WELDING RESEARCH
SUPPLEMENT TO THE WELDING JOURNAL, MAY 1996
Sponsored by the American Welding Society and the Welding Research Council
Corrosion Behavior of Welded Stainless Steel
Corrosion resistance of stainless steel weld metal depends on an
understanding of the nature of the protective passive film
that gives the steel its "stainless" characteristic
BY T. G. G O O C H
ABSTRACT. The corrosion resistance of
stainless steels depends on a protective,
passive film being fornled at the steel surface on exposure to the service environment. The use of fusion welding for fabrication leads to local compositional
variations within the material, which
may significantly alter the stability of the
passive layer and hence the corrosion behavior. The paper reviews published
work and practical experience on such
effects of a weld thermal cycle.
Compositional heterogeneity and locally reduced passive film stability can
stem from three main causes. Attention is
given first to the consequences of alloy
element segregation during weld metal
solidification and to the formation of a fusion boundary unmixed zone. Element
partitioning as a result of solid-state
phase change is then considered, with
particular
reference
to
duplex
ferritic/austenitic materials. Finally, the
effects are described of precipitation,
both of carbide (or carbonitride) particles
and of intermetallic phases, with resultant development of alloy depleted regions. Throughout, recognition is made
of the different grades of stainless steel
employed by industry in terms of overall
composition and microstructure.
1 G. GOOCH is with T~/Vh Abin~4ton Hall,
Abington, Cambrid,~e, U.K. Paper presented
as the Comfort A. Adams Lecture at the AVVS
76th Annual Meetin,G April 3-7, 1995, Cleveland, Ohio.
Study remains necessary to ensure
that corrosion resistance of welded joints
is adequate for service in a range of
media. However, controlling factors
have in large part been defined for both
established and recently developed
steels, and quantitative recommendations on preferred welding procedures
can generally be given.
Introduction
There can be no doubt that the development of stainless steel represents one
of the major technical achievements of
the 20th century. The demonstration and
understanding that the addition of
chromium to steels endowed thenl with
the property of passivity led to the development of materials with remarkable
KEY WORDS
Stainless Steel
Austenitic
Martensitic
Ferritic
Duplex
Welding
Corrosion
corrosion resistance, coupled with excellent mechanical properties (Ref. 1).
Hence, "stainless steels" have reached a
dominant position for a very wide range
of applications. They give resistance to
aggressive media, provide an inert surface such that a product stream is not
contaminated, and can offer high
strength and toughness for service at extremes of temperature. The alloys are employed in chemical all([ process plants,
for power generation and other energy
industry equipment, for the food industry,
and for medical duties.
The fact that passivity can be induced
in an iron-based system is extrenlely fortunate, since, by appropriate adjustment
of elements other than chronlium, ferritic, martensitic or austenitic structures
can be produced singly or in any combination. The result is that alloys can be tailored to produce specific properties, depending on practical requirements and
costs of alternative materials.
Welding, especially by a fusion
process, plays an essential role in fabrication of stainless steel products because of
the economy and flexibility afforded to
design and manufacturing. A continuous
item is produced, facilitating direct transmission of stresses through the stru(ture,
and avoiding through-wall perforations
and potential leakage resultant on other
methods of joining. However, welding
necessarily involves a thermal cycle with
localized heating and cooling, and expansion and contraction. Regardless of
WELDING RESEARCH SUPPLEMENT I 135-s
!
1
I
I
0.50
m
30~ ~ 0
MO
'ZII
lie
~.-0.25
I
10
I
I
10z
10 ~
Currenf densify, p.A/cm z
I
10 ~
105
I
E
I
WI
]
Fig. I - - Potentiostatic polarization curves for 12% Cr (410), 17%
Cr (430) and 18% Cr (304) stainless steels in 1N H2SO 4 + 1N NaCI
at room temperature (Ref. 2).
WI
El
i
Z Effect of chromium: Cr.8%Ni steel- I
-0.1 ~
I
LI
01
-Ok
I
Effect of nickeh 18%Cr-Ni steel
-~ - o . 2 1 ~
al
iilUiii!iI
-OI,
~iEi!i!ill
-0.I
Effect of molybdenum: 18%Cr-lO%Ni steel I
-03
o Mo
-04
'
,
1
2
Currenf densih/,mA/cm z
0
:
3
I
Fig. 2 - - Effects o f varying Cr, N i and M e on active range o f stainless steels in 10% H2SO 4 at 2 0 ° C (Ref. 3).
