Simulation of Dissimilar Weld Joints of Steel P91

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Simulation of Dissimilar Weld Joints of Steel P91

Jiří Sopoušek
1
, Rudolf Foret
2
, Vít Jan
2

1
Masaryk University


Brno, Faculty of Science, Department of Theoretical and Physical Chemistry,
Kotlářská 2, CZ
-
61137 Brno, Czech Rep.

2
Dept. of Materials Scienc
e, Faculty of Mechanical Engineering, BUT, Technická 2, Brno CZ
-
61669 Brno, Czech Rep.

E
-
mail addresses:
foret@umi.fme.vutbr.cz
,
sopousek@chemi.muni.cz

(correspon
dence author),
jan@pime.fme.vutbr.cz

.


Abstract

Results of computer simulations of long
-
term service exposure for weldments of the CSN

15

128

/

P91 and
SK3STC

/

P91 steels are presented and compared with cor
responding results of phase and composition
experiments. The welded material P91 (EU standard: X10CrMoVNb 9
-
1) represents progressive chromium
steel alloyed with Mo, V, C, and N. The CSN

15

128 (EU standard: X10CrMoVNb 9
-
1) material is low
-
alloy Cr
-
Mo
-
V st
eel. The SK3STC alloy (EU standard: 12CrMo 10
-
10) represents the consumable electrode material.
The stability of the weldment microstructure is investigated at elevated temperatures (500
-
700ºC). The
simulation method is based on the CALPHAD approach
comple
mented with the theory of

multi
-
component
bulk diffusion, local condition of phase equilibria and the
assumption that diffusion is the process that
controls the rate of phase transformation (the DICTRA software is used). S
ignificant phase profiles,
concent
ration profiles, and phase transformation processes in diffusion
-
affected zone are simulated,
investigated, and compared with experimental results. The potential risky carbon depleted region inside each
weld joint is discussed. The method described can be
used to predict microstructure instability in weld joints.


Keywords: DICTRA, diffusion, profile, redistribution, consumable electrode, weld.


2


3

Introduction

A weld joint of dissimilar materials under external stress usually represents a critical point in m
any technical
solutions involving elevated temperatures. The investigation of the relationships among the element/phase
redistributions, the microstructure at various points across the weldments, and local mechanical properties
represents a method that is
suitable for the evaluation of long
-
term mechanical/microstructure stability of
weld joints
[98Pil]
. In the case of weld joint applications at elevated temperatures, the mechanical properties
can be related to chemical concentration and phase transformatio
n processes in the diffusion
-
affected zone
[89Buch], [90Wit] [95Mil]. Information on the time evolution of both the phase and the element
redistributions at a given temperature treatment is therefore very important.

The main factors that influence the sta
bility or instability of the weld joint of steels are, above all, carbide
nucleation, phase transformation, rate of diffusion [00Rin], and carbon depletion [98For]. These factors are
significantly dependent on temperature
.
Phase transformations in weld joi
nts of multi
-
component alloys have
become very complex due to the high degree of freedom of the multi
-
component system [90And], [86Por].

The diffusion fundamentally affects the rate of phase transformations at elevated temperatures [85Kir] and
the changes
in chemical potentials of the elements are consequently the cause of phase precipitations,
growth, phase dissolutions, and/or phase boundary replacement in the weld diffusion zone. In many cases,
regions parallel to the initial weld interface are formed in

the diffusion
-
affected zone, with different
microstructures having different mechanical and corrosion
-
resistant properties.

The theoretical and experimental methods for multi
-
component alloy weldment investigations were limited in
the past. The main theor
etical problem was the absence of a more complex model for multi
-
component welded
alloy system with dispersed phases that can describe the processes occurring in the weld joint with
satisfactory accuracy. If we accept that a diffusion couple [98Sau] may ap
proximate a weld joint, it is then
possible to perform a simulation of weldment temperature exposure using a model based on the CALPHAD
approach [98Sau] and complemented with the assumption that diffusion is the controlling process of the
phase transformat
ion rate [85Kir],
[
94Eng], [00Bor], theory of multi
-
component bulk diffusion [90And],
thermodynamic evaluation of the driving force for phase transformations, and the assumption of local
condition of phase equilibria. This complex model is also implemented

in the DICTRA software package
[98DIC] and it enables a prediction of the sequence of phase regions across the investigated weldment,
which are formed during the temperature exposure. The different creep strain rates are generally in different
regions of
the weldment under external stress and this leads to the generation of increased mismatch stresses
and earlier failure.

