced, low-heat input and high-speed processes such as E lect ron Beam

weld ing ( E B ) and laser are also being developed. Th is k ind of process has

several at tract ive characteristics for weld ing of th ick sections. The heating and

69

400

Standard 9 C r - l M o

• A S -We lded O P W H T : 7,32°C for 1 h •

200

150

W e l d Me ta l

F i g . 4 M i c r o h a r d n e s s t r a v e r s e a c r o s s a G T A w e l d b e f o r e a n d a f t e r P W H T

a t 732 °C f o r l h ； s t a n d a r d 9 C r - l M o s tee l base a n d f i l l e r w i r e

cooling rates are much higher compared to arc weld ing, so narrower weld zones

are produced and thermal damage to the material adjacent to the welds is

therefore minimized. But weldment in modif ied 9 C r - I M o produced by EB contained large porosity

presumably f rom its relatively high ni trogen c o n t e n t w . Ni t rogen content has

been shown to affect the occurrence of porosity in other steels. Welding

parameter variations may reduce or eliminate this porosity but i t has not been

evaluated.

4. 2 Welding consumable The design of appropriate welding consumables is important. On many

occasions, i t is wise to make the composit ion of welding consumable as similar

to the parent steel type as possible. However , for the modif ied 9Cr - lMo， i t is

realised that adjustments to the Nb and N i content are crucial to obtain

acceptable toughness values, give rise to significant changes in microstructure,

hardness, room temperature tensile strength and creep rupture duct i l i ty.

The weldments of modif ied 9 C r - l M o w i t h fu l ly martensite structure which

70

L r

o o

o 5

3 2

(qdp) SS3UPSH

O 2 4 6 8 10 12 14 16 18 20 22 24 26 28 30 32 34 Distance ( m m )

F i g . 5 M i c r o h a r d n e s s t r a v e r s e ac ross t h e s u b m e r g e d a rc w e l d b y s t r u c t u r e w e l l s .

T h e base p l a t e w a s m o d i f i e d 9 C r - l M o a n d t h e f i l l e r w i r e w a s s t a n d a r d 9 C r - l M o .

T h e p l a t e w a s p o s t e d hea t t r e a t e d f o r 2 , 1 0 , 2 0 , 4 0 , a n d 80 h at 732 °C

are subjected to elevated temperatures and rapid cooling frequency contain small

amounts of 8 - ferr i te in both the weld metal and Heat Af fected Zone ( H A Z ) . I t

has been established that d - ferr i te has a number of detr imental effects on the

properties of modif ied 9Cr_ lMo， i nc lud ing reduction of creep duct i l i ty and

toughness of weldments. So the compositions of modif ied 9 C r - l M o consumables

should be selected to minimise d - ferr i te format ion. Panton-kent7 s wo rk gives

the expression to estimate the amount of retained ^ - fe r r i t e in modif ied 9Cr-.IMo.

F F is called the Kaltenhauser ferr i te factor.

F F = % C r + 6 % S i + 8 % T i + 4 % M o + 2 % A l + 4 % N b - 2 % M n - 4 % N i - 4 0 %

( C + N )

Increasing FF gives rise to an increase in d - ferr i te content, fu l l y martensite

structure being obtained by F F value below approximately 8. d - ferr i te content is

also dependent on welding condit ions, w i t h increasing arc energy and preheat

temperature causing a reduct ion in weld metal 8 - ferr i te level.

For G T A W weld ing, there is no need to modi fy the composit ion of weld

metal. Weld wire meeting composit ional ranges similar to those of the base

O

O

O

O

O

O

O

O

O

O

O

A

2

0

8

6

4

2

0

8

6

4

2

A

3

3

2

2

2

2

2 1

1

1 1

1

(qdp) SS3UPS-BHOJ0I5

71

mater ial specif ication, i t results in welds w i t h acceptable weldabi l i ty and weld

properties.

For SMAW[5—6] S A W [ 5 ] ， M M A , some fur ther modif icat ions of composit ion

is necessary to achieve an acceptable balance of weld metal creep st rength and

room temperature notch toughness properties.

N iob ium is impor tant in achieving creep s t rength , but is detr imental to the

toughness properties and should be at the lower l im i t of the specifications.

N icke l has a favourable influence on the duct i l i t y and should be at the upper

l im i t of the specifications. I t has been suggested commonly that the N b content

is lower than in base meta l , is l im i ted to 0. 04% 〜0 . 0 8 %， a n d N i content is

higher in weld metal than in base meta l , up to 1%. On the other hand the

M n + N i contents is not al lowed to exceed 1. 5%C6]，because not only N i but also

M n influence the lower t ransformat ion point A c l .

T w o consumable choices for weld ing P91 by Esab[7] are one M M A welding

electrode designated O K 76. 98 and one solid wi re called O K T ig rod 13. 38.