11
|
:)
1.0
,
,
' r=j...~~-~/~
0.8
|
w
I
~s 0,6
:):?: ,/t0: I
J;:l
:
U,i; I
~ o,4
02
O,
0
I
÷lOOmVsce • •
I
20
I
I
I
40
50
80
Steel Cr content, af %
100
Fig. 3 - - Relationship between Cr content and Fe-Cr base metal. A
- - Cr cation fraction in surface film; B - - atomic fraction in underlying material: I M H2SO 4.
136-s
I MAY 1996
the grade of stainless
steel, and whether it be
wrought or cast, the base
material will be supplied
following careful processing to achieve optimum properties. The additional welding cycle
will differ appreciably
from that imposed during material manufacture, and can markedly
change the properties
around the joint. Specifically, this is the case
with corrosion resistance, and the present
paper is intended to review the effects of welding as a method of fabrication on the service
behavior of stainless
steels in aggressive
media.
It is generally considered that the development of passivity in
iron-chromium alloys
occurs virtually as a step
change as chromium
level
is increased:
below some 10% Cr, the
material remains active
in normal aqueous
media, while above
12% Cr, a passive film
develops and corrosion
rate is greatly decreased
(Ref. 1). While this is essentially true, the influence of the environment
must be clearly recognized in terms of redox
potential and content of
specific anions such as
chlorides (Ref. 2). Under
aggressive conditions, a
chromium level well in
excess of 12% will be
needed to maintain a
stable passive film - Fig. 1. There are, therefore, varying degrees of
"stainlessness" (Fig. 2)
(Ref. 3), and a major effect of welding is that the
associated metallurgical
changes may lead to
some parts of the material having lower levels of
chromium (or other elements which enhance
corrosion resistance, primarily molybdenum and
nitrogen) than elsewhere
on the metal surface.
This situation can stem from segregation
of alloy elements during solidification,
from element partitioning during phase
changes, or from precipitation of second
phase particles. These aspects are considered below.
Passivity of Stainless Steels
Passive Film Formation
The nature of the passive film formed
on stainless steels has been researched
for many years (Refs. 1,4, 5). Although
commonly, and conveniently, regarded
as an oxide layer, it is in fact considerably
more complex. Simple application of a
Pourbaix diagram, for example, would
not predict the observed fact of passivation in acid media on the basis of a thermodynamically stable oxide (or hydroxide) layer. The constitution of the passive
film has not been precisely defined, but
the development of x-ray photoelectron
spectroscopy (XPS) has enabled study to
be carried out on stainless steels surfaces
exposed to various environmental conditions. The approach is in principle more
direct than interpretation of electrochemical data or examination of film
stripped by aggressive etchants, although
the possibility of some change between
removal from the test environment and
XPS analysis must be recognized.
Accepting this limitation, Asami, e t a l .
(Ref. 4), working with a range of Fe-Cr alloys, found that the surface film formed
after exposure to 1M H2SO 4 in the passive range became chromium-rich immediately after the substrate chromium
content exceeded 12%, the minimum
level commonly associated with passivity and hence "stainless" steels-- Fig. 3.
Despite the high chromium content in
the film, the composition of the metal
surface corresponded to that of the bulk
alloy, indicating that the film resulted
from preferential dissolution of iron.
Chromium enrichment in the surface
film has been shown (Ref. 5) for a range
of potential, and in both chloride-free
and chloride-containing media, but only
under passive conditions. Below the passive/active transition, the surface film analyzed reflected the Cr and Fe contents of
the base metal. Hashimoto, e t a l . (Ref. 5),
concluded that the passive film was essentially a hydrated chromium oxy-hydroxide (CrOx(OH)3_2x-nH20), with a
small amount of hexavalent molybdenum if the substrate steel contained this
element. Molybdenum was found to be
beneficial primarily by suppressing active sites via formation of an oxy-hydroxide or molybdate, and this was further indicated by Halada, et al. (Ref. 6).
200 1[
I
I
,
,
b.t 2"
A
"
.
lime
,
,~,,
,'o
'
I
r
?°
0
i~.;l'z
sec
Fig. 4 - - Current transients on austenitic stainless steel in chloride media (Ref. 11). A - - Metastable pitting; B - - stable pitting.
The beneficial effect of nitrogen has
also been explored (Refs. 6, 7). Available
evidence has indicated that surface enrichment of nitrogen takes place at breaks
in the passive film, thereby blocking active dissolution. This view is consistent
with the observed synergistic effect between molybdenum and nitrogen (Ref. 6).