The aims of this study are: to present the results of weldment simulations for two weldments (CSN

15

128/P91
and SK3STC/P91), to compar
e the theoretical results with the foregoing experiments
[01For] [99Mil] [01Svob]
,
and to suggest the relation between the simulated phase region sequence, weldment microstructure and
mechanical properties of the weld joint. It is hoped that this approach
will be accepted as a progressive tool
for weld design.

Subject of investigation

The following materials and their weld joints will be in the focus of our study (see
Table 1

for detailed
compositions): P91 steel alloyed with chromium and small amounts of M
o, V, and N, 15 128 low
-
alloy Cr
-
Mo
-
V steel (according to Czech standard CSN

41

5128), and the consumable electrode material (Chromocord
SK3STC). The progressive P91 steel represents creep
-
resistant chromium steel for industrial application
[01Wil],

[99Orl
], [98Cad]. The 15 128 low
-
alloy steel mentioned above is a frequently used steel because it
gives a good combination of mechanical properties, resistance to corrosion, and cost.

The foregoing long
-
term (up to 10

000

hrs) experiments
[01For],
[01Svo], [
99
Mil] in the
temperature range from

500 to 700ºC

confirmed that the matrix of each
investigated diffusion couple preserved the BCC
_A2 (referred
to as α in the following)
structure
with time
-

and distance
-
dependent arrangement of carbides and/or
carbonitrides. The development of the following dispersed phases was observed in the weldments:
chromium
-

rich carbides M
7
C
3

and M
23
C
6
, speci
al chromium
-
poor carbide (M
23
C
6
), molybdenum
-

rich carbide
M
6
C, molybdenum
-

and vanadium
-
rich carbide (M
2
C), vanadium
-
rich carbonitride (MX, where X stands for C
or N).


4

Simulation model

The following simplifying fundamental assumptions were used in our mo
del. The metastable phases or the
stable phases with non
-
equilibrium compositions formed before post
-
weld heat exposure were excluded from
the simulation. The width of the fusion zone was put equal to zero because the experimental fusion zone was
also smal
l (about 3∙10
-
6

m). No heat
-
affected zone generated during the weldment preparation was supposed.
The initial state of the weldment was approximated by two systems, each in equilibrium.

In our simulations, the CALPHAD approach
[
98Sau] was used for the sol
ution of both local and global phase
equilibrium problems in the steels investigated. Every investigated steel represents a multi
-
element system
with several phases: the matrix phase including dispersed phases (carbides and/or carbonitrides). The
CALPHAD a
pproach enabled a solution based on constrained minimization of the total Gibbs energy in a
closed system at a given composition, temperature and pressure. In short, the total Gibbs energy is calculated

as a sum of molar phase Gibbs energies weighted by ph
ase ratios.
Using an appropriate thermodynamic
model
[81Sun],

[48Red],

[70Hil], [78Hil], [75Ind], [81Ind]

the molar phase Gibbs energy can be evaluated as a
function of temperature, pressure, phase composition and so
-
called thermodynamic parameters (TP). T
he
parameters for selected phases and selected systems that represent the subsystems of the steels under
examination can be found in the relevant literature. Numerically, this problem can be solved using available
software packages (
[
97TC],
[
89Pan],
[
93Sop
], [E
-
FACT], [E
-
malt], etc
.).