Tab. 2 Chemical composition mass fraction of OK 76. 98 and OK Tigrod 13. 38

C Mn Si P s Cr Ni M o V Nb N

OK 0. 08- 0. 40- 0. 20- max max 8 . 0 - 0. 4- 0. 85- 0. 15- 0. 04- 0. 03-

76. 98 0. 13 1 . 0 0. 50 0. 02 0. 02 10. 0 1. 0 1. 1 0. 30 0. 08 0. 07

OK Tigrod 0. 08- 0. 35- 0. 20- max max 8. 6- 0. 6- 0. 85- 0. 18- 0. 04- 0. 03-

13. 38 0. 12 0. 60 0. 50 0. 01 0. 01 9. 3 0. 9 1. 05 0. 25 0. 08 0. 07

Tab. 3 Chemical composition (mass fraction)of the pure weld metal[6]

C Mn Si P S Cr Ni M o V Al Nb N

0. 08- 0. 40- 0. 20- max max 8. 0 - 0. 4- 0. 85- 0. 15- max 0. 04- 0. 03-

0. 12 1. 0 0. 50 0. 02 0. 02 10 .0 1. 0 1. 10 0. 30 0. 04 0. 08 0. 07

The weld metal in reference [ 6 ] for S M A W is as follows：

The reference [ 8 ] investigates 12mm th ick plates of modif ied 9 C r - l M o steel

which are welded by M M A using three di f ferent consumables： standard

9 C r - l M o steel (P9)，modi f ied 9 C r - l M o ( P 9 1 )，a n d nickel-base al loy Inconel 182.

The toughness in the case of P9 is not superior to that of P91 when the lat ter is

used. I n the case of the P91 weld meta l , acceptable fusion zone mechanical

72

propert ies, especially duc t i l i t y and toughness, are obtained only after a postweld

tempering t reatment at 760 °C for 6 h. The relat ively poor duc t i l i t y and

toughness of the P91 welds is most l i ke ly a consequence of the presence of small

amounts of fer r i te in local regions of the we ld metal microstructure. T h e ferr i te

islands are believed to occur because of segregation effects associated w i t h the

in t roduct ion of a l loy ing elements th rough the f l u x coating. I t is suggested to use

f ine ferroal loy part icles in the f l u x coating. The weld ing posi t ion has inf luence on

toughness. By experience the best toughness propert ies are achieved i f the

weaving technique is used instead of the usual str inger bead welding1 6 '9 ] .

4 . 3 PWHT

The A S M E Boi ler Code, Section I X has defined 9 C r - l M o - V as P-number 5，

Group 4 material. Section I of the code stipulates that P W H T at 704 °C minimum is

mandatory after any and all welding to this alloy. I t is prudent to require that sections

greater than 1/2 inch in thickness be given a P W H T in a temperature range of 760 °C

to 788 °C. H igher P W H T r isks the chance of v io lat ing the lower t ransformat ion

temperature wh ich is as l ow as 830 °C. Mos t fabricators specify 760 °C for 2h ,

but indiv idual specifications require lower temperature for longer t ime.

4 . 4 Preheat

Comprehensive we ldab i l i t y test programs are conducted using Y-groove

restraint tests to study the effects of preheat temperature. No cracking is

observed in test welds preheated to 204 °C. A l t h o u g h the A S M E does not

specif ically address i t , i t is recommended that preheat be maintained at a

m i n i m u m temperature of 204 °C on sections thicker than 1/2 inch. For thinner

sections, 149 °C m i n i m u m is acceptable. For jo ints in excess of 1 /2 inch in

th ickness, preheat temperature should not be al lowed to drop below the

m i n i m u m preheat temperature un t i l at least two- th i rds of jo in t is completed, and

then only after either a post bake at 204 °C for four hours fo l lowed by an air

cool，or an instantaneous intermediate post we ld heat t reatment at a temperature

of 704 °C to 732 °C is performed. I n the next section, i t can be seen that the

preheat temperature is best to be selected w i t h mutua l consideration of u l t imate

tensile s t reng th , specific e longat ion and martensi te content.

73

Cracking during or after welding

The weld ing of the 9%、〜12% chromium steels requires a h igh degree of

preparat ion and cont ro l to avoid cracking dur ing the we ld ing , P W H T ， o r

service. Cracking processes include ： sol id i f icat ion c r a c k i n g，H A Z l iquat ion

cracking，hydrogen or cold cracking，reheat cracking and Type I V cracking.

5 . 1 Solidification cracking

The results of an extensive program using the T igamaj ig test on many

exper imental heats of modi f ied 9 C r - l M o steel have indicated l i t t l e or no

suscept ib i l i ty to sol id i f icat ion cracking.

5. 2 HAZ liquation cracking

The h igh ch romium martensi t ic steels are thought to be proned to l iquat ion

cracking in the H A Z immediate ly adjacent to the FZ. I n the coarse-grained

region of the H A Z , l iquat ion cracks can develop as a resul t of acting stresses and

impu r i t y segregations on grain boundaries, fo rming low mel t ing phase.

Temperatures are determined at wh ich the duc t i l i t y ( reduc t ion of area) and

s t rength drop to zero - so called n i l -duc t i l i t y ( T N D ) and n i l -s t rength ( T N S )

temperatures. The difference AT 二 TN S — TN D is a measure for l iquat ion cracking

sensi t iv i ty.

