Influence of precipitation on dislocation substructure and creep properties of P91 steel weld joints

Influence of precipitation on dislocation

substructure and creep properties of P91 steel

weld joints

Dagmar Jandova* and Josef Kasl

SKODA VYZKUM s.r.o., Tylova 1y57, 316 00, Plzen, Czech Republic

*E-mail: [email protected]

ABSTRACT

Two trial weld joints were prepared using the GTAW and SMAW methods and they underwent creep

testing at temperatures between 525 and 625�C. The longest time to rupture was 45,811 h. Two main

processes occurred during creep exposures: recovery and precipitation of secondary phases. Slight

coarsenings of the M23C6 carbide, precipitation of Laves phase and Z-phase were observed after long

tests at high temperatures. Some differences in microstructure and creep failure were found in the

individual zones of weldments. After long exposure at temperatures up to 600�C, fractures occurred

in the fine-grain heat-affected zone as a result of a low density of fine vanadium nitride and a high

density of coarse particles at grain and subgrain boundaries. At 625�C, growth of Laves phase caused

a softening of the ferritic matrix and crack propagation in the weld metal.

Keywords: creep resistant steels, weld joints, microstructure, precipitation

1. INTRODUCTION

Grade P91 is the main representative of creep resistant

modified 9Cr – 1Mo steels. It is currently used for manufac-

turing boiler and turbine components of fossil fuel power

plants, especially steam piping and rotors, but also turbine

casings, outlets and so on. As these components have to

operate at severe conditions for years, high long-term

structural stability is required.

Modified 9Cr – 1Mo steels are generally used in conditions

after quenching and tempering. Two main processes take

place during tempering. First, recovery causes a reduction in

the high dislocation density and the formation of sub-grains.

Second, precipitation of carbides, nitrides or carbonitrides

occurs. The M3C carbides are dissolved and the more stable

phases arise. Chromium rich M23C6 carbides precipitate at

prior austenite grain boundaries, ferrite lath and subgrain

boundaries, while vanadium and niobium rich MX nitrides or

carbonitrides are spread mainly within subgrains [1,2].

The high creep rupture strength of P91 steel relies on the

martensitic transformation hardening and additionally: (i)

substructural strengthening of coarse M23C6 carbides,

which pin the grain and subgrain boundaries; (ii) precipita-

tion strengthening of fine MX particles; and (iii) strength-

ening of the solid solution by molybdenum atoms. Changes

in size and distribution of secondary phases taking place at

creep conditions cause a decrease in the creep rupture

strength and are strongly dependent on the temperature and

duration of the creep exposure. Both M23C6 carbides and

MX nitrides are considered to be stable during long-term

creep exposures at temperatures up to 600�C [3]. Significant

changes in substructure can occur during loading at 550�Cand higher temperatures as a result of precipitation of two

undesirable phases: (i) the Fe2Mo Laves phase that causes

depletion of ferritic matrix by molybdenum; and (ii) the

modified Z-phase that causes dissolution of MX particles [4].

These microstructural processes taking place in the creep-

exposed weld joints can induce an unexpected decrease in

creep strength. The kinetics of precipitation depends on the

local chemical composition and the initial structure of the

individual zones of the weld joint. Thus, the final structures

and also the creep properties of the weld metal (WM), heat

affected zones (HAZ) and the base material unaffected by

welding (BM) can be different. Consequently, creep failure

is concentrated in specific regions of the weld joints and

different types of cracking occur [5,6].

The study investigates trial weld joints which underwent

long-term creep testing. As received and crept specimens

were compared and microstructural changes in WM, HAZ

and BM were evaluated. The goal of the study was to

elucidate the microstructural processes taking place in

different zones of the weld joints which result in specific

types of fractures. Partial results have already been presented

at several international conferences [7 – 10].