The following elements were considered for thermodynamic description of the materials examined: Fe, Cr, Mo,
V, C, N (cf. steel composition in
Table 1
). The Gibbs energy of the phases existing in this 6
-
element system
was desc
ribed using thermodynamic parameters from the STEEL12.TDB database [01Kro]. This database
contains data for the following subsystems: Fe
-
Cr
-
C, Fe
-
Mo
-
C, Fe
-
V
-
C, Fe
-
N
-
C, Fe
-
Cr
-
N, Fe
-
C
-
N, Cr
-
V
-
C, Cr
-
Mo
-
N, Mo
-
V
-
C, Cr
-
Fe
-
V and some data for higher subsystems: F
e
-
Cr
-
Mo
-
C, Fe
-
Cr
-
V
-
C, Cr
-
V
-
Mo
-
C, Cr
-
Fe
-
Mo
-
V
-
C. The phases were described using a regular solution model for phases with several components and
sublattices [81Sun]: b.c.c phase as the magnetic ALFA??? BCC_A2 phase with two sublattices:
(Fe,Cr,Mo,V)
1
(C,N,Va)
3
, where Va stands for vacancies, M
23
C
6

as the stoichiometric phase with 3 sublattices
(Cr,Fe,V)
20
(Cr,Fe,Mo,V)
3
C
6

, M
7
C
3

carbide as the stoichiometric phase with two sublattices (Cr,Fe,Mo,V)
7
C
3
,
M
6
C using 4 sublattice description Fe
2
Mo
2
(Cr,Fe,Mo,V)
2
C, vana
dium reach carbide M
2
C and vanadium
-
rich
carbonitride MX were modelled as a non
-

stoichiometric phase (Fe,Cr,Mo,V)
1
(C,N,Va)
1

using the parameters
for the FCC_A1 phase [92Lee]. This description was used to evaluate all the equilibrium quantities or
function
s presented (equilibrium composition, chemical potential, activity, etc).

In the case of heterogeneous weldment, a gradient of the chemical potential can be observed and the
diffusion [85Kir] of the species is going on at elevated temperatures. The changes

in the chemical potentials
of the species form the thermodynamic ground for phase transformations, phase precipitation, growth and/or
phase dissolution. In our paper the DICTRA [98DIC] program (including the ThermoCalc routines [97TC])
was used for the si
mulation of the post
-
weld heat exposure. This model deals with element mobilities
[92And]
for diffusion matrix evaluation.

The same selection of the species as for the thermodynamic description was considered for the kinetic
description.
The major intersti
tially diffusing species are carbon and nitrogen but the diffusion of the other
substitutional species (Fe, Cr, Mo, and V) was also taken into account. The weldments were simulated as
one
-
dimensional diffusion couples. S
teel kinetics was approximated using

the kinetic parameters for carbon
and nitrogen mobility evaluations in the ALFA???? BCC_A2 phase based on Fe
-
Cr
-
C
and

Fe
-
Cr
-
N
[94Jo1],
[02Jan] subsystems. The trace [82Ku
č
] [82Oik]
diffusions of substitutional elements in iron were included
too. The carbi
de and carbonitride phases were treated as non
-
diffusion spheroid phases suspended in the
matrix. A local equilibrium between the matrix and the carbide/carbonitride phases was assumed in each part
of the diffusion couple.


Simulation results

The simulati
on method used enabled us to predict any of kind of phase diagram cross
-
section of all the
investigated alloys. The presented phase diagram cross
-
sections (P91:
Fig. 1,
15

128:
Fig. 2
) show
equilibrated phases. The activities of the elements were calculate
d for the investigated alloys too. The carbon
activity in the investigated steels is given in
Fig. 3
.

The simulation using the DICTRA program offers the time
-

and distance
-
dependence (profiles) for many
values (molar phase ratio, overall concentration, bot
h matrix and carbide element concentrations, activity,
chemical potential, etc.). With respect to the experiment, the phase profiles and overall element concentration

5

profiles are the most interesting. Sometimes, the element composition profile of the indi
vidual phase can be in
the focus of attention too.

A simultaneous plot of the phase profiles is best
-
suited to distinguish the phase regions across the
weldment, the loading of the phases in the regions, as well as the width of the regions. The simulation
results
for two selected weldments are given in
Fig. 4
and

Fig. 5.
Generally, the sequence of the phase regions
observed in the couple is time independent, the rate of the region interface motion increases non
-
linearly with
the distance from initial weld i
nterface, and the region width increases approximately with the square root of
the time. The phase region sequences for simulated weldment combinations are summarized in
Table

2.

These
sequences enabled us to illustrate carbide reactions running in the wel
dments at given temperatures.