Reference [ 1 0 ] studies the propert ies of P91 steel in the f o r m of a pipe

於275X45 by M M A bu t t weld ing. The test in [ 1 0 ] can be regarded as very

resistant to l iquat ion cracking for AT ^ 0 •

5. 3 Hydrogen or cold cracking

Cold cracking of the b r i t t l e martensi t ic phase can occur dur ing cool ing after

weld ing of low and h igh al loy martensi t ic steels, par t icu lar ly in th ick sections，

as a resul t of the stresses induced by thermal condi t ion and volume expansion

associated w i t h the austenite to martensi te t ransformat ion. The incidence of

hydrogen or cold cracking is min imized by the use of covered electrodes w i t h low

hydrogen contents and protect ion of the wires and electrodes by storage at

elevated temperatures pr ior to weld ing. A n effective way of prevent ing cold and

hydrogen-assisted cracking is by cont ro l l ing the preheat, interpass，and P W H T

temperatures.

I n general, preheat temperature and interpass temperature is 200 to

300。C[8]. Reference [ 1 ] indicates that the preheat temperature has inf luence on

mater ia l crack suscept ib i l i ty because the martensi te content is contro l led by the

74

preheat, investigates the relationship between preheat, interpass and M s

temperatures, and recommends to calculate the preheat temperature w i t h the

chemical composit ion of steel in order to attain the opt imum strength parameters

of the jo int .

The relationship between Ms and composit ion of steel is as fo l lows

(concentrations in w t % )：

M s = 454 - 210 • C + ^ - 27 • N i - 7. 8 • M n - 9. 5

• (Cr + Mo + V + W + 1. 5 . Si) — 21 • Cu

From the fo l lowing Tab. 4，it can be seen that there is 50 to 60 °C difference

in the M s temperatures，which fo rm the upper and lower content l imi ts.

Tab. 4 composition and Ms temperature of Mod. 9Cr-lMo

C(%) M n ( % ) S i ( % ) C r ( % ) N i ( % ) Mo(%) V ( % ) N b ( % ) M s ( ° C )

0. 08 0. 30 0. 20 8. 0 0 . 2 0. 85 0. 18 0. 06 393

0. 12 0. 6 0. 50 9. 5 0 . 4 1. 05 0. 25 0. 1 339

I f the preheat temperature is chosen independently f rom the composit ion of

the steel heat, in the case of welding steels at the composit ion l im i ts , there may

be a 50% 〜 6 0 % difference in martensite content. The steel w i t h the higher

martensite quant i ty may crack.

The austenite present at the austenitizing temperature should t ransform

fu l l y to martensite on cooling. Neither a high preheat temperature is favourable

because of the crack sensit iv i ty dur ing cooling , nor one that is too low because

of the r isk of cracking dur ing welding due to excessive martensite format ion. If

the preheat temperature is above the M s temperature (austenitic welding)， there

is no crack susceptibi l i ty un t i l welding ceases, but considerable austenite

transforms into martensite dur ing cooling. As a resul t , h igh residual stress can

result in cracking. I f martensite welding is carried ou t , some martensite has

already formed dur ing welding. For a correctly selected preheat temperature,

there is a considerable amount of ducti le austenite that is the reason for the

lower crack susceptibi l i ty. Since martensite welding gives lower stress levels

which tends to reduce crack suscept ib i l i ty , i t is commonly preferred to austenitic

welding.

75

1 600

1 200

800

M s ( M s --50) (Ms—100) (Ms—150)

Temperature be low Ms , (Ms—7), (°C)

400 (Ms—200)

F i g . 6 M a r t e n s i t e f r a c t i o n , e l o n g a t i o n a n d t e n s i l e s t r e n g t h o f a 0. l C - 9 C r - l M o - t y p e

c r e e p - r e s i s t a n t s tee l c o o l e d f r o m t h e a u s t e n i t i z i n g t e m p e r a t u r e

t o t h e t e s t i n g t e m p e r a t u r e

- x X : ^ j c r ^ J ^ n

5. 4 Reheat cracking

Reheat or stress-relief cracking may occur in the H A Z and sometimes in the

weld of al loy steels dur ing P W H T or service at elevated temperatures. The

cracking results f rom increased 歹olution of al loy carbides in those parts of the

H A Z and weld metal heated to temperatures of ^ 1 320 °C， fo l lowed by strain-

induced precipi tat ion of f in particles on the dislocations and stacking faults

w i th in the coarse pr ior austenite grains when the residual stresses relax by creep

at temperatures in the range 400 to 750 °C. Th is leads to marked strengthening

of the grains such that the deformation is concentrated at or near the grain

The preheat temperature should be selected w i t h mutual consideration of

ul t imate tensile strength，specific elongation and martensite content. See Fig. 6.