It has also been suggested that nitrogen
encourages iron dissolution, thus enhancing repassivation by the resultant
chromium enrichment (Ref. 8), and the
two mechanisms may be complementary.
ble hexavalent chromium ions. Nevertheless, the passive region is sufficiently extensive for the materials to be employed in
a wide variety of service media.
However, stainless steels are sensitive
to local passive film breakdown and attack, most notably by chloride ions (Ref.
1). This limitation is well recognized, and
has led to significant developments in
material composition to achieve resistance to chloride pitting (and crevice attack), exemplified by the increased Mo
and N levels in high-alloy "super-
austenitic" and "superduplex" grades
marketed over the last decade or so. Pitting occurs at potentials corresponding to
general passivity, becoming more pronounced at higher chloride levels and
applied potential. The formation of a pit
is preceded by electrochemical noise,
and, from measurement of small corrosion current transients (typically less than
l n A to over 101aA), the concept has
emerged of the passive layer being frequently penetrated by chlorides (or other
sufficiently mobile anions), but with re-
Passive Film Breakdown
Stainless steels display passivity only
over a certain range of potential (Ref. 1).
At low potentials, the passive film cannot
form and the material corrodes in the active state (Fig. 1), whereas under high potential conditions (in chloride-free media),
the passive film is oxidized to form solu-
Fig. 5 - - Segregation and pitting in fully austenitic weld metal
(125X).
i|O l
l
I
I w
! I
[.
ou
,
Fig. 6 - - Critical pitting temperature data from ferric chloride tests on stainless
steel base metal and weld metals (Ref. 17).
WELDING RESEARCH SUPPLEMENT I 137-s
pair rapidly taking place after dissolution
of a small volume of metal (Refs. 9-11).
As chloride activity or applied potential
is increased, the volume of metal dissolved is sufficient that local hydrolysis
and acidification prevent repassivation,
and a stable pit is formed which will then
propagate - - Fig. 4. The critical volume
of metal constituting the pit nucleus depends on the environmental conditions,
but has been estimated as a hemisphere
from about 10 nm to 11Jm diameter (Ref.
9). In this model, the initial passive film
rupture is considered to take place at
Table 1 - - EDX Analyses (wt-%) on GTA
Root Passes in $31254 Pipe Butt Joint Welds
Filler
Autogenous
ERNiCrMo-3
ERNiCrMo-4
Element Bulk
Mo
Cr
Ni
Mo
Cr
Ni
Mo
Cr
Ni
Dendrite interCore dendritic
6.1
4.2
20.2 19.3
17.6 17.6
7.6
5.8
21.4 20.7
47.2 47.7
11.3
9.8
17.3 17.2
40.7 43.1
6.9
20.5
19.2
10.5
20.9
43.6
14.4
18.1
39.2
I
Fig. 7 - - Dependence of stainless
steel weld metal
passivation on interclendritic Mo
segregation: 30%
H2SO4 with 0.1
g/L NH4CNS at
25°C (Ref. 15).
I
/
.!
~.
I
0
0.5
I
I
1.0
1.5
!
2.5
20
Molybdenum segmc~hon ratio
I
I
I
I
I
.~lid i fication mode
Pffmary [ Mixed
2.S "
Primary ousfenife
I
I
~~
¢3
1.0
Fig. 8 - - Relationship between element segregation,
partitioning and
bulk composition
and solidification
mode (Ref. 15).
1 3 8 - s l M A Y 1996
b-----o~
1-
0 l
0,4
,
0.5
ii
~,
8 o-o"--o---o
I, ~',.~1
0.6
0.7
0.8
' w
/P
#
cr~
I
I
,
0,9
1.0
1.1
~ I
1.6
weak points. Mattin associated rupture
events with inclusions in the steel surface
(Ref. 9), but in principle any local alloy
element depletion would be expected to
constitute a weak point and thus a pit initiation site.
Solidification
Alloy Element Segregation
A weld metal effectively constitutes a
chill casting and is subject to the normal
segregation of alloy elements during solidification of the molten weld pool (Fig.
5) (Refs. 12, 13) in the first instance, the
degree of segregation will depend upon
the equilibrium partition coefficient of
individual elements between the solid
and liquid states, and this can vary considerably depending on the element and
the primary
solidification
phase.
Chromium and molybdenum are rejected from the solid to a greater extent
with primary austenite solidification than
in weld metals freezing to ferrite (Ref.