The simulation enabled the calculation of the time
-

and distance
-
dependence for the composition of the
dispersed phases too. The carbonitride composition profile for the 15128

/

P91 weldment after post
-
weld heat
treatment is
given in
Fig. 6

and the composition profile for the M
23
C
6

carbide is in
Fig. 7
.

The overall concentration profiles can also be obtained from the simulation and they can be easily compared
with experimental measurement
[01For],
[01Svo], [
99Mil]
. Two experi
mental and simulated overall carbon
concentration profiles are given in
Fig. 8,

where numerous experimental observations
[01For] are supplied by
fitted experimental curves.

Discussion

The CALPHAD method enables us to calculate the phase diagram of steels
with high reliability and without
much experimental work. The phase diagram predictions in
Figs 1

and
2

obtained using the STEEL12
thermodynamic database [01Kou] can be performed for technical steels too. The phase diagram predictions
may differ if we use
a different thermodynamic database (for example SSOL [97TC]) especially in low
temperature ranges. However, we believe that our thermodynamic database yields more accurate results
because the STEEL12 database is optimized for lower temperatures. The same e
mphasis is placed on the
consistency of the DIF12 kinetic database [02Sop] at low temperatures.

The high
-
temperature microstructure stability or instability of dissimilar weldments may be considered in
different ways. In the first approximation, the elemen
t activity in steels can be used as a determinative value
for weld joint stability. The carbon and/or nitrogen is of utmost importance (see
Fig. 3
). From this simple
view, the SK3STC steel represents a convenient electrode material for the fabrication of

t
he
15

128

/

SK3STC

/

P91 complex weld joint because the curve of carbon activity of the consumable lies
between the curves for joined materials. However, this approach is not sufficient for a long
-
term
stability/instability judgement.

In the second approx
imation, the simulated phase and element profiles can be used for an estimation of the
life time of weld joints and consequently of the time remaining for all technological facilities. The phase
region sequence (see
Table 2
) and phase profile (see
Figs 4

a
nd
5
) can help us to find the weakest point
inside the weldment because the mechanical properties of the region can be estimated from the phase
microstructure [01Raj]. This approach assumes that the weldment simulation results are correct.

The simulation
result corresponds with experimental observations [01For]
,
[01Svo], [
99Mil]

and any possible
discrepancies observed can be easily explained by the limitations of experimental technique, pre
-
weld heat
treatment, or simulation model limitations. The carbide
types found experimentally in steels weldments
correspond well with the diffusion couple simulation (see
Figs. 4

and
5
). All this confirms the stability of MX
carbides during post
-
weld temperature exposure across the whole
15

128

/

P91
weldment (
Fig. 4
), M
23
C
6

carbide nucleation and coarsening in the P91steel, carbon
-
depleted zone close to the initial weld interface
(see
Fig. 8
), etc. The simulation also shows the same predominant M
6
C carbide that formed a narrow band of
coarse carbides on the experimental

15

128

/

P91
weld interface
[01For]

(
Fig. 4
,
Table 2
). Further, t
he
simulations confirmed the existence of the phase regions: α+MX (15

128

/

P91) and α+M
6
C (SK3STC

/

P91),
which were found experimentally inside the carbon
-
depleted zone (see the temperatur
e range including 625˚C

in

Table 2
).

The experimental and simulated carbide/carbonitride compositions are in a good agreement too. For example,
the simulation confirmed the variation in the M
23
C
6

carbide composition in the experimental

15

128/P91
weldment
[01For] (see
Fig. 6
), the simulated chemical composition of MX carbonitride (
Fig. 7
) reflects the
experimental metal content, t
he simulations predict carbon
-
rich and nitrogen
-
rich carbonitrides, etc.

Examples of the experimental and the simulated bulk car
bon profiles are in
Fig. 8
. The extreme experimental
values (peak maximum/minimum) agree in the limit of accuracy of the experimental measurement with the
simulated values. In
Fig. 8
,

we can see the lower simulated maximum of carbon concentration in the
15

128

/

P91 weldment carburised zone. It can be easily explained by grain boundary diffusion, which is not
included in the simulation model.