The recommended preheat temperatureD ] when 9Cr~lMo type steel is

welded is as fo l lows：

Tp9Cr

the interpass temperature [ 1 ] is ：

(M s — 9 0 ) ± 1 0 °C

T{ = ( M s — 1 9 0 ) ± 1 0 °C

Temperature below M s , ( M s - 7 ) , (°C)

80 M s ( M s - 1 0 0 ) (Ms—200) (Ms—300)

100

(1)

(2)

2 000

ods)

Mbcis 31111

S o

o

o

8 6

4

2

— I

o o

o

6 4

2

(0/0)

ilow

76

boundaries and can result in low-duct i l i ty? in tergranular fai lures.

I n general，the presence in the steels of n iob ium and vanadium w i l l promote

stress-relief cracking due to the int ragranular precipi tat ion of n iob ium and

vanadium carbides. H o w e v e r , the modi f ied 9 C r - I M o is proved to have resistance

to the reheat cracking15’10’12] • The test in reference [5 ] has been per formed on

modi f ied 9 C r - I M o steel plate w i t h 300 m m th ick and coarse grain area of H A Z is

simulated by temperature peaks at 1 300 and 1 350 °C. The specimen is then

machined as a tensile specimen, and tensile test is per formed after P W H T up to

760 °C. The duc t i l i t y is characterized by the reduct ion of area of the tensile

specimen. I f reduct ion of area > 2 0 % for any temperature peak, the mater ia l is

considered to be not susceptible to reheat cracking. The result shows that the

steel offers a very h igh duc t i l i t y , reduct ion of area is > 7 0 % .

5. 5 Type IV cracking

Th is phenomenon is not f u l l y understood，circumstant ia l evidence suggests

that the cracking results f r o m h igh stresses across the weldment and the

accumulat ion of creep damage in the in tercr i t ica l ly t ransformed zone-region4 of

Fig. 2 of the we ld H A Z , wh ich has a low rup ture duct i l i ty . I t has been

recommended that the modi f ied 9 C r - I M o steel should be par t ia l ly tempered at

600 to 700 °C before weld ing to remove the soft zone? and the effectiveness of

th is approach in prevent ing premature cracking has been demonstrated.

6. Material properties

6 .1 Tensile properties

The m i n i m u m values for y ie ld and u l t imate tensile strengths of modif ied.

9 C r - I M o at room temperature are 414 and 585 MPa. The m i n i m u m value of

reduct ion of area is 55 %. The strengths propert ies of weldments indicated by

the rat io of weldment to base metal and reduct ion of area value are shown

I t can be seen that the y ie ld strengths of al l weldments made by G T A ,

SMA，and SA processes exceeded 1. 10 t imes m in imum values for the base metal

except on ly two points at very h igh test temperatures. U l t ima te tensile s t rength

for al l but two weldments met the base metal m in imum values，they ranged f r om

1 to 1. 35 t imes the base metal m i n i m u m value. The reduct ion of area values for

al l weldments exceeded the m i n i m u m value of 55% for the base metal for al l test

temperatures.

77

F i g . 7 R a t i o ( i ? ) o f y i e l d s t r e n g t h o f aged t o y i e l d s t r e n g t h

o f u n a g e d m a t e r i a l as a f u n c t i o n o f t i m e a n d t e m p e r a t u r e

i n a p a r a m e t e r i z e d ( P ) f o r m f o r t h r e e h e a t s o f m o d i f i e d 9 C r - l M o

The comparision between measured and predicted values to be made in

Tab. 5.

So it is concluded that modified 9Cr - lMo weldments w i l l have no problem in

meeting the base metal.

The influence of prolonged exposure on the elevated temperature short term

tensile properties is also of interest. Brinkman [ 1 ] develops the equations that

would analysis the yield and tensile strengths for aged to unaged material as a

function of a time-temperature parameter P where P is：

P = T( log t+lO)/l 000，

T : t empera tu re (K ) , ^ : t ime(A) .

Ry 二 14. 143 • P- 1 1 0 2 9 (3)

R t 二 12.418 • _P-L。58 4 (4)

R y is the ratio of yield strength of aged to yield strength of unaged material,

R t is the ratio of tensile strength of aged to tensile strength of unaged material.

Predicted values are based on the parametric expression (R equations) given in

Fig. 7-8. Agreement between predicted and measured values is good.

i?:14..143-P-1.1029 P:T( logt+10)/1.000

r :Temperature(k)

/ : t ime(h)

• H E A T 30176

m H E A T 30383

• H E A T 30394 A g i n g and Test Temper ature Were the Same

14

1

.

0

.

9

8

7

6

5

.

4

l.Lo.o.,ao.0,.0,,

(sp&gn Hauss /p&v

llaualls

PR!A

78

Tab. 5 measured and predicted values of yield strength (YS)

and ultimate tensile strength (UTS) at several temperatures (at strain rate of 6. 7X 10"5s_1)

A g m g and 八 ㈣ 7 5 0 0 0 h

testing Measured Measured Predicted

temperatureC °C) Y S ( M P a ) U T S ( M P a ) Y S ( M P a ) U T S ( M P a ) Y S ( M P a ) U T S ( M P a )

482 427 490 428 486 418 470

593 303 324 256 275 256 269

649 208 233 163 187 164 181

649 209 243 170 185 165 189

These equations are then used to estimate the reduct ion in y ield and tensile

strengths that wou ld occur fo l l ow ing prolonged service at elevated temperature.

