Report number: 1-FSFT
Evaluation of phase relations in weld
overlays of 316, 309MoL and SKWAM
Fredrik Stenarson
Fritjof Tibblin
Supervisors: Professor Malin Selleby, Tomislav Buzancic
and PhD Sten Wessman
2013
Dept. of Material Science and Engineering
Royal Institute of Technology
Stockholm, Sweden

2
Abstract
AREVA
NP
Uddcomb
AB
wants
to
replace
the
material
used
for
a
specific
valve
seat
used
in
boiling
water
reactors,
BWR.
Their
solution
is
a
weld
overlay
of
different
stainless
steels
composed
of
two
buffer
layers
of
the
steel
309
MoL
followed
by
two
layers
of
the
filler
material
SKWAM
welded
on
type
316
stainless
steel
or
carbon
steel.
The
report
focuses
on
the
long
term
structural
effects
in
the
weld
overlay
due
to
the
operating
temperature
in
BWRs,
in
this
case
270
°C.
To
investigate
the
thermodynamic
stability
in
the
weld
overlay
the
computer
software
Thermo-­‐Calc
was
used
and
a
metallographic
examination
was
carried
out.
The
results
from
these
procedures
were
compared
and
possible
long
term
effects
were
discussed.
Most
likely
spinodal
decomposition
is
the
most
severe
structural
change
that
may
appear
in
the
material.
At
equilibrium
conditions
at
the
operating
temperature
ferrite
is
decomposed
into
Fe-­‐rich
and
Cr-­‐rich
ferrite
but
since
the
kinetics
is
not
included
in
the
calculations
it
is
not
possible
to
determine
the
rate
of
decomposition.
Keywords:
SKWAM,
316,
309MoL,
475°C
embrittlement,
spinodal
decomposition,
intermetallic
phase,
welding
filler
material,
Thermo-­‐Calc,
stainless
steel,
thermodynamic
stability.

3
Table
of
Contents
Abstract
.........................................................................................................................................................
2
Introduction
..................................................................................................................................................
4
Background
...................................................................................................................................................
5
Materials
...................................................................................................................................................
5
Austenitic
stainless
steels
.....................................................................................................................
5
Ferritic
stainless
steels
..........................................................................................................................
5
Sigma
phase
..............................................................................................................................................
5
475
°C
embrittlement
...............................................................................................................................
6
Carbides
....................................................................................................................................................
8
Computations
and
experiments
....................................................................................................................
8
Thermo-­‐Calc
calculations
..........................................................................................................................
8
Metallographic
examination
.....................................................................................................................
9
Light
optical
microscope
.......................................................................................................................
9
Scanning
electron
microscope,
SEM
.....................................................................................................
9
Assumptions
............................................................................................................................................
10
Results
and
discussion
................................................................................................................................
11
Sigma
phase
at
300
°C
.............................................................................................................................
11
Carbides
at
300
°C
...................................................................................................................................
13
Spinodal
decomposition
at
300
°C
..........................................................................................................
14
Comparing
layers
in
the
weld
overlay
.................................................................................................
17
Metallographic
examination
...................................................................................................................
17
Sources
of
error
......................................................................................................................................
21
Conclusions
.................................................................................................................................................
22
Acknowledgements
.....................................................................................................................................
22
References
..................................................................................................................................................
23

4
Introduction
Cobalt
based
Stellite
alloys,
have
traditionally
been
used
as
hard
facing
materials
for
nuclear
plant
valves
(mainly
gate
valves)
owing
to
their
high
corrosion
resistance
and
superior
wear
resistance
under
sliding
conditions.
However,
the
need
to
avoid
the
use
of
Stellite
alloys
has
emerged
since
they
are
the
main
source
of
cobalt,
which
is
the
largest
contributor
to
the
occupational
radiation
exposure.
Isotope
cobolt59
,
which
may
be
released
from
cobalt
containing
surfaces
in
the
form
of
wear
and
corrosion
products,
is
transported
to
the
reactor
vessel
where
it
is
activated
to
the
radioactive
isotope
cobalt60
by
neutron
capture
in
the
fission
process.
In
the
light
of
these
findings
and
as
a
most
effective
way
to
reduce
cobalt
contamination,
many
cobalt-­‐free
hard
facing
alloys,
such
as
iron-­‐based
and
nickel-­‐based
alloys,
have
been
developed
in
order
to
replace
Stellite.
After
annual
routine
testing
of
the
BWRs
security
system,
cracks
were
detected
in
manually
welded
valve
seats.
An
investigation
took
its
start
to
find
a
new
material
combination
with
better
resistance
against
crack
formation.
Repeatedly
welding
maintenances
were
done
but
every
year
new
cracks
were
detected.
Under
normal
conditions
the
valve
seats
are
exposed
to
69
bar
and
270
°C
and
in
worst-­‐case
scenarios
80
bar
and
300
°C.
The
old
weld
overlay
consisted
of
SKWAM
welded
directly
on
carbon
steel.
In
this
case
the
structure
probably
gets
completely
martensitic
and
therefore
brittle.
The
new
material
combination
that
is
under
consideration
is
done
with
mechanized
welding
consist
of
a
base
material
of
type
316
covered
with
309MoL
and
a
few
layers
of
the
filler
material
SKWAM.
All
compositions
including
Stellite
6
can
be
found
in
table
1.
Type
316
and
309MoL
are
primarily
austenitic
and
SKWAM
is
predominantly
ferritic.
They
are
all
stainless
steels
and
are
highly
alloyed
with
chromium
and
none
of
them
contains
cobalt.
When
joining
these
grades
there
will
be
a
mixing
between
the
materials
and
new
phases
may
occur
that
can
cause
problems.
Table
1:
Composition
of
used
materials
in
wt%
Element
Carbon
steel
316
309MoL
SKWAM
Stellite
6
Fe
97.65
65.495
58.83
80.48
0
C
0.25
0.08
0.02
0.02
1.2
Si
0.5
0.75
0.45
0.7
0
Mn
1.6
2
1.5
0.7
0
Cr
0
17
21.5
17
30
Ni
0
12
15
0
0
Mo
0
2.5
2.7
1.1
0
S
0
0.03
0
0
0
P
0
0.045
0
0
0
N
0
0.1
0
0
0
Co
0
0
0
0
63.8
W
0
0
0
0
5