14), whilst the reverse is true for nickel.
Nitrogen displays a somewhat unusual
behavior in that, thermodynamically, it is
more soluble in solid austenite than in
liquid steel, and thus tends to segregate
in the reverse sense to other elements
(Ref. 1 5).
Because of segregation, the corrosion
resistance of weld metal is inferior to that
of base metal of identical bulk composition (Refs. 15-17), In effect, the behavior
of the weld deposit is determined by the
minimum content of "passivating" elements at a point w i t h i n the solidified
weld metal, generally the dendrite centers since these are the first to solidify.
The consequences of segregation have
long been recognized, and consumable
specifications for common grades of
stainless steel are generally slightly overalloyed primarily in Cr and Mo, so that
even depleted regions retain sufficient
levels of these elements for satisfactory
corrosion resistance to be obtained.
Segregation is not commonly seen as
a practical problem with ferritic alloys.
The liquid/solid partition coefficients for
chromium and molybdenum approach
unity under solidification to ferrite, while
high-temperature diffusion in the bcc ferrite lattice is very rapid, so that some homogenization takes place in the solid
state during the weld cooling cycle. Nitrogen might segregate strongly, but this
element is usually held at a very low level
in ferritic stainless steels. In contrast, segregation becomes much more significant
in high-alloy superaustenitic grades. The
materials typically contain 20-25%Cr,
20-25%Ni, 6%Mo and 0.2%N, and seg-
regation of molybdenum can be sufficiently marked that some regions of the
weld metal contain only about 4% (Ref.
15), leading to an appreciable loss in corrosion resistance - - Table 1. This has
most frequently been highlighted with
reference to chloride pitting attack, and
indeed any reduction in passive film stability would be expected to have a pronounced effect on pitting behavior; given
passive film breakdown at alloy-depleted
regions and stable pit initiation, local hydrolysis and acidification will render
subsequent pit growth autocatalytic - Fig. 5. Over the last decade, it has become common practice to describe the
pitting resistance of stainless steels by an
empirical compositional parameter (Ref.
17), variously termed a "pitting index" or
"pitting resistance equivalent" (PRE),
such as (Ref. 18)
PRE = %Cr + 3.3%Cr + 16%N
I
800
I
I
I
I
n
600
'j"
I
I
400
200
I
0
0
/ I -
I
20
I
J
4O
I
I
50
Cr+ 3,SMo- 14N÷ (1.TMo, 23N)
-
8O
-
Fig. 9 - - Relationship between weld
metal pitting resistance and modified
pitting index (MPI)
(Re£ 15). Potentiodynamic passive
flint in 3% NaCI at
50°C.
(1)
A direct equivalence in pitting behavior can be drawn between the minimum
alloy level in the weld metal solidification structure and a base metal of similar
"lean" bulk composition - - Fig. 6. At the
same time, the adverse influence of segregation on passive film stability has
been demonstrated also in other media,
notably sulphuric acid (Ref. 15) in which
general rather than pitting corrosion will
occur
Fig.7.
Various factors for nitrogen in Equation 1 have been proposed. While Suutala and Kurkela used 13 (Ref. 17), 16 is
most commonly employed (Ref. 18), although a higher value of 30 has been recommended for high-alloy superduplex
and superaustenitic steels (Refs. 18, 19).
Recognizing that the PRE concept is empirical, Walker and Gooch showed that
the "optimum" compositional relationship, in fact, depends on the environmental conditions causing attack (Ref.
20).
D IQ! i
N
#
H
#
O
~
60
.4r/3Nz
L.j
0
40
Ar
Fig. 10 - - Effect o f
arc energy and
shielding gas on
FeCI 3 pitting resistance o f GTA welds
in $31254 steel
(Ref. 24).
20
I
0.5
o,24
I
I
20
Arc ene~jy, k J / ~
i
J
I
1.5
I
t
.
o z2
Effect of Composition
0.20
The degree of segregation is influenced by the total composition of the material, and tends to be more pronounced
in iron-based than in nickel-based systems (Refs. 15, 21 ). As illustrated in Fig.