F r o m fo l low ing Fig. 9-10， i t can be seen that prolonged thermal exposure w i l l

not degrade tensile propert ies below S„ values.

The modi f ied 9 C r - l M o is strain-rate sensitive for aged and unaged material .

The yield and u l t imate strengths increase w i t h increasing st ra in rate, the results

in [13 ] show a st ra in aging peak at temperature of 538 °C to 593 °C.

Results f r om analyt ical electron microscopy are presented wh ich show that

prolonged thermal aging causes some al terat ion in the microst ructure wh ich can

account for the changes observed in the mechanical propert ies. The rma l aging

general ly can cause a number of changes in the microst ructure inc luding

increased precipitate densities of several phases ： M2 3 C 6 ，M C and Laves, and

recovery depending upon the t ime , temperature and specific composi t ion of heat

involved. Laves phase ( p r i m a r i l y S i - , M o - , Fe- and in some cases P - r i c h )

precip i tat ion wh ich is not found in unaged materials occurred over the

temperature range of 482 °C to 649 °C. Laves phase is normal ly thought of as an

embr i t t l i ng component , the presence of wh ich can reduce room temperature

toughness and long te rm creep-rupture duct i l i t y . Recovery, for example,

reduct ion in dislocat ion density and sharpening of subgrain boundaries occurred

par t icu lar ly w i t h longer t imes and at the higher temperatures. Some fine

precipitate d issolut ion may have also occurred in the process of g row ing larger

precipitates. These processes of recovery and precipitate d issolut ion wou ld be

expected to cause s t rength at the aging temperatures to decline.

79

1.2

10 11 12 13 14 15 P

8 R a t i o (R) o f u l t i m a t e s t r e n g t h o f a g e d t o u l t i m a t e s t r e n g t h o f u n a g e d m a t e r i a l

as a f u n c t i o n o f t i m e a n d t e m p e r a t u r e i n a p a r a m e t e r i z e d (P)

f o r m f o r t h r e e hea t s o f m o d i f i e d 9 C r - l M o

F i g . 9 Y i l d e s t r e n g t h as a f u n c t i o n o f t e m p e r a t u r e c o m p a r i n g u n a g e d

^ :12.418-PA1.0584

P: T ( l o g / 1 0 y l . 0 0 0 )

s . T: Temperature(k) • • • _ t ime(h)

• • •

• • ^ ^ ^ • • •

• ^ ^

• • •

• 755 k (482 C )

El 811 k ( 5 3 8 ° C )

® 866 k (.593 °C)

• 922 k (649 C ) A g i n g and Test Temperature

A 977 k (704。C) Were the Same

o o

o o o

o o

o o o

5 4

3

2

1

(BP-IAO

石 sgfcsscuf

00 丨00 00

T('C)

00

(^pagBun IPMUoJls

3sssn/p<u§v qauallsutsUISn

t o e s t i m a t e s o f aged m a t e r i a l

F i g . 10 U l t i m a t e s t r e n g t h as a f u n c t i o n o f t e m p e r a t u r e c o m p a r i n g u n a g e d

t o e s t i m a t e s o f aged m a t e r i a l

6. 2 Creep rupture properties

The creep behavior of base metal modif ied 9 C r - l M o is given for

temperatures ranging f rom 427 °C to 704 °C [ 1 ] .

Booker114] develops the equation to predict the stress-rupture properties of

material at di f ferent temperatures, wh ich is as follows：

log tr = Ch-0. 231<T— 2. 385log ^ + 3 1 . 0 8 0 / 7 (5) where

tr ： rupture l i fe ( h )

(7：stress ( M P a )

T： temperature ( K )

The parameter Ch is a " l o t constant" that reflects the relative strengths of

dif ferent lots of mater ia l , assuming that the stress and temperature dependence

is the same for al l lots. The average value of Ch was -23. 737. The analysis

yielded an overal l Standard Er ro r of Est imate ( S E E ) of 0. 324 and a m in imum

Ch of-24. 272.

Predictions of rupture behavior of base metal using above rupture equations

are very good over the temperature range of 427 °C to 704 °C.