5
In
this
project
the
thermodynamic
stability
of
the
weld
overlay
was
investigated
with
Thermo-­‐Calc
a
software
for
equilibrium
calculations
[1],
metallographic
examination
of
weld
overlay
was
performed
and
possible
upcoming
phases
and
problems
that
may
occur
in
the
future
after
long
operation
times
were
discussed.
Evaluation
of
thermodynamic
stability
was
focused
on
different
compositions
depending
on
cooling
conditions,
which
phases
were
to
be
expected
after
long
periods
of
time
and
how
chromium
is
distributed
in
the
different
phases.
Metallographic
examination
was
made
with
focus
on
determining
the
phases
in
the
overlay
and
comparing
with
thermodynamic
calculations.
Background
Materials
The
materials
of
interest
are
309MoL,
316
and
SKWAM
which
are
all
stainless
steels.
SKWAM
is
a
product
name
and
the
others
are
material
groups.
309MoL
and
316
are
both
austenitic
steels
while
SKWAM
is
ferritic-­‐martensitic.
Austenitic
stainless
steels
Austenitic
stainless
steels
are
the
most
produced
and
largest
category
of
stainless
steels.
Generally
austenitic
steels
have
good
mechanical
properties
such
as
high
toughness
and
ductility.
The
corrosion
resistance
is
good
in
most
environments
but
decreases
when
exposed
to
elevated
temperatures,
the
maximum
service
temperature
is
approximately
760
°C.
The
austenite
phase
is
promoted
by
addition
of
nickel,
carbon,
nitrogen
and
manganese
where
the
most
important
addition
is
nickel.
Austenitic
stainless
steels
generally
contain
about
8-­‐20
wt%
nickel
but
some
austenitic
stainless
steels
are
free
of
nickel.
In
this
case
nickel
is
replaced
with
manganese
and
nitrogen.
Austenitic
stainless
steels
are
used
in
a
series
of
applications,
most
of
them
at
low
temperatures,
such
as
structural
support,
kitchen
equipment
and
medical
products.
Austenitic
stainless
steel
is
considered
to
have
good
weldability
[2].
Ferritic
stainless
steels
Ferritic
stainless
steels
consist
mainly
of
ferrite
phase.
Ferritic
stainless
steels
generally
have
better
corrosion
resistance
compared
to
austenitic
stainless
steels
but
do
not
have
as
good
mechanical
properties.
The
corrosion
resistance
does
not
depend
on
the
ferritic
phase
but
rather
the
chromium
and
molybdenum
content.
Ferritic
stainless
steels
are
used
for
applications
where
corrosion
resistance
is
more
important
than
good
mechanical
properties,
such
as
exhaust
systems
for
cars
and
in
chemical
industries.
The
ferrite
phase
in
stainless
steels
is
favored
by
high
chromium
and
molybdenum
contents
and
low
nickel
content.
Compared
to
austenitic
stainless
steels
they
are
cheaper
due
to
fewer
alloy
elements
but
they
are
relatively
more
expensive
since
they
are
hard
to
manufacture.
Ferritic
stainless
steels
are
sensitive
to
embrittlement,
such
as
475
°C
embrittlement
and
are
therefore
used
at
relatively
low
temperatures,
up
to
400
°C
but
as
low
as
280
°C
pressure
vessel.
The
weldability
of
ferritic
stainless
steel
is
not
as
good
as
for
austenitic
because
grain
growth
reduces
toughness
and
ductility
[2].
Sigma
phase
When
stainless
steels
are
exposed
to
elevated
temperatures
for
an
extended
period
of
time
intermetallic
phases
may
precipitate
e.g.
sigma
phase.
The
sigma
phase
consists
of
mainly
iron
and
chromium
but
its