6

It can be said that both weldment combinations (
15

128/P91, SK3STC/P91) represent quite solid joints under
the moni
tored conditions. However, there is a risk of decarburising the BCC_A2 zone formed during long
-
term exposure. This risk is reduced by carbide and/or carbonitride particles. It was found by simulation and
also by experiment that the decarburised 15

128/P91

zone contains the MX carbonitride and the decarburised
SK3STC/P91 zone contains the M6C carbide. The M
23
C
6
carbide predominates in the carburised zones of both

weldments. The 15

128/P91 (see
Table 2
) simulated phase region sequence includes a central (

+M
X) |
(

+M6+MX) subsequence with a high amount of M
6
C carbide. It reflects the experiment [01For]. These M
6
C
carbides formed a narrow band of huge particles.

Other phases (LAVES, M
2
C, V
4
C
3

, PI ) were also considered in weldment simulation. One of them, th
e LAVES
phase, represents a great risk. This phase has not been mentioned because it is not stable under the
simulated conditions but a small concentration shift or addition of another LAVES phase stabilising element
(namely Mn) may be the cause for the LA
VES phase precipitating in the P91 steel or in the decarburised zone.
The simulation confirms that the LAVES phase occurs at
temperatures below 600˚C and at lower carbon
concentrations (approx.
<0.01wt %).

Conclusion

The above approach and the simulation results enabled us to better understand the microstructure processes
and carbide and/or carbonitride phase transformations o
ccurring in the diffusion zone of the investigated
weldments. It can be generally said that the method applied yields results that reflect the weldment phase
structure quite correctly.

A good agreement was found in the sequence of the carbide and/or carbo
nitride phase regions.

The performed kinetic simulations provided information that can be used for failure
-
risk prediction for weld
joints. The simulated phase profiles for the vanadium
-
rich MX carbonitride and the fine
-
size M
6
C carbide
showed their high
microstructure
-
stabilising effect IN nebo ON ??? the carbon
-
depleted zone. Special
attention must be paid to the concentration balance, which may be responsible for the LAVES phase
precipitation in weldments that include the P91 steel.

The simulation resul
ts save time and expenses in the evaluation of long
-
term microstructure stability of the
weldments examined. The profiles presented can serve for the prediction of microstructure development and
estimation of the mechanical stability of the weldments under

external stress during long
-
term service.

Acknowledgements

The support of the Grant Agency of the Czech Republic (No: 106/00/0855) is gratefully acknowledged for
funding. The calculations were performed with the aid of the ThermoCalc and DICTRA programs.

Sincere
thanks are extended to P
rof. Vřešťál for thermodynamics and kinetics calculations.

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Engineering, Cambridge Int. Science Publ., Cambridge (1998)

[98Sau]

Saunders, N.; Miodovnik, A. P.: CALPHAD (Calculation of Phase Diagram)
-

A Comprehensive
Guide (Pergamon Materials Series, Vol. 1), Elsevier Sc
ience, Amsterodam (1998).

[99Mil]

B. Million, R. Foret, A. Rek, et al.: Carbon redistribution during long
-
term operation of weld
joints of creep
-
resisting chrome steels,KOVOVE MATER 37: (5) 314
-
323 1999. P91/15128
P91/Cr3Mo1C

[99Orl]

Orlova A, Bursik J, K
ucharova K, et al. Microstructural development during high temperature
creep of 9% Cr steel MAT SCI ENG A
-
STRUCT 245: (1) 39
-
48 APR 30 1998

[E
-
FACT] Web Sites in Inorganic Chemical Thermodynamics, FACT:


http://www.crct.polymtl.ca/fact/websites.htm


[E
-
Malt]

Thermodynamic database : MALT2:


http://www.kagaku.com/malt/emalt2.html



8



Table 1:


Chemical composition (in wt%) of the steels.

Ste
el

C

Mn

Si

P

S

Cr

Ni

Mo

V

N

Nb

P91

0.10

0.40

0.43

0.015

0.006

8.50

0.10

0.88

0.23

0.045

0.018

CSN

15

12
8

0.13

0.60

0.31

0.012

0.022

0.58

0.07

0.47

0.25

-

-

SK3STC

0.12

0.80

0.07

0.008

0.011

2.73

0.02

0.96

0.01

-

0.013


Table 2:

Weldment combinations,
annealing conditions and simulated phase region sequences. P
hase (used abr.):
BCC_A2 (α), M
7
C
3

(M7)
, M
23
C
6

(M23), M
6
C (M6)
, carbonitride (
MX). Symbol “
<” means: interface is moving
to the left with time; “>” means: moving to the right; “
|” means: movement

not observed.