Long term creep-rupture duct i l i ty is as an indication of the resistance of the

material to creep-fatigue interact ion and to such creep phenomena as stress

relaxation induced cracking. ORNL C l ] gives a plot (F ig . 11) of creep rupture

duct i l i ty for mul t ip le heats at several temperatures as a fuct ion of t ime. The plot

shows that modif ied 9 C r - l M o has excellent short term creep-rupture duc t i l i t y , 1 81

Average Unaged

M a x i m u m Stress A l l owab le ,

300

T('C

00

o

o

o

o

o

o

o

o

o

8 6

4

2

(cdds) 二tbca>i3s

31}SS1(D^UIPln

400

1 1 1 1 1 1

V •

_ A A

i 1 1 1 1 i ‘

7礙灣

1 ‘ 1 1 ‘

祖 A _

1 1 1 1 1 1

V •

_ A A 厶 •

i 1 1 1 1 i ‘

7礙灣 o •

. Test Temperatures (°C)

Heats F5349 30182 30176 30383

0 • % -

O — 4 8 2 0 -— A - 5 3 8 6 —

• — 593 V -

-649 -677 -704

.30394 10148 91887 X A 3 6 0 2 14361

I . . . • 1 •

• -• _

A V v Oj

•

10° 101 102 1 03 1 04 1 05

T ime to Rupture (h)

F i g . 11 C r e e p - r u p t u r e d u c t i l i t y d a t a as a f u n c t i o n o f r u p t u r e t i m e

a t v a r i o u s t e m p e r a t u r e s f o r c o m m e r c i a l hea ts

and decreased duct i l i t y beyond about 20 000 h，b u t no values below 10% in

terms of reduct ion of area.

Creep rupture s t rength behavior is also changed by pre-thermal. The

results [ 1 ] show that pre-aging does reduce creep rupture st rength somewhat and

that for 75 000 h exposure rupture strengths can fa l l below m in imum values for

unaged material. See Fig. 12.

Creep fracture occurs main ly at the softening zone in the heat affected zone

( H A Z ) and only rarely in the weld metal. A soft zone that forms at the outer

region of the H A Z adjacent to the base metal is called Type I V zone. I n this

V U

O

O

O

O

O

O

O

A

U

9

8

7

6

5

4

3

0

(。/0)B9JVJo§

^

j

i

r

o o

o

o

o

5 4

3

2

1

(0/。)UOSB§UOI3 -SOH

82

zone, carbides coarsening takes place and this minimizes the precip i tat ion

strengthening114 ] .

The reduct ion of creep s t rength of modi f ied 9 C r - I M o welded j o in t is due to

this creep weak zone[16"18]. The creep weak zone is very narrow. The l imi ted

w i d t h of the zone leads to macroscopical ly low duct i l i t y in the type I V fractures.

I t is inevitable that the Type I V zone w i l l always be present irrespective of

weld ing process. A higher heat input w i l l lead to a wider zone wh ich is more

detr imenta l to creep propert ies than a th inner one.

000

100

Average for Heats Tested i n Unaged Cond i t ion

M i n i m r n for Heats Tested Inunaged Cond i t ion

• Preaged 50 000 h B Preaged 75 000 h

T w o Heats Thermal A g i n g and Temperatures Were the Same and Var ied f r o m 482 to 649 °C

- 4 0 - 3 8 - 3 6 - 3 4 - 3 2 log ( 0 -31080 /7

- 3 0 - 2 8

F i g . 12 A c o m p a r i s o n o f t h e c reep e q u a t i o n p r e d i c t e d r u p t u r e s t r e n g t h s v a l u e s

a n d da ta o b t a i n e d f r o m t e s t s c o n d u c t e d o n p r e - a g e d m a t e r i a l

Fig. 13 indicates the creep rup ture s t rength as a funct ion of Larson Mi l le r

Parameter ( L . M . P. ) [ 1 9 ]：

P = T ( 3 0 + log t)/l 000

where t is the rupture t ime in hours spent at the test ing temperature T ( K ) .

I n the case of higher stress and short t e rm side where L . M . P. values are

smal l , creep rup ture occurred in either the we ld metal or the base meta l ,

depending on the s t rength balance between weld metal and base metal. As

L. M . P increased, e. g. greater than around 29. 5 X 10 3， the creep rupture

locat ion tended to sh i f t f r o m the base metal or we ld metal in to the softened H A Z

for the change in s t rength balance at the over-al l welded jo in t . H A Z softening 83

25 26 27 28 29 30 .31 32 33 34 35

p=r(iogH30)xi0"3

F i g . 13 C r e e p r u p t u r e p r o p e r t i e s o f M o d i f i e d 9 C r _ l M o s tee l w e l d e d j o i n t ( S A W )

a n d 1 2 C r s tee l w e l d e d j o i n t ( G T A W ) a f t e r P W H T

behavior is probably due to the lack of fine Nb and V precipitates coherent w i t h

the mat r ix in addit ion to the fine grained microstructure formed by recovery and

polygonalization of the tempered martensite.

The reference [20] uses the creep l i fe assessment method to study the creep

properties of modif ied 9 C r - l M o steel. I t is shown that the creep damage (t/tr)

(t is interrupted t ime, tr is creep rupture t ime) in the base metal corresponds to

the decrease in hardness. I n the base of the kinetic of the dislocation in the

creep，the hardness is correlated w i t h creep damage. So the hardness could be a

good measure of the creep damage in the base metal. I n the weld jo in ts , voids

are formed and grow in the fine grained H A Z region dur ing creep. The area

fract ion of void and void density are increased w i t h the creep damage. See Fig.

14—15.