6
composition
is
varying
depending
on
the
alloying
elements.
It
is
precipitated
when
heat
treated
at
about
570-­‐1000
°C [2].
The
same
applies
for
the
interpass
temperature
during
welding
and
the
heat
input
should
be
minimized
[8].
Precipitation
of
sigma
phase
causes
embrittlement
of
the
material
since
the
sigma
particles
are
harder
than
the
surrounding
matrix
and
therefore
reduces
the
ductility
and
toughness.
Since
the
sigma
phase
contains
15-­‐70
%
chromium
it
is
most
likely
that
precipitation
will
occur
in
chromium
rich
environments
within
the
material
e.g.
the
ferrite
phase
[2].
Generally
the
chromium
content
needs
to
be
above
20
wt%
for
the
precipitation
to
take
place
and
if
the
chromium
content
is
raised
to
25-­‐30
wt%
sigma
phase
forms
rapidly
[3].
The
main
transformation
mechanism
for
the
precipitation
of
sigma
phase
is
the
transformation
of
ferrite
to
sigma
phase.
Sigma
phase
will
form
directly
in
chromium
rich
regions
of
the
ferrite
grains.
It
is
possible
for
sigma
phase
to
form
in
austenite
but
it
is
not
as
usual
since
it
is
harder
for
chromium
to
diffuse
in
FCC
than
BCC.
Except
for
chromium
other
ferrite
stabilizing
elements
such
as
silica
and
molybdenum
will
accelerate
the
formation
of
sigma
phase [3].
475
°C
embrittlement
It
is
mainly
ferritic
stainless
steels
that
experience
475
°C
embrittlement
if
they
are
exposed
to
temperatures
in
the
interval
of
425-­‐550
°C
[2].
475
°C
embrittlement
only
takes
place
in
stainless
steels
in
the
ferritic
phase
during
annealing
around
475
°C
[5].
Most
common
is
that
475
°C
embrittlement
does
not
occur
when
welding
because
long
time
exposure
to
high
temperatures
is
required.
It
will
therefore
be
important
to
know
the
environment,
such
as
the
temperature
range
the
material
will
be
exposed
to
in
its
application
[2].
A
broader
framing
of
the
material
suggests
that
475
°C
embrittlement
can
occur
in
steels
that
are
ferritic,
austenitic-­‐ferritic
and
in
filler
materials
that
contain
δ-­‐ferrite.
475
°C
embrittlement
is
due
to
spinodal
decomposition
[5].
Alloying
elements
affect
time
and
temperature
for
the
maximum
embrittlement
of
the
ferritic
phase.
Silicon,
aluminum,
chromium
and
molybdenum
do
all
accelerate
the
maximum
embrittlement.
Carbon
has
the
opposite
effect
and
reduces
the
maximum
effect
of
embrittlement
when
forming
chromium-­‐
carbides.
Alloys
that
contain
titanium
and
niobium
form
stable
carbides
before
chromium-­‐carbides
are
formed
so
the
embrittlement
effect
is
enhanced
as
long
as
there
are
such
stable
carbides
formed
instead
of
chromium-­‐carbides
[5].
Nitrogen
and
manganese
seems
to
have
no
impact
on
the
475
°C
embrittlement,
while
nickel
increases
the
effect
[5], [6].
The
dominant
theory
of
why
475
°C
embrittlement
occurs
is
the
coherent
precipitate
below
550
°C
because
of
the
miscibility
gap
in
the
iron-­‐chromium
phase
diagram
as
can
be
seen
in
Fig
1.
Iron-­‐
chromium
alloys
with
compositions
in
the
range
of
the
miscibility
gap
and
being
annealed
below
550
°C
tend
to
precipitate
two
phases,
α-­‐ferrite
and
α’-­‐ferrite.
α-­‐ferrite
is
an
iron-­‐rich
phase
with
BCC-­‐lattice.
α’-­‐
ferrite
is
a
chromium-­‐rich
phase
with
BCC-­‐lattice,
which
contains
about
61-­‐83
%
chromium
and
is
nonmagnetic
[2].
The
two
phases
are
said
to
have
different
morphologies
where
the
newly
formed
phase
of
α’-­‐ferrite
is
embedded
in
the
chromium
depleted
α-­‐ferrite.
There
are
two
ways
for
α’-­‐ferrite
to
form,
either
through
nucleation
and
growth
or
through
spinodal
decomposition
[7].

7
Figure
1.
Parts
of
the
iron-­‐chromium
system [1].
Velocity
and
rate
of
embrittlement
can
be
seen
as
a
function
of
the
chromium
content.
High
chromium
content
and
higher
temperatures
results
in
more
embrittlement
whereas
stainless
steels
with
low
chromium
content
can
be
almost
exempt
from
475
°C
embrittlement
[2].
In
this
mechanism,
activation
energy
of
aging
is
similar
to
the
activation
energy
of
Cr
diffusion
in
the
ferrite
phase.
The
kinetics
for
475
°C
embrittlement
precipitation
can
be
tested
by
measuring
the
hardness
and
impact
strength
in
ferrite
with
Charpy-­‐V.
The
kinetics
of
the
embrittlement
can
be
of
significant
importance
in
certain
construction
parts
in
BWRs
[6].
Studies
of
both
ferritic-­‐
and
duplex
stainless
steels
have
shown
that
spinodal
decomposition
is
faster
in
duplex
steels.
Radiation
has
been
found
to
accelerate
the
spinodal
decomposition
and
also
effect
volume
fraction
and
morphology
[7].
Cold
working
affects
stainless
steels
so
that
precipitation
of
α’-­‐ferrite
increases
which
accelerates
the
embrittlement.
475
°C
embrittlement
also
makes
the
steel
less
resistant
to
corrosion
since
the
chromium
depleted
α-­‐ferrite
is
particularly
susceptible
to
corrosion.
There
are
some
alternatives
to
reduce
the
embrittlement
and
restore
the
mechanical
and
corrosion
properties.
By
heat
treatment
of
the
embrittled
material
in
the
temperature
interval
of
550-­‐600
°C
for
a
short
period
of
time
the
original
properties
of
the
stainless
steel
can
be
restored
and
α-­‐ferrite
and
α’-­‐ferrite
can
form
ferrite
again [2].
There
will
be
more
475
°C
embrittlement
in
materials
with
high
chromium
content
when
it
has
been
exposed
to
elevated
temperatures
for
long
periods
of
time.
Therefore
stainless
steels
with
high
chromium
content
should
not
be
heat-­‐treated
at
too
high
temperature[8].