2. EXPERIMENTAL MATERIAL AND

PROCEDURES

Two trial weld joints were fabricated from wrought or cast

P91 steel using GTAW and SMAW methods. The weld

assigned as ‘C’ was produced by joining cast plates

5006150625 mm in size, and the C1 weld was made by

MATERIALS AT HIGH TEMPERATURES 27(2) 1–000 1# 2010 Science Reviews 2000 Ltddoi: 10.3184/096034010X12710730545552

MHT090297

FIRST PROOF

joining pipe segments with an outer diameter of 325 mm, a

wall thickness of 25 mm and a length of 400 mm. The

chemical composition of the base materials and the weld

metals is shown in Table 1. Both base materials were

austenitized at 1050�C for 1.5 h, and then were oil quenched

and tempered at 750�C for 3.5 h. The post-weld heat

treatment (PWHT) was applied as follows: (740 –

750)�Cy2.5 h for the C weld joint and 760�Cy2.5 h for the

C1 weld joint. Smooth cross-weld specimens underwent

creep testing to rupture in a temperature range from 525 to

625�C at stresses from 50 to 240 MPa. The longest time to

rupture was 45,811 h. Creep data were evaluated in relation

to the Larson – Miller parameter P ¼ T � ðCþ log tÞ, where

T is temperature (K), C is a material constant equal to 25 and

t is time to rupture (h). The creep strength of both weld joints

fell into the usually permitted + 20% scatter band of creep

strength for the corresponding base material in a temperature

range up to 575�C. At higher temperatures, it decreases

below the bottom of that scatter band. Applied stresses

versus Larson – Miller parameter are graphically represented

in Figure 1. In addition, the type of fracture is indicated.

It was found that the type of fracture depends on the creep

test conditions and on the kind of weldment. The pipe weld

joint ruptured in the weld metal or in the fine-grain heat-

affected zone (FG HAZ) if the Larson– Miller parameter

were lower or higher than 24,000, respectively. The plate

weld joint fractured in the FG HAZ across almost the whole

range of temperatures and applied stresses with one excep-

tion; the fracture was located in the weld metal after testing

at the highest temperature and the lowest stress, which

corresponded to the highest Larson – Miller parameter. All

the fracture surfaces were observed using scanning electron

microscopy. No defects were detected which would cause

premature rupture of weld joints. This means that fractures

occurred as a result of gradual microstructural changes

taking place during creep exposures. Different strain rates

and development of cavitation failure in individual zones of

the weldment resulted in crack propagation in specific

regions.

Microstructural investigation was carried out using light

metallography, scanning (SEM) and transmission electron

microscopy (TEM). Secondary phases were identified using

energy dispersive X-ray microanalysis (EDX) and selected

area electron diffraction (SED). Carbon extraction replicas

were prepared from the weld metals and the base materials.

Thin foils were prepared from the weld metal, the coarse

prior austenitic grain heat affected zone (CG HAZ), the fine

prior austenitic grain heat affected zone (FG HAZ) and the

base material. Foils were jet electropolished to transparency

in a 6% solution of perchloric acid in methanol at tempera-

tures ranging from � 60 to � 40�C.

Cross-weld hardness profiles were measured before and

after the creep tests.

3. RESULTS AND DISCUSSION

The microstructure in all the weld joint zones was tempered

martensite formed with a ferritic matrix and particles of

secondary phases. Differences in individual zones could be

observed on etched metallographic samples using light

microscopy and SEM. After PWHT a lath-like structure

was evident in the WM, the CG HAZ and in the BM,

while a featureless structure occurred in the FG HAZ.

Some evidence of dendritic segregation was observed in

the cast base material and small globular-shaped inclusions

were present in the weld metal. Coarse chromium rich M23C6

carbides were located at prior austenite grain boundaries,

boundaries of ferritic laths and also at subgrain boundaries.

The density of coarse carbides was relatively low in the WM

and the CG HAZ and markedly higher in the FG HAZ.