The creep rupture occurs at this softest H A Z zone. Assuming the creep

strain is associated w i t h the void format ion and g rowth in the fine grained H A Z ,

[20] gives the relat ionship between void area fract ion and creep damage(Fig. 15).

• 550 °C

• • 593/600 °C

O O 650 °C

open : ruptured in H A Z

so l id : ruptured in w e l d metal or base metal

(Bl) ils

84

• 873 K,108 M P a

tr = 5 353 h

• 923 K ’ 4 9 MPa

" = 7 5 9 5 h

•

•

20 40 60 80 100

Creep Damage tltr (%)

R e l a t i o n b e t w e e n v o i d d e n s i t y a n d t/tr i n w e l d e d j o i n t

160 0 20 40 60 80 100

Creep Damage t/tr (%)

F i g . 14 V i c k e r ' s h a r d n e s s o f c r e p t base m e t a l

For modif ied 9 C r - l M o steel, i t is shown that both the base metal and

weldment have the best creep rupture strength.

The creep rupture strengths of al l weldments are similar to the average

creep rupture strength of the base metal. The equation ( 5 ) developed for base

metal is assumed to be appropriated for weldment w i t h a new lot constant Ch，

400 O aged,873 K 9 aged,92.3 K

• crept,873 K - 1 2 5 M P a • crept,923 K - 7 6 M P a

A crept,873 K - 1 5 7 M P a A crept,923 K - 9 8 M P a

• O ° n 〇 • 〇 〇 i • ©O 0 • • • • •

• 口 a

B • •

2 8

6

4

2

o

5

1

(％ )

POA JO U.210B43CS9J

V .lg.

F

]

i

I

o o

o

o

4 2

o

8

2 2

2

1

(A//)SS9UPJBKg ^

85

which is equal to-24. 257. Th is assumption seemed val id since al l fai lures

occurred at the edge of the H A Z in the base metal. The average creep rupture

st rength of the weldment roughly coincides w i t h the m in imum strength of the

base meta l , and is somewhat less than that of the average st rength of base

meta l , for the reason of a weakened or soft region at the edge of the H A Z ⑴ . A t

593 °C and for 30 000 h , the ratios of average weldment to average base metal

strengths is equal to 0. 84.

JWES [ 2 1 ] (Japan Weld ing Engineering Society) researches the creep, fat igue

and creep-fatigue propert ies of modif ied 9 C r - l M o steel including weldments

tested in air and found that the forg ing material of 550 m m thickness has almost

the same creep rupture st rength and duct i l i t y as that of the 25 m m thickness

plate, and creep strength of these materials and their weldments agreed

relat ively closely w i t h Booker 's rupture equation ( 5 ) . The results in [22 ] are

similar： the creep rupture s t rength of the tube is almost the same as that of the

plate and the creep rupture st rength of the weldment is almost the same as that

of the base metal at 500 and 550 °C，and s l ight ly lower than that of the base

metal at 600 °C. See Fig. 16.

6. 3 Creep fatigue properties

Long te rm thermal aging at temperatures results in a reduct ion in the

continuous cycl ing fat igue and creep-fatigue lives of modif ied 9 C r - I M o steels [23]，

as a result of the changes in the microst ructure dur ing aging. Dur ing aging,

N b C carbides disappear and the volume fract ion of V C carbides reduces, wh ich

carbides are main ly responsible for the al loy 's strength in the normalized and

tempered condit ion. Precipi tat ion of laves phase also occurs dur ing aging,

removing M o , wh ich is a potent solid solut ion strengtheners and retardant to

dislocation recovery/recrystal l izat ion.

The fatigue l i fe of the forg ing material is approximately the same as that of

the plate, and the fat igue l i fe of the weld metal or weldment is shorter than that

of the base metal The l i fe reduct ion in the weldment is caused by strain

concentrat ion in the base metal [ 2 1 ] . See Fig. 17-18.

The cyclic stress-strain relations for the base metal and weld metal are given

Fig. 18 [21 ]. The plate and forg ing mater ial show the same behavior, whereas the

base metal shows a softer behavior than the weld metal.

Local strain measurements reveal that the strain range in the H A Z become

86

100

70

50

10° 102 103 104

Rupture t ime t, (h)

105 106

F i g . 16 s t r ess vs . r u p t u r e t i m e o f base m e t a l a n d w e l d m e n t

smaller w i t h increasing cycles, the st ra in concentrat ion occurs in the base metal

rather than in the H A Z . I t is consistent w i t h the experimental results showing

that the fai lure in the weldments occurred in the base metal. See Fig. 19 [ 21 ] .

Examinat ions of the effect of strain hold posi t ion on creep-fatigue l ife

indicated that for modif ied 9Cr l M o the compression peak hold gave shorter

lives than the tensile-peak or intermediate hold tests. Fatigue l ife reduct ion w i t h

hold period for the modif ied 9 C r - I M o steel are generally smaller than that for

the conventional austenitic stainless steels.