8
Carbides
In
cases
when
the
amount
of
carbides
in
austenitic
stainless
steels
is
critical,
a
solution
annealing
treatment
can
bring
carbides
back
into
solution.
By
quenching,
a
low
amount
of
carbides
can
be
obtained.
The
alloying
elements
that
precipitated
as
carbides
earlier
are
now
in
a
non-­‐equilibrium
state.
Depending
on
how
high
temperature
the
stainless
steel
will
be
exposed
to
in
its
application
the
diffusion
coefficient
changes
and
the
kinetics
for
alloying
elements
determine
if
stable
carbides
can
form
ones
again
[5].
For
stainless
steels
carbon
has
low
solubility
at
low
temperatures.
Excess
of
carbon
may
result
in
precipitation
of
iron-­‐chromium-­‐carbides
such
as
M23C6
and
M6C.
The
chromium
content
in
M23C6
is
often
in
the
range
of
42-­‐65
wt%.
Since
the
chromium
content
in
M23C6
is
two
to
four
times
as
much
as
the
average
matrix
content
the
close
surroundings
of
M23C6
will
be
depleted
in
chromium.
Variation
of
chromium
content
can
be
evened
out
by
heat
treatment.
During
heat
treatment
the
temperature
should
be
higher
than
the
temperature
range
where
M23C6
is
precipitated
otherwise
the
diffusion
for
chromium
and
iron
is
to
slow.
Precipitation
of
M23C6
is
mostly
concentrated
to
grain
boundaries
which
make
adjacent
areas
chromium
depleted.
Chromium
contents
below
11,5
wt%
increases
the
risk
of
corrosion.
Since
chromium
depletion
is
concentrated
to
grain
boundaries
activation
potential
of
intergranular
corrosion
increases
and
propagation
will
progress
along
chromium
depleted
grain
boundaries
[5].
Increased
carbon
content
increases
the
risk
of
intergranular
corrosion.
Nickel
contributes
to
increased
precipitation
of
M23C6
because
it
reduces
the
solubility
of
carbon
and
increases
the
carbon
activity.
Silicon
influence
carbide
precipitation
the
same
way
as
nickel
but
with
stronger
effect.
M23C6
precipitation
is
mildly
affected
of
increased
chromium
content,
the
intergranular
corrosion
resistance
increases
since
the
closest
surroundings
have
enhanced
chromium
content.
Molybdenum
reduces
carbon
solubility
and
carbides
can
be
precipitated
to
a
greater
extent.
Manganese
increases
carbon
solubility
and
reduces
carbon
activity
but
seems
to
have
no
influence
on
corrosion
resistance
[5].
Computations
and
experiments
Thermo-­‐Calc
calculations
The
composition
of
each
layer
in
the
weld
overlay
was
calculated
from
a
principle
of
70
%-­‐30
%
mixing
between
layers.
Each
layer
composition
was
input
data
in
Thermo-­‐Calc
3.0
beta
2,
thermodynamic
calculations
were
made
and
output
data
such
as
plots
and
tables
were
extracted.
Thermo-­‐Calc
was
set
to
use
TCFE7 [1]
database
in
all
calculations.
Calculations
were
carried
out
with
the
main
goal
to
extract
plots
of
stable
phases
in
each
layer
considering
two
different
methods,
70
%-­‐30
%
principle
and
70
%-­‐30
%
principle
with
subsequent
Scheil
calculations
when
3
%
melt
remained.
The
composition
used
in
the
last
method
was
the
composition
of
the
melt
when
97
%
had
solidified.
Reality
is
expected
to
be
somewhere
between
equilibrium
of
the
70
%-­‐30
%
principle
and
70
%-­‐30%
with
subsequent
Scheil
calculations.
Focus
has
also
been
on
determining
how
much
chromium
that
was
expected
to
be
distributed
in
each
phase
since
it
affects
475
°C
embrittlement.
Chromium
distribution
diagrams
for
each
phase
and
all
layers
were
calculated
using
Thermo-­‐Calc.
The
driving
force
for
precipitation
of
other
phases
than
BCC
and
FCC
were
also
calculated
with
Thermo-­‐Calc.
By
using
previous
equilibrium

9
calculations
stable
phases
at
300
°C
were
detected.
All
stable
phases
except
BCC
and
FCC
were
excluded
in
equilibrium
calculations
but
still
the
driving
force
was
calculated,
the
rest
was
suspended
from
all
calculations.
Metallographic
examination
The
examined
sample
was
a
weld
overlay
consisting
of
a
base
layer
of
carbon
steel
covered
with
two
buffer
layers
of
309
MoL
and
two
layers
of
SKWAM.
Light
optical
microscope
Sample
preparation
started
with
cutting
a
piece
of
the
weld
overlay
with
a
saw.
Then
the
piece
was
casted
in
a
polymer
matrix.
The
piece
was
later
grinded
with
two
different
papers
and
later
polished.
Last
step
in
the
preparation
was
etching
with
a
10
%
solution
of
electrolytic
chromic
acid
until
phases
could
be
easily
detected.
The
sample
was
examined
with
light
optical
microscope
and
pictures
were
taken
to
examine
included
phases,
for
further
discussions
and
results.
Scanning
electron
microscope,
SEM
The
polished
and
etched
sample
was
put
in
a
beaker
containing
ethanol.
This
was
done
to
clean
the
sample.
Then
the
sample
was
horizontally
fixated
with
conducting
clay.
The
sample
was
then
examined
with
a
Hitachi
S-­‐3700N
scanning
electron
microscope.
The
composition
in
each
layer
was
measured
using
the
software
Brunker
Quantax
800.
Also
a
picture
of
each
layer
was
taken
with
the
SEM.
Carbon
and
nitrogen
cannot
be
measured
with
this
instrument
since
these
elements
are
too
light.
Figure
2.
Schaeffler-­‐diagram,
phases
to
be
expected
in
each
layer
of
the
weld
overlay.
[9]