8, segregation is likely to be most marked
in high-alloy austenitic steels containing
around 20-25%Ni, decreasing at higher
and lower Ni levels. Figure 9 shows the
correlation between weld metal pitting
resistance and a PRE parameter incorporating an empirical factor to recognize
the overall composition (Ref. 15), derived from studies of stainless steel solidification. From Figs. 8 and 9, it appears
0"18
.•E
0.16
0.11,
0.12
0.~
I
I
I
I
2
4
6
8
Nitrogen in A r shield,%
I
10
Fig. 11 - - Effect o f
nitrogen level in Ar
shielding gas on
22% Cr duplex
weld metal nitrogen
content (Ref. 25).
W E L D I N G RESEARCH SUPPLEMENT I 139-s
that segregation and the effect on passive
film stability is greatest in weld metals of
composition such that the primary solidification phase is austenite, but in which
the Cr and Mo rejection is sufficient for
the interdendritic regions to approach
final solidification to ferrite.
Experience with plant made from
316L-type steels and producing urea or
some organic acids, such as pteraphthalic acid, has indicated that preferential corrosion can occur in weld metals
containing a small amount of ferrite
(Refs. 22, 23). Consequently, it is normal
to specify low or zero ferrite contents,
with attendant sensitivity to solidification
cracking and microfissuring. Almost certainly, these media result in a metal/environment potential close to the active/passive transition, depending on the oxygen
level of the process stream. Passive film
!
c.I
z
CO
Fig. 12
Variation o f
c o m p o s i t i o n across an
-
-
interdendritic volume
element allowing complete mixing in the
liquid (Re£ 13).
Distance - - - ~ .
I
I
I
I
I
I
I
EE
I
f=O.18 sec ~
!
f~
•~
20
&
m
~3
Ni
b~
10
I
0
0
I
0.2
I
I
i
i
i
i
0.4
0.6
08
NOrl~lised d~-h~nce,r/R
A
,1
CO
kco
D m ~ n ~ ~
tO
Fig. 13 - - Calculated interdendritic microsegregation patterns for GTA welds in stainless steel (Ref. 28). A
ifying to primary austenite; B - - 23% Cr, 12% Ni, 65% Fe alloy, solidifying to primary ferrite.
1.0
,
,
0.9
0.8
0.7
,/
,
/I-~......--~
I
80
/"
50
I
O - -
~Itzl
6O
cont,
I
Base
70
~Z
O6
21% Cr, 14% Ni, 65% Fe alloy, solid-
•
QO
0.5
04
I
0310.~
--.~
Aziz
m ~
Wood
,
40
|
10"1
I
10
Dimensionless velocify
102
Fig. 14 - - Calculated variation of partition coefficient with velocity
using various models (Rei~ 13).
140-s
MAY1996
,/.
.........
I0
100
I000
Laser veloc#y,mm /min
10000
Fig. 15 - - Effect of laser travel speed on ferric chloride critical pitting
temperature of surface melted UNS $31254 steel (Ref. 31).
stability is therefore marginal, and it is
probable that the observed corrosion,
which tends to be initiated at ferriteaustenite interfaces, is due not to the
presence of ferrite per se, but rather to the
associated segregation and alloy depletion. Interfacial effects, such as residual
element segregation, may also be contributory. Comparable restriction on
weld metal ferrite content has seldom
been found necessary in other environments, and is apparently a consequence
of the specific cathodic corrosion reactions in the organic media concerned. It
seems that improved behavior could
equally well be obtained with weld metals of composition giving either fully
austenitic or primary ferritic solidification. In the latter regard, satisfactory service has been reported to TWl of materials such as 309LMo or 22% Cr duplex
steel which solidify initially to ferrite, although these are also somewhat more
highly alloyed than the 316L base metals
generally used.
A further point regarding weld metal
corrosion behavior is that the possibility
must be recognized of compositional
changes during welding. With flux
processes, factors such as element transfer efficiencies, alloying addition particle
sizing, etc., are well understood by consumable manufacturers, and the required
deposit composition can normally be obtained for a wide range of welding conditions. However, recent years have seen
increased use of nitrogen as an alloying
addition to enhance corrosion resistance,
and, when GTA welding high nitrogen
steels with a conventional inert gas
shield, this element tends to be lost from
the molten pool because of the relatively
low partial pressure in the surrounding atmosphere. In consequence, shielding
gases are finding favor that contain a few
percent nitrogen to counter the nitrogen
loss, and even to promote pickup by the
weld metal and improve corrosion behavior (Fig. 10) (Refs. 24, 25). The permissible nitrogen content in the shielding
gas depends on the solubility in the
molten pool, and hence on the composition, especially Cr and Mn levels. Excessive shielding gas nitrogen levels may
lead to porosity, and to evolution during
solidification, causing "spitting" and
degradation of the tungsten electrode. At
least for duplex ferritic-austenitic grades,
up to some 2% N 2 in the gas shield is normally acceptable for a 22%Cr steel, and
possibly 3% for a 25%Cr alloy - - Fig. 11.