The dislocation substructure and fine precipitates were

studied using TEM. Dislocation walls were present in some

Influence of precipitation on P91 steel weld joints: Dagmar Jandova and Josef Kasl

2 MATERIALS AT HIGH TEMPERATURES www.materialshightemp.co.uk

Table 1 Chemical composition of weld metals and base materials given in wt%

C Mn Si Cr Mo V Ni Nb Al N P S

BM pipe 0.12 0.49 0.29 8.6 0.96 0.21 0.30 0.07 0.011 0.06 0.011 0.004BM plate 0.12 0.41 0.21 8.8 0.92 0.22 0.10 0.09 0.004 0.06 0.014 0.008WM pipe 0.12 0.70 0.21 9.1 1.05 0.21 0.73 0.05 0.003 0.04 0.010 0.006WM plate 0.12 0.71 0.23 9.1 1.05 0.20 0.68 0.05 0.005 0.04 0.012 0.009

Figure 1 The creep rupture strength of trial weld joints: (a) the C plate weld joint and (b) the C1 pipe weld joint in comparison to the creep

strength of the base materials (wrought P91 steel– X10CrMoVNb 9 1 W and cast P91 steel – GX12 CrMoVNbN 9 1).

(a) (b)

laths that formed subgrain boundaries. The density of free

dislocations within the subgrains was lower than the density

of the non-arranged dislocation nets in other laths. The

highest dislocation density was observed near the fusion

line. A relatively high dislocation density was also observed

in the weld metal and in the base material. A significantly

lower dislocation density was found in the FG HAZ. The

occurrence of fine intragranular V(C,N) precipitate was

proved using electron diffraction in all zones of the weld-

ments. Some small orthorhombic (Cr,V)2C carbides were

also detected in the weld metal of the plate weld joint.

The precipitation reaction taking place during creep

exposures was strongly dependent on the creep loading

conditions. After long creep tests, no M2C carbide was

identified in the weld metal of the plate weld joint. This

non-stable phase dissolved. No important changes in the size

of coarse M23C6 carbides and fine vanadium nitrides were

observed after creep tests (Figure 2). A slight increase in

hardness after long creep tests at 600�C (Figure 3) was

probably caused by the precipitation of new particles of

both these phases, which have a positive influence on the

creep strength of P91 steel.

During long creep tests at temperatures ranging from 550

to 625�C for about 30,000 h, Fe2Mo Laves phase precipi-

tates. It often nucleates at the surface of M23C6 carbides or

small oxides in the weld metal, and forms large irregular-

shaped particles and clusters with other particles. In such a

form, the Laves phase cannot contribute to precipitation and

substructural strengthening. On the contrary, they can serve

as nucleation centres for cavities and cracks. In addition, the

growth of Fe2Mo particles causes depletion of the solid

solution by the molybdenum and softening of the ferritic

matrix. Laves phase was observed in all zones of the plate

weld joint, however, the largest particles (up to 1 mm in size)

are present only in the weld metal (Figure 4). This fact is the

main reason for the rupturing of the plate weld joint creep

tested at 625�C for 29,312 h in the weld metal.

A great effort has been made to find the Z-phase. Some

plate-like particles with dimensions of aproximately 100 nm

were detected in the WM, the HAZ and the base materials of

both weld joints after creep tests at 625�C for about 30,000

hours. The modified Z-phase corresponded to Cr(V,Nb)N

nitride, which can crystallize in two modifications with a

tetragonal or a cubic lattice [11]. Diffraction patterns of both

modifications were observed in the weld joints. The EDX

spectrum and diffraction patterns of M23C6 carbide and Z-

phase are shown in Figure 5. The Z-phase is considered to be

very harmful for creep resistance of P91 steel as it causes

dissolution of vanadium nitrides. Research has been

published that shows one Z-phase particle can destroy


www.materialshightemp.co.uk MATERIALS AT HIGH TEMPERATURES 3

Figure 2 The base material of the C plate weld joint after creep test

at 625�Cy50MPay29,312 h. TEM micrograph.

Figure 3 Cross-weld hardness profiles before creep tests and after long creep exposures at 600 and 625�C: (a) the C plate weld joint and (b)

the C1 pipe weld joint.

(a)(b)

about 1,000 fine MX precipitates [12]. Although only a few

Z-phase particles were found in this experiment, it can be

expected that they would cause a drop in creep strength after

exposures at temperatures above 625�C significantly longer

than 30,000 h.