10
Assumptions
The
weld
overlay
consists
of
several
layers
that
will
interact
during
welding.
Whenever
a
new
layer
is
added
the
heat
will
partially
melt
the
base
layer
and
the
two
will
mix.
In
this
report
the
mixture
is
assumed
to
be
70
%-­‐30
%
between
the
layers.
This
means
that
if
material
A
is
welded
onto
material
B
the
new
layer
will
consist
of
70
%
material
A
and
30
%
material
B.
This
percentage
was
used
after
recommendations
from
AREVA
NP
Uddcomb
AB.
It
is
also
assumed
that
the
mixture
is
70
%-­‐30
%
all
over
the
layer.
The
operating
temperature
for
the
valve
seat
in
a
BWR
is
about
270
°C
and
the
calculated
worst-­‐case
scenario
gives
a
temperature
of
about
300
°C.
All
tables
in
this
report
are
calculated
at
300
°C.
The
most
interesting
temperature
is
the
operating
temperature
since
this
report
is
focusing
on
long
term
effects
but
since
the
difference
between
operating
and
worst
case
temperature
is
small
and
temperatures
are
low
it
is
assumed
that
300
°C
is
representative.
During
operation
the
valve
seat
experiences
a
pressure
of
about
69
bar
and
in
the
worst-­‐case
scenario
the
pressure
increases
to
about
80
bar.
In
this
report
it
is
assumed
that
the
pressure
does
not
affect
the
calculations
and
all
calculations
are
done
using
atmospheric
pressure.
Calculations
using
Thermo-­‐Calc
with
a
pressure
of
80
bar
were
carried
out
and
there
was
negligible
difference
as
when
carried
out
with
atmospheric
pressure.
During
the
Scheil
calculations
it
was
assumed
that
carbon
is
fast
diffusing.
From
the
Scheil
calculations
the
composition
of
the
liquid
phase
were
acquired,
which
was
used
to
create
plots
of
stable
phases.
The
composition
that
was
used
in
this
report
is
for
the
liquid
phase
when
97
%
of
the
system
is
solid.
In
this
case
it
is
assumed
that
the
diffusion
rate
will
be
low
and
the
remaining
3
%
will
solidify
with
another
composition
than
the
rest
of
the
system.
This
composition
is
assumed
to
be
a
worst-­‐case
scenario.
It
is
taken
into
account
that
it
is
not
possible
to
perform
calculations
on
diffusion
free
phase
transformations
using
Thermo-­‐Calc.
In
this
particular
case
irradiation
effects
on
the
weld
overlay
can
be
excluded
since
the
valve
seat
is
situated
in
an
area
of
the
plant
with
low
radiation.
This
assumption
was
made
after
discussions
with
AREVA
NP
Uddcomb
AB.

11
Results
and
discussion
Sigma
phase
at
300
°C
After
evaluation
of
Fig
3
it
can
be
stated
that
the
sigma
phase
has
no
thermodynamic
stability
at
300
°C.
Therefore
sigma
phase
will
not
be
precipitated
even
after
long
periods
of
time
at
300
°C.
If
any
sigma
phase
is
present
it
has
been
an
effect
from
the
welding
thermal
cycle
but
sigma
phase
precipitate
after
long
time
and
welding
usually
concerns
rapid
cooling.
Buffer
layer
First
SKWAM
layer
Second
SKWAM
layer
Third
SKWAM
layer
Figure
3.
Amount
vs.
temperature
of
all
stable
phases
for
all
layers. [1]

12
The
equilibrium
calculations
using
the
composition
from
Scheil
calculations
of
the
third
SKWAM
layer
when
3
%
melt
remains
show
that
sigma
phase
is
thermodynamically
stable,
as
can
be
seen
in
Fig
4.
Even
if
the
sigma
phase
would
precipitate
in
this
layer
the
volume
of
sigma
phase
would
be
small
considering
the
whole
sample,
also
the
kinetics
of
the
reaction
must
be
taken
in
to
account
since
the
temperature
is
low.
Buffer
layer
First
SKWAM
layer
Second
SKWAM
layer
Third
SKWAM
layer
Figure
4:
Amount
vs.
temperature
of
all
stable
phases
for
liquid
composition
in
all
layers
during
Scheil
calculations
when
3
%
of
the
system
is
in
liquid
phase. [1]

13
Carbides
at
300
°C
The
carbon
content
is
decreasing
from
base
material
to
top
layer.
At
equilibrium
the
amount
of
carbides
at
300
°C
follows
the
carbon
content
tendency.
Carbides
M23C6
and
M6C
are
both
thermodynamically
stable
but
the
total
amount
of
them
never
surpass
1
mole%.
M23C6
and
M6C
do
not
seem
to
coexist
in
the
same
layer
at
equilibrium.
In
the
buffer
layer
and
the
first
SKWAM
layer
M23C6
is
thermodynamically
stable,
for
the
second
and
the
third
SKWAM
layers
M6C
is
thermodynamically
stable.
The
equilibrium
calculation
using
the
composition
from
the
Scheil
calculations
when
3
%
of
melt
remains
shows
that
all
layers
contain
a
higher
amount
of
carbides,
both
M23C6
and
M6C
can
coexist.
The
increased
concentration
of
carbides
is
an
effect
of
about
six
time’s
higher
carbon
content.
Even
though
the
carbide
content
is
high
in
the
3
%
melt
the
carbide
concentration
in
the
whole
sample
is
low.
The
amount
of
carbon
and
carbides
decreases
from
the
base
layer
to
the
top
layer
in
the
same
way
as
in
the
calculations
at
equilibrium.
The
opposite
is
true
considering
the
driving
force
for
precipitation
of
carbides.
The
driving
force
for
carbide
precipitation
increases
from
the
base
layer
to
the
top
layer
as
seen
in
table
2.
Even
though
carbides
are
thermodynamically
stable
at
300
°C
AREVA
NP
Uddcomb
AB
has
not
had
any
problem
with
carbides
in
the
weld
overlay.
This
states
that
no
substantial
amount
is
formed
during
welding
and
that
the
kinetics
is
slow
at
the
operating
temperature.
Table
2:
Driving
force
for
precipitation
of
carbides
at
300
°C
for
each
layer.
Buffer
layer
1st
SKWAM
layer
2nd
SKWAM
layer
3rd
SKWAM
layer
M23C6
0,045
0
0
3,7
M6C
0
2,09
3,14
3,14

14
Spinodal
decomposition
at
300
°C
Spinodal
decomposition
is
thermodynamically
stable
at
300
°C.
If
the
weld
overlay
reaches
equilibrium
Fe-­‐rich
BCC
and
Cr-­‐rich
BCC
will
be
dominating
phases
in
all
layers
which
can
be
seen
in
Fig
3.
If
the
system
reaches
equilibrium
the
absolute
majority
of
the
total
Cr-­‐content
will
be
in
the
Cr-­‐rich
BCC
phase
as
can
be
seen
in
Fig
5.
Figure
5.
Weight
percentage
of
total
chromium
content
in
α’-­‐ferrite
for
all
layers
at
300
°C
when
equilibrium
is
reached.
67.7%
77.1%
78.2%
78.2%
0.0%
20.0%
40.0%
60.0%
80.0%
100.0%
Buffer
layer
First
SKWAM
layer
Second
SKWAM
layer
Third
SKWAM
layer
Weight-­‐%
of
total
Cr-­‐content
in
Cr-­‐rich
BCC
at
equilibrium
for
each
layer