Higher nitrogen levels may be tolerable
with the plasma arc process, used in keyhole mode (Ref. 26).
Effect of Welding
I
/
50
"
q
I
I
Fig. 16 - - Relationship between FeCI3 pitting
resistance o f welded joints
in UNS $31254 steel and
composition (Ref. 32).
A - - Weld metal bulk composition: O@ = GTA weld
metal; O~D = SMA weld
metal. H o l l o w symbols =
15% M o filler metal; solid
symbols = autogenous or
9% M o filler metal;
numbers = % dilution;
B - - minimum alloying
level in solidification structure: A = base metal; 17 =
autogenous weld metal; Ci =
GTA unmixed zone; ll =
SMA unmixed zone.
I
I
!I
I
!
50
~lIlr~ /60
4#
30
........................................................
45
i
I
I
40
50
60
Pl = % E r ÷ 93°/ot.fo+16°/oN
•
7T-).
-
J f ' L
%
~-
" "~'~
"'" ~.~x;-
:~.'.>'
/7).
- -.
_
.....
• <~
.
' >
.-,~_
V >r-';.-<<-.<,-d-.:~-,<I~J,4tD>.Q t \ ' ~ , . ~ : ~ '
, ~ ~ - z :<.',.,~.,
',
"\
~ ,-t
. . . .
..-J
/~
>~>.," .,/.,
.
•
,
~ t
,.1~I~!
~.. ' W ' " R ~ E V Z : z :
O01xm
.<t
~- '
"..'. . . .
,
Fig. 1 7 Unmixed
zone at weld interface
o f welds in 531254
steel: N i - 2 0 C r - 9 M o
filler metal. A - - SMA,
50)0
L
+t
(_b) ~ . ~
B - - center o f A, 320X;
Conditions
The residual segregation in a weld
c - - GTA root toe,
160X.
WELDING RESEARCH SUPPLEMENT I 141-s
allow for solid state diffusion after solidification, and found (Ref. 28) to give good
qualitative prediction of final segregation
gradients in stainless steels - - Fig. 13. A
problem arises, however, in that no direct
allowance is made for a time factor during the solidification process, with the resultant implication that the degree of segregation is unaffected by welding
procedure. From studies on castings, segregation is known to depend upon solidFig. 18 - - Pitting attack initiification rate because this influences eletared at U M Z in GTA weld in
$31254 steel: ERNiCrMo-3
ment redistribution apart from diffusion
filler metal, 25X.
and intermixing in liquid and solid
phases. Most alloy weld metals in which
the
effects of segregation are greatest soand
Vitek
(Ref.
13).
In
large
part,
the
metal at room temperature will depend
lidify primarily by dendritic growth, and
Scheil hypothesis has been followed, viz.
on element rejection during solidificavarious models for solute redistribution
tion, on the extent of intermixing in the
during such solidification have been deliquid phase ahead of the liquid/solid inC s = k C o ( 1 - f s ) k-1
(2)
rived (Ref. 13). These have been based
terface, and on the degree of diffusion and
mainly on the assumption of segregation
homogenization in the solid phase during
where Cs = composition of solid; Co = iniunder equilibrium conditions, and this is
cooling and any reheating by further weld
tial
alloy
composition;
fs
=
volume
fracunlikely to be the case over the range of
runs. The last few years have seen consolidification rates normally encomtion
of
solid;
k
=
partition
coefficient.
siderable advances in understanding the
passed by fusion welding.
This assumes that pool stirring by condependence of the solidification seAccordingly, models have been devection
and
electromagnetic
and
surface
quence in stainless steels on the bulk
veloped in which the segregation coeffitension forces is sufficient to give comcomposition (Ref. 27). This has been ascient varies between the equilibrium
plete mixing in the liquid phase - - Fig:
sociated with evaluation of element segvalue and unity as growth rate is in12.
The
approach
has
been
modified
to
regation behavior, as reviewed by David
creased from low to high velocities
(Fig.14) (Refs. 13, 29). On this basis, element segregation will be reduced and
weld metal corrosion resistance improved as solidification rate is increased.