Weldment
combinations

Temp. [ºC]

phase region sequence (predominant carbide/carbonitride bold)

CSN

15

128

/P91

500
-
561

(

+M7+
MX
) < (

+M23+M7+
MX
) < (

+M7+
MX
) < (

+
MX
) | (

+
M6
+MX)
>

(

+M6+
M7
+MX)
>

(

+
M23
+M6+M7+MX)
>

(

+
M23
+M6+MX)

CSN

15

12
8

/P91

561
-
577

(

+M7+
MX
) < (

+M23+M7+
MX
) < (

+M7+
MX
) < (

+
MX
) | (

+
M6
+MX)
>

(

+M6+
M7
+MX)
>

(

+
M23
+ M6+ M7+MX)

>

(

+
M23
+M7+MX)
>

(

+
M23
+MX)

CSN

15

128

/P91

577
-
655

(

+M23+M7+
MX
) < (

+M7+
MX
) < (

+
MX
) | (

+
M6
+MX)
>

(

+M6+
M7
+MX)
>

(

+
M23
+ M6+ M7+MX)
>

(

+
M23
+
M7+MX)
>

(

+
M23
+MX)

CSN

15

128

/P91

655
-
700

(

+
M23
+MX) < (

+
MX
) | (

+
M6
+MX)
>

(

+M6+
M7
+MX)
>

(

+
M23
+ M6+
M7+MX)
>

(

+
M23
+M7+MX)
>

(

+
M23
+MX)

SK3STCC /P91

500
-
554

(

+
M23
+M6) < (

+
M6
) < (

+
M23
+M6) | (

+M6+
M23
+MX)
>
(

+
M23
+MX+M6)
>
(

+M7+MX+
M23
+M6)
>
(

+MX+
M23
+M6)

SK3STCC /P91

554
-
556

(

+
M7
+M6+M23) < (

+
M23
+M6) < (

+
M6
) < (

+
M23
+M6) |
(

+M6+
M23
+MX)
>
(

+
M23
+MX+M6)
>
(

+M7+MX+
M23
+M6)
>
(

+MX+
M23
+M6)

SK3STCC /P91

556
-
561

(

+
M7
+M6) < (

+M7+M6+
M23
) < (

+
M23
+M6) < (

+
M6
) < (

+
M23
+M6) |
(

+M6+
M23
+MX)
>
(

+
M23
+MX
+M6)
>
(

+M7+MX+
M23
+M6)
>
(

+MX+
M23
+M6)

SK3STCC /P91

561
-
700

(

+
M7
+M6) < (

+M7+M6+
M23
) < (

+
M23
+M6) < (

+
M6
) < (

+
M23
+M6) |
(

+M6+
M23
+MX)
>
(

+
M23
+MX)
>
(

+M7+MX+
M23
)
>
(

+MX+
M23
)



9


Figure captions and figures

Figure 1:
Phase diagram cross
-
section for

steel P91.















Figure 2:
Phase diagram cross
-
section for steel SK3STC.













10


Figure 3:
Carbon activity of the steels at different temperatures.



Figure 4:

Simulated phase profiles for CSN

15

128/P91 weldment at 625ºC after 5000hrs .



11

Figure 5:

Simulated phase profiles for SK3STC/P91 weldment at 625ºC after 5000hrs .



Figure 6:

The simulated chemical composition of the MX carbonitride in dependence on distance
from interface (CSN

15

128/P91 weldments, 625ºC/5000hrs) .



12

Figure 7:

Th
e simulated chemical composition of the M23C6 carbide in dependence on distance
from interface (15128/P91 weldments, 625ºC/5000hrs) .



Figure 8:

The simulated and experimental overall concentration profiles at 625ºC/5000hrs.



Correspondence addresses

J
iří Sopoušek (corresponding author):

Masaryk University, Faculty of Science, Department of Theoretical and Physical Chemistry, Kotlářská 2, CZ
-
611 37 Brno, Czech Republic

sopousek
@chemi.muni.cz
, fax: 00420 541
211214