15
Fig
6
is
displaying
the
Cr-­‐content
in
all
phases
for
the
different
layers.
The
Cr-­‐content
in
the
Cr-­‐rich
BCC
is
increasing
while
decreasing
in
the
Fe-­‐rich
BCC.
This
shows
that
there
is
a
driving
force
for
spinodal
decomposition
as
the
temperature
decreases.
Buffer
layer
First
SKWAM
layer
Second
SKWAM
layer
Third
SKWAM
layer
Figure
6.
Amount
of
Cr
in
all
stable
phases
at
equilibrium
for
all
layers. [1]
Considering
only
the
thermodynamics
the
separation
between
iron
and
chromium
into
two
different
BCC
phases
will
be
greater
as
temperature
decreases.
Fig
7
shows
that
there
is
a
substantial
amount
of
Cr-­‐
rich
BCC
in
all
layers.
It
also
shows
that
the
amount
of
Cr-­‐rich
BCC
stays
basically
the
same
even
though
the
total
Cr-­‐amount
in
each
layer
is
decreasing
towards
the
third
SKWAM
layer.
The
decreased

16
chromium
content
in
the
SKWAM
layers
is
probably
compensated
by
an
increased
amount
of
ferrite,
which
can
decompose.
In
the
buffer
layer
a
large
part
of
the
system
is
austenite.
Figure
7:
Amount
of
Cr-­‐rich
and
Fe-­‐rich
BCC
in
each
layer
at
equilibrium.
Even
though
the
thermodynamics
states
that
the
ferrite
should
be
separated
into
one
Cr-­‐rich
and
one
Fe-­‐rich
phase
at
300
°C
the
calculations
do
not
consider
the
kinetics
for
the
reactions.
For
instance
it
is
not
likely
to
have
spinodal
decomposition
right
after
welding
since
high
temperatures
under
longer
periods
of
time
is
required.
In
reality
the
reaction
for
spinodal
decomposition
is
slow
and
requires
chromium
diffusion
in
solid
state.
The
valve
seats
within
the
nuclear
plant
will
be
exposed
to
a
somewhat
elevated
temperature,
270
°C
under
normal
circumstances,
which
will
enhance
spinodal
decomposition
but
it
still
is
below
the
most
critical
temperatures.
The
most
critical
temperature
according
to
literature
is
approximately
475
°C,
the
reaction
rate
for
spinodal
decomposition
is
highest
at
this
temperature.
The
operating
temperature
for
the
valve
seat
is
lower
than
475
°C
but
since
nuclear
plants
run
day
and
night
all
year
around
it
will
be
exposed
to
this
elevated
temperature
for
long
periods
of
time.
With
all
certainty
the
kinetics
is
lower
at
the
operating
temperature
but
since
all
calculations
in
this
project
is
done
assuming
equilibrium
it
is
not
possible
to
determine
the
decomposition
rate
at
270
°C.
The
results
from
the
Scheil
calculations
are
not
relevant
when
talking
about
spinodal
decomposition
since
the
composition
used
in
calculations
only
represent
the
3
%
of
liquid
phase
remaining.
The
composition
of
the
remaining
97
%
that
is
solidified
has
almost
the
same
composition
as
at
the
original
composition
and
is
assumed
to
behave
the
same
way.
0,151
0,155
0,154
0,152
0,575
0,785
0,833
0,837
0
0,1
0,2
0,3
0,4
0,5
0,6
0,7
0,8
0,9
Buffer
layer
1st
SKWAM
layer
2nd
SKWAM
layer
3rd
SKWAM
layer
Cr-­‐rich
BCC
Fe-­‐rich
BCC

17
Comparing
layers
in
the
weld
overlay
According
to
Fig
7
the
spinodal
decomposition
is
similar
in
all
SKWAM
layers,
but
the
buffer
layer
differs
and
has
lower
amount
of
Fe-­‐rich
BCC.
The
main
reason
that
the
buffer
layer
does
not
contain
much
Fe-­‐
rich
BCC
is
because
there
is
large
amount
of
austenite
present,
which
does
not
decompose.
Since
the
Cr-­‐
content
is
higher
in
the
buffer
layer
it
suggests
that
the
amount
of
Cr-­‐rich
BCC
should
be
higher
compared
to
the
SKWAM
layers.
But
it
follows
the
opposite
trend,
the
buffer
layer
does
contain
more
chromium
but
much
of
it
is
found
in
austenite
and
other
Cr-­‐rich
phases.
Fig
5
shows
that
in
the
buffer
layer
less
chromium
are
absorbed
in
Cr-­‐rich
BCC.
In
the
SKWAM
layers
lower
amount
of
austenite
is
found
and
other
Cr-­‐rich
phases
are
also
found
in
smaller
amounts,
this
result
in
more
BCC.
The
general
trend
for
all
layers
is
that
Fe-­‐rich
BCC
is
reduced
and
more
Cr-­‐rich
BCC
is
precipitated
at
lower
temperatures
as
can
be
seen
in
Fig
3.
Metallographic
examination
In
Fig
2
the
composition
for
each
layer
is
pointed
out
in
a
Schaeffler-­‐diagram.
In
the
metallographic
examination
no
precise
determination
of
the
amount
of
each
phase
was
performed
so
there
can
only
be
a
brief
discussion
of
expected
and
actual
precipitated
phases.
• All
layers:
An
overview
of
all
layers
in
the
weld
overlay
can
be
seen
in
Fig
8.
Figure
8.
All
layers.
Magnification
x12.5.