This has been experimentally demonstrated by Nakao and Nishimoto (Ref. 30)
and by Woollin (Ref. 31), for the laser
process - - Fig. 15. With laser welding,
the solidification rate will be high, although rather lower than the travel speed
since solidification will occur epitaxially
from the base metal, with preferential
growth parallel to the maximum temperature gradient, and perpendicular to the
Fig. 19 - - EDX
scans across weld
travel direction. For travel speeds associinterface o f SMA
ated with more common arc welding
weld in UNS
procedures, growth rate will be relatively
531254 steel using
low, and will vary around the periphery
ENiCrMo-3 elecof the weld pool, so that a lesser effect of
Oisfonce,~m
trodes (Ref. 32).
solidification rate might be anticipated.
Nonetheless,
improvement in corrosion resistance has
been found at low
arc energy (Ref. 24),
with the rider that,
unfortunately, the
necessary arc energy
limits to achieve
practical benefit for
high-alloy austenitic
deposits may be so
low as to be unacceptable from the
productivity viewFig. 20 - - H A Z microstructures in duplex ferritic-austenitic stainless steels, 40X. A - - Predominantly ferritic; B - - austenp o i n t - - Fig. 10.
ite reformation.
142-s I MAY 1996
Unmixed Zone
Because of the problem of segregation
in high-alloy austenitic steels, it is normal
industrial practice to weld these materials using nonmatching nickel-based filler
metals such as ENiCrMo-3 (Ref. 19).
Fig. 21 - - CalcuThese were selected originally simply by
lated PRE values o f
virtue of their high molybdenum content,
ferrite and austenand although segregation still occurred,
ite in alloys with
even the dendrite cores had an alloy con25%Cr and
tent and corrosion resistance equivalent
4%Mo. Ni was
to that of the base metal (Table 1). The apvaried to keep a
constant ferrite
proach is now almost invariably emcontent. PRE = Cr
ployed for fabrication. Seam welded tube
+ 3.3%Mo +
may be welded autogenously (at, it may
16%N (Ref. 18)•
be noted, high speed), but this product
form will be given a high-temperature
G48 ferric chloride pitting test, attack
net nitrogen pickup in the bulk weld
homogenization treatment by the tubemetal (Fig. 10), and weldment pit initiagenerally takes place first on the unmixed
maker prior to delivery. Use of nontion then takes place preferentially at the
zone (Fig. 18), resulting in a weldment
matching consumables requires close atUMZ (Ref. 24).
critical pitting temperature (CPT) some
tention to welding procedure to obtain
The unmixed zone may be essentially
15-20°C (27°-36°F) below that of the
adequate filler metal addition (Ref. 32).
stagnant when molten, or may display rebase metal. In fact, it is the presence of
Dilution by base metal can be tolerated,
stricted laminar flow (Ref. 33). In either
the unmixed zone (UMZ) that renders
up to perhaps 60% (illustrated by the
case, its thickness is not uniform around
fairly high dilution levels acceptable in
hatched area in Fig. 16), but this still
the weld. This is partly because fluid flow
joints made with nickel-based filler metmeans that a root opening is preferable,
in the molten pool can lead to swirls of
als (Fig. 16), and further means that there
to ensure that filler metal addition takes
base metal becoming detached but not
is no practical advantage in employing
place during welding.
fully mixed into the bulk weld deposit,
even more highly alloyed filler metal
The application of "overalloyed" conbut in addition there will be an effect of
types for improved weldment corrosion
sumables has enabled satisfactory serthe drag induced by the surface adjacent
behavior (Ref. 32)
vice to be obtained from high-alloy
to the molten metal. This drag is less at
Available corrosion test data indicate
austenitic types in a range of media, but
the pool/gas interface than at the
that the UMZ is not as detrimental as
does not give weldment corrosion resisconventional autogenous weld metal.
pool/base metal junction, and in general
tance completely matching the base
This is indicated by the dotted line in
the stagnant zone tends to be thinner tometal. It has been known for many years
Fig.16 (Ref. 32). Solidification in the
ward the surface of the material as opthat, during arc welding, a stagnant zone
UMZ may be by cellular as opposed to
posed to the mid-depth of the weld pool
can exist at the weld interface (Fig. 17),
dendritic growth, but microanalyses
(Fig. 17A)(Ref. 32).
constituted by melted base metal that has
have not shown any marked difference in
This may not be the case actually at
not mixed into the bulk weld metal (Ref.