No important changes in size and distribution of

secondary phases were observed in the pipe weld joint

after creep test at 625�Cy60 MPay13,673 h, nevertheless a

growth of subgrains was evident in the FG HAZ. After creep

test at 625�Cy50 MPay29,962 h subgrains in the FG HAZ

were significantly coarser than those in the base material and

the weld metal. The pipe weld joint fractured in FG HAZ as a

result of extensive cavitation failure.

In the plate, weld joint that underwent rather longer creep

testing, growth of subgrains in the FG HAZ was also

observed. However, after the longest creep exposure at

625�C, growth of subgrains in the weld metal was faster

and the weld metal became the softest zone in the weld joint

(Figure 6).

4. CONCLUSIONS

Precipitation processes taking place in trial weld joints

during long-term creep exposures caused creep strength

decrease in specific regions: in the FG HAZ or in the WM.

Coarse chromium rich M23C6 carbide, fine (V,Nb)N preci-

pitates and atoms of molybdenum dissolved in the ferritic

matrix were responsible for strengthening the tempered

martensite.

Ferritic grains, less lath-like and more equiaxed, with

many coarse carbides at grain and subgrain boundaries,

relatively low density of fine intragranular precipitate and



Figure 4 The weld metal of the C plate weld joint after creep test at 550�Cy50 MPay29,312 h: (a) the substructure with precipitates, TEM

micrograph; (b) diffraction patterns of Fe2Mo particle; (c) schema of diffraction patterns – spots of two ferrite grains and the zone z ¼ ½1��21��3�

of Fe2Mo Laves phase; and (d) EDX spectrum of Laves phase.

(a)

(b)

(c)

(d)

low dislocation density, occurred in the FG HAZ. The

hardness of such structures was lower than that of the lath-

like structures in other parts of the weld joints. These soft

zones are generally susceptible to fracture during long-term

creep exposures. Plastic deformation of individual grains was

caused by dislocation creep and resulted in the formation of

steps at grain and subgrain boundaries that could serve as

nucleation centres of cavities. Cavities were also formed as a

result of interaction between dislocations and coarse carbides

at grain and subgrain boundaries. Subsequent growth and

coalescence of cavities was promoted by grainysubgrain

boundary diffusion. Cavities were concentrated in the HAZ

after creep tests at temperatures up to 600�C. Individual

cavities also appeared in the weld metal and the base material

after exposures at 625�C and in the WM after exposures at

550�C for more than 30,000 h.

During long creep exposures at temperatures above 550�Cprecipitation of Laves phase occurred, which caused a drop

in the strengthening of the ferritic matrix by molybdenum.

Laves phase was present especially in the weld metal and

often formed clusters with M23C6 carbides and silicon

oxides. Consequently, fracture occurred in the weld metal,

although extended cavitation was also present in the FG

HAZ.

At long exposures at temperatures above 600�C, two

competitive processes of creep failure are expected in

similar P91 steel weld joints: cavitation failure in the FG

HAZ and softening of the ferritic matrix as a result of the

growth of Laves phase particles. In addition, under creep

loading at temperatures above 600�C for more than

30,000 h, creep strength decreases as a result of the

precipitation of Z-phase.


www.materialshightemp.co.uk MATERIALS AT HIGH TEMPERATURES 5

Figure 5 The weld metal of the C plate weld joint after creep test at 625�Cy50 MPay29,312 h: (a) substructure with precipitates, TEM

micrograph; (b) diffraction patterns of tetragonal Z-phase and ferrite matrix; (c) diffraction patterns of the zone z ¼ ½012� of M23C6 carbide; (d)

EDX spectrum of Z-phase; and (e) M23C6 carbide.

(a)

(b)

(c)

(d) (e)

ACKNOWLEDGEMENTS

This work was supported by grant project SMS 4771868401

from the Ministry of Education, Youth and Sports of the

Czech Republic.

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Figure 6 Substructure of the C plate weld joint after creep test at 625�Cy50 MPay29,312 h: (a) the base material and (b) the weld metal. TEM

micrographs.

(a) (b)

Influence of precipitation on dislocation substructure and creep properties of P91 steel weld joints

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