18
• Buffer
layer
1:
From
Schaeffler-­‐diagram,
100
%
austenite
was
to
be
expected
and
Fig
9
show
that
there
is
probably
a
few
percent
of
ferrite
present
in
the
sample.
Figure
9.
First
buffer
layer,
material
309MoL.
Magnification
x200.
• Buffer
layer
2:
From
Schaeffler-­‐diagram,
5
%
ferrite
and
95
%
austenite
were
to
be
expected
and
Fig
10
shows
that
austenite
and
ferrite
are
present.
The
two
buffer
layers
has
approximately
same
ratio
between
austenite
and
ferrite.
Figure
10.
Second
buffer
layer,
material
309MoL.
Magnification
x200.
White areas:
Dendrite of austenite
Dark areas:
Ferrite
Dark areas:
Primary
precipitation
of ferrite
White areas:
Dendrites of
austenite

19
• SKWAM
layer
1:
From
Schaeffler-­‐diagram,
a
mixture
of
austenite,
ferrite
and
martensite
with
approximately
80
%
ferrite
can
be
expected.
The
ratio
is
hard
to
determine
from
Fig
11
but
it
is
clear
that
ferrite,
austenite
and
martensite
is
present.
Ferrite
seems
to
be
the
dominating
phase.
Figure
11.
First
SKWAM
layer.
Magnification
x500.
• SKWAM
layer
2:
From
Schaeffler-­‐diagram,
only
ferrite
should
be
present.
Fig
12
shows
that
there
are
three
phases
present,
ferrite,
austenite
and
martensite.
Ferrite
is
the
dominating
phase.
Figure
12.
Second
SKWAM
layer.
Magnification
x100.
White areas:
Ferrite
Grey areas:
Dendrites of
Austenite
Dark areas:
Martensite
White area:
Ferrite
Grey/Dark area:
Martensite and
austenite

20
15,05
19,57
17,77
17,23
9,29
17,71
17,16
16,96
Calculated
wt%
Cr
in
each
layer
SEM
valvue
of
wt%
Cr
in
each
layer
Buffer
layer
SKWAM
1
SKWAM
2
SKWAM
3
Martensite
is
present
in
all
SKWAM
layers
as
seen
in
Fig
11
and
12.
It
is
not
possible
to
see
martensite
in
the
Thermo-­‐Calc
calculations
since
it
is
not
thermodynamically
stable
but
if
precipitated
the
decomposition
is
slow.
Martensite
is
an
effect
of
welding
and
rapid
cooling
from
the
austenitic
region.
During
operation
in
the
nuclear
plant
more
martensite
will
not
form
in
the
SKWAM
layers
since
rapid
cooling
from
high
temperatures
is
required
to
form
martensite.
In
this
case
the
valve
seat
will
be
exposed
to
a
somewhat
elevated
temperature
for
a
long
time
but
not
high
enough.
Fig
13
shows
that
the
assumption
of
a
70
%-­‐30
%
mixture
is
quite
accurate
for
the
chromium
content
in
each
layer.
Figure
13.
Calculated
wt%
chromium
with
70
%-­‐30
%
mixture
in
each
layer
of
the
examined
sample
and
measured
wt%
chromium
from
SEM.

21
The
calculated
values
and
the
values
measured
with
SEM
for
all
elements
are
summarized
in
table
3.
Table
3:
Calculated
values
from
70
%-­‐30
%
mixture
and
composition
of
elements
using
SEM.
Fe
Si
Mn
Cr
Ni
Mo
Buffer
layer
1
calculated
70,47
0,47
1,53
15,05
10,5
1,89
Buffer
layer
1
from
SEM
82,07
0,18
1,06
9,29
6,65
0,76
Buffer
layer
2
calculated
62,32
0,45
1,51
19,57
13,65
2,46
Buffer
layer
2
from
SEM
66,22
0,3
1,47
17,71
11,55
2,67
SKWAM
layer
1
calculated
75,03
0,63
0,94
17,77
4,1
1,51
SKWAM
layer
1
from
SEM
72,66
0,33
0,84
17,16
7,4
1,55
SKWAM
layer
2
calculated
78,85
0,68
0,77
17,23
1,23
1,22
SKWAM
layer
2
from
SEM
78,91
0,45
0,6
16,96
2,04
0,95
Sources
of
error
The
mixture
between
layers
in
the
weld
overlay
was
assumed
to
be
exactly
70
%-­‐30
%
mixture.
It
is
unreasonable
that
the
mixture
is
exactly
70
%-­‐30
%
in
the
whole
layer.
Reasonable
is
that
the
area
closest
to
the
layer
beneath
is
more
mixed
than
at
the
top
of
the
new
layer,
as
a
gradient.
During
Scheil
calculations
it
was
assumed
that
the
melt
segregates
until
3
%
of
the
melt
is
remaining.
When
the
3
%
melt
remains
calculations
were
aborted
because
otherwise
temperature
of
solidification
would
be
unrealistically
low.
3
%
melt
were
discussed
with
our
supervisors
and
it
was
decided
that
it
was
a
reasonable
amount.
This
was
thought
to
be
a
worst
case
scenario
for
the
weld
overlays
composition.
When
the
Scheil
calculations
were
performed
carbon
was
assumed
to
be
a
fast
diffusing
element
because
of
its
small
size.
Since
nitrogen
has
approximately
the
same
size
as
carbon
it
is
possible
that
it
also
should
have
been
considered
to
be
fast
diffusing.
All
calculations
in
Thermo-­‐Calc
were
done
with
the
constitution
of
three
SKWAM
layers
and
one
buffer
layer.
Unfortunately
the
samples
from
AREVA
NP
Uddcomb
AB
consisted
of
two
SKWAM
layers
and
two
buffer
layers.
Also
the
base
material
was
carbon
steel
instead
of
stainless
steel
type
316.
This
makes
the
comparison
between
the
calculations
and
the
samples
less
meaningful.
The
samples
from
AREVA
NP
Uddcomb
AB
have
not
been
in
operation
in
a
nuclear
plant.
Comparing
samples
with
the
calculations
makes
them
less
accurate
since
calculations
are
focusing
on
long-­‐term
effect
due
to
an
elevated
temperature.
The
samples
only
show
the
structure
right
after
welding.