6 .4 Toughness

The weld metal toughness may be thought to be irrelevant in assemblies

designed for operation in the temperature range of 500 〜600 °C，since this is

def ini tely far in excess of the temperature range at which br i t t le f racture could

occur. However , the components may very we l l also be stressed at ambient

temperature dur ing test ing or start-up. T o minimise the r isk of b r i t t le fracture

in these s i tuat ions, a m in imum toughness average for weld metal of 47 J at 20 °C

has been introduced in the European specification E N 1599 ： 1997 [ y ] .

M o d i f i e d 9 C r - l M o steel

\ Base metal

Weldment

Plate 〇 • Forging A •

550 'C

600。C

Base metal

Weldments

g

o

o o

o

g

o

o o

o

o 7

5

3

2

(Bp-2)

to sis

87

0 0.4 0.8 1.2 1.6 2 Total strain range A st (%)

F i g . 18 C y c l i c s t r e s s - s t r a i n c u r v e s f o r base a n d w e l d m e t a l s

As mentioned in section 4， the chemical composit ion of we ld metal has an

impor tant influence on the toughness [ 9 ] . N icke l has a great influence on the

toughness： the higher the N i content， the higher the toughness. Th is is caused

by the decrease of the A c l temperature towards the tempering temperature w i t h

102 103 104 105

Number o f cycies to fai lure N t (cycle)

106

F i g . 17 L o w - c y c l e f a t i g u e d a t a f o r t h e base m e t a l ,

w e l d m e n t a n d w e l d m e t a l

000

M o d i f i e d 9 C r - l M o steel, 550 °C

Weld metal

Base metal

\ Base metal

Weld metal

Plate O • Forg ing A —

o o

o

o

o o

o

o

8 6

4

2

(Bds)bv 3gc2

Ss9hs

(％)co<1

&PS ulBhs

PSOH

88

1..2

M o d i f i e d 9 C r - l M o Weldment,550 °C

. 10° 101 102 103 104

Number o f cycles N (cycles)

F i g . 19 E x a m p l e o f m e a s u r e d l o c a l s t r a i n b e h a v i o r o f a w e l d m e n t

the increasing N i content. A l t h o u g h the N iob ium content is very sma l l， i t is

very impor tan t for the creep propert ies，but has a large effect on the duc t i l i t y of

the we ld jo ints. The N b content should be adjusted at the lower l im i t at 0. 03%

~ 0 . 0 5 % ( 0 . 04% in section 4 )，w h i c h is an op t imum for deposited we ld metal.

Fabricators str ive to minimize the durat ion of the P W H T for economic

reasons. Th i s is done by adjust ing the P W H T temperature as h igh as possible

wh i le at the same t ime considering the A c l temperature. F u r t h e r m o r e，P W H T

temperature close to the A c l point improves the duct i l i ty . I t can be seen that the

toughness increases w i t h increased temperatures and hold ing t imes. For a

P W H T 760 °C，2 hours m i n i m u m should be used to assure suff ic ient toughness

(41 J at 20。C). On the other hand， i t is possible to achieve a h igh level of

toughness w i t h lower temperatures using longer hold ing t imes， fo r example，

730 °C for 6 h.

The weld ing posi t ion also has inf luence on toughness [ 9 ] . By experience the

best toughness propert ies are achieved if the weaving technique is used instead of

the usual str inger bead weld ing.

A l t h o u g h the weld ing process has no inf luence on the creep s t reng th , i t is

signif icant to impact toughness. The highest toughness is normal ly achieved

o.

(o/o)

3

如 us

u'a-qs lisOJL

89

using G T A W , whi le f l ux processes such as M M A and S A W produce lower

toughness values. I t is because of vary ing oxygen content. A n oxygen content of

less than 100〜200X10_ 6 may be reached w i t h G T A W procedures, wh i le M M A

and S A W w i l l produce typical oxygen contents in the range of 400 〜800 X

1 0 - 6 [ 7 ] .

7. Conclusion

Modi f ied 9 C r - I M o is the best candidate for higher service temperature. I t

has attract ive properties： h igh creep strength w i t h good duc t i l i t y , h igh

resistance to cracking, good we ldab i l i t y , h igh thermal conduct iv i ty and low

thermal expansion coefficient. As far as selecting op t imum preheat temperature

and suitable P W H T , cont ro l l ing the chemical composit ion of weld meta l , good

mater ial properties of modi f ied 9 C r - I M o weldment w i l l be obtained.

• The N b content in we ld metal is lower than in base meta l , is l imi ted to

0. 0 4 % ~ 0 . 0 8 %， a n d N i content in weld metal is higher in weld metal

than in base meta l , up to 1 %. On the other hand the M n + N i contents

in weld metal is not al lowed to exceed 1. 5 % .

• The preheat temperature is best to be selected w i t h mutua l consideration

of ul t imate tensile s t rength , specific elongation and marten site content.

• P W H T 760 °C for 2 hours m in imum should be used.

Note

Th is wo rk was performed by Dr. Xiaot ian L i dur ing her vis i t in C E A

(Cadarache) f rom A p r i l , 2002 to October , 2002.

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