22
Conclusions
The
method
was
to
perform
equilibrium
calculations
using
Thermo-­‐Calc
to
gain
information
on
which
phases
that
are
present
in
the
different
layers
of
this
particular
weld
overlay.
A
metallographic
examination
was
carried
out
to
compare
the
calculations
with
the
samples.
One
shortcoming
in
this
project
was
that
the
sample
that
was
examined
has
not
been
in
operation
and
because
of
that
no
long-­‐
term
effects
could
be
observed.
One
way
to
improve
the
method
would
be
to
use
a
sample
that
had
been
in
operation.
During
the
metallographic
examination
martensite
was
observed
in
the
SKWAM
layers.
This
was
assumed
to
be
an
effect
from
welding
and
is
not
possible
to
predict
using
Thermo-­‐Calc.
Since
martensite
will
influence
the
mechanical
properties
of
the
valve
seat
an
improvement
would
be
to
find
a
way
to
predict
the
amount
of
martensite
formed.
Among
the
thermodynamic
effects
that
occur
after
long
time
exposure
to
the
operating
temperature
spinodal
decomposition
seems
to
be
the
most
severe.
At
equilibrium
the
spinodal
decomposition
is
extensive
but
in
the
calculations
performed
in
Thermo-­‐Calc
the
kinetics
was
not
considered.
This
is
a
shortcoming
with
the
method
and
to
get
more
accurate
results
kinetic
calculations
should
be
performed.
For
example
if
the
kinetics
for
the
spinodal
decomposition
at
the
operating
temperature
is
slow
this
might
not
be
a
problem
but
it
can
have
large
impact
on
the
mechanical
properties
if
the
kinetics
is
fast.
The
chromium
composition
is
crucial
for
the
spinodal
decomposition
since
it
is
depending
on
chromium
diffusion.
By
using
SEM
the
calculated
wt%
of
chromium
in
each
layer
could
be
controlled.
Fig
13
shows
that
the
approximation
is
good
when
the
composition
between
layers
is
similar
but
between
the
carbon
steel
and
the
highly
alloyed
buffer
layer
the
difference
is
large.
The
method
also
offers
some
advantages.
Phases
that
do
not
exist
in
the
weld
overlay
after
welding
can
be
disregarded
if
they
are
not
thermodynamically
stable
at
the
operation
temperature.
For
example
the
sigma
phase
will
not
be
a
problem
in
this
case
since
it
is
not
stable
at
the
operating
temperature
according
to
Fig
3
and
was
not
detected
in
the
samples.
Using
this
method
it
is
possible
to
exclude
several
phases
but
not
to
get
an
exact
result.
The
most
important
improvement
in
this
case
would
be
to
learn
more
about
the
kinetics
for
spinodal
decomposition
at
the
operating
temperature.
Acknowledgements
Thanks
for
all
help
and
support
from
supervisors’
professor
Malin
Selleby
and
PhD
Sten
Wessman
at
Dept.
of
Material
Science
and
Engineering
at
KTH.
Thanks
to
Wenli
Long
for
your
help
with
SEM.
For
helping
us
with
the
preparation
of
the
samples
thanks
to
Ian
Patterson
and
Jonas
Guldbrandsson
for
demonstrating
welding
procedure
at
AREVA
NP
Uddcomb
AB.
Most
of
all
thanks
to
Tomislav
Buzancic
for
assigning
us
this
project,
support
and
the
field
trip
to
AREVA
NP
Uddcomb
office
in
Karlskrona.

23
References
[1]
Thomas
Helander,
Lars
Höglund,
Pingfang
Shi,
Bo
Sundman
J-­‐O
Andersson,
"THERMO-­‐CALC
&
DICTRA,
Computational
Tools
For
Materials
Science,"
Calphad,
vol.
26,
pp.
273-­‐312,
2002.
[2]
John
C.
Lippold
and
Damian
J
Kotecki,
Welding
metallurgy
and
weldability
of
stainless
steels.
Hoboken,
NJ,
USA:
John
Wiley,
2005.
[3]
Chih-­‐Chun
Hsieh
and
Weite
Wu,
"Overview
of
Intermetallic
sigma
phase
precipitation
in
stainless
steels,"
ISRN
Metallurgy,
vol.
2012,
january
2012.
[4]
T.H.
Chen
and
J.R.
Yang,
"Effects
of
solution
treatment
and
continuous
cooling
on
sigma-­‐phase
precipitation
in
a
2205
duplex
stainless
steel,"
Materials
science
and
engineering,
vol.
311,
2000.
[5]
Erich
Folkhard,
Welding
metallurgy
of
stainless
steels.
Wien,
Austria:
Spinger-­‐Verlag,
1988.
[6]
J.K.
Sahu,
U.
Krupp,
R.N.
Ghosh,
and
H.J.
Christ,
"Effect
of
475
°C
embrittlement
on
the
mechanical
properties
of
duplex
stainless
steels,"
Materials
science
and
engineering,
vol.
508,
pp.
1-­‐14,
january
2009.
[8]
Kobe
Steel,
LTD,
Arc
welding
of
specific
steels
and
cast
irons,
4th
ed.
Tokyo,
Japan:
Kobe
steel,
LTD,
2011.
[7]
K.H.
Lo,
C.H.
Shek,
and
J.K.L.
Lai,
"Recent
development
in
stainless
steels,"
Physics
and
Materials
Science,
City
University
of
Hong
Kong,
april
26,
2009.
[9]
Schaeffler,
Welding
journal,
1947.