minimum Cr and Mo levels from segrethe toe of a root run, where a more pro33, 34). Following solidification, the regation in the UMZ (Fig. 19) and those in
nounced UMZ has been noted for GTA
bulk autogenous weld metal• The reason
gion displays segregation analogous to
than SMA welds, presumably reflecting a
for the disparate corrosion behavior has
that in an autogenous weld metal and is
local stagnant region - - Fig. 17C. Unnot been established, but it may be that
thus sensitive to localized attack in apderstanding of weld pool flow and the
nitrogen loss from the UMZ is less than
propriately aggressive media (Refs. 32,
relative effects of buoyancy, electromagfrom an autogenous weld bead. Cer34, 35). This effect was identified originetic and surface tension forces has intainly, pitting resistance of autogenous
nally in welds of ordinary 300 series
creased substantially (Reg. 37, 38), but
GTA
weld
metal
can
be
increased
by
use
stainless steels (Ref. 36), causing a minor
modeling remains a complex problem
of an Ar/5% N 2 shielding gas to promote
increase in the active/passive transition
potential, but it
becomes much
.~.'-..,~
,?
more significant
~,~
~
with high-alloy
grades to the extent that it is lim~ ' - ~ i
iting on corrosion
resistance
of
joints made with
overalloyed consumables. With
~
~
increasing environmental corrosivity, for exama ~ ' ~ ~ ~ ' 4 ' ~
0Y875
ple
higher
(a)
exposure temperFig. 22
Preterential weM metal attack in FeCI3. A - - Austenite, 70X; B - - ferrite, 250X.
atures in an ASTM
~o,
'
I
o
~
WELDING RESEARCH SUPPLEMENT I 143-s
[
~~.
10
'
z
'
I
.
, /~ • ~ •
,;
?-
Is
Ferri~ content, %
.
Fig. 23 - - Secondary austenite in superduplex weld metal,
1000X.
especially when filler metal additions are
made. In this respect, there will be some
disruption with GTA welding, as filler
metal is added, but the resultant stirring
will normally be low compared to gas
metal arc processes, whether metal transfer takes place by short circuiting or
under droplet spray conditions. The situation thus exists that, although GTA
welding may be preferred for control of
bead shape in root runs, this is the
process in which the unmixed zone is
likely to be most marked. In principle, the
width of the UMZ may be reduced via
welding conditions giving steep thermal
gradients (Refs. 33, 39). There may also
be benefit from use of electromagnetic
stirring, either from external induction
coils, or perhaps by high current, high
travel speed conditions to increase the
Lorentz force in the pool. Further work is
Fig. 24 - - Effect of ferrite-austenite balance on pitting resistance of 22%Cr0.12%N GTA weld metal (Ref. 49).
necessary, and this should recognize also
the possible influence of relative base
steel and weld metal melting points.
Solid State Transformation
Given the possible coexistence of ferrite, martensite and austenite, consideration must be given to the effects of elemental partitioning between the phases.
Since martensite formation is essentially
diffusionless, concern attaches primarily
to the ferrite and austenite phases, and
particularly to recently produced ferritic/
austenitic grades.
These alloys were developed primarily to avoid the problem of chloride stress
corrosion cracking associated with fully
austenitic steels. Originally, the composition was designed to achieve a twophase structure that could be accommodated
by
minimum
changes
to
existing
rolling
I
I
I
schedules to produce a
80
wrought product. This apJI Illlf
proach did not adequately
allow for subsequent trans.,drl I 1%1
.~rl
i i IJF
_/
6O
formation during a welding
~ l l l ~
cycle. The annealed base
P
metal might have a 50/50
.
phase balance, but weld
i ' ~
metals displayed primary
ferritic solidification, while
q
-./
the high-temperature heat20
f
affected zone (HAZ) trans---/
formed to ferrite close to
the solidus. Under the
0
I
I
I
I
fairly rapid cooling associ30
34
38
42
/'6
ated with welding, transPI= %Cr+~3(% Mo+OSx% W ~ 16x% N
formation to austenite was
suppressed, leading to a
Fig. 25 - - Relationship between duplex steel and weld metal predominantly ferritic weld
composition and critical pitting temperature in EeCI3 (Ref 46).
area structure (Fig. 20) with
-
I//
144-s
MAY 1996
reduction in corrosion resistance and
mechanical properties (Ref. 40). Hence,
it is now normal for welding consumables to be slightly overalloyed in nickel
to promote austenite reformation, while
nitrogen levels have been increased in
base and weld metals to promote transformation to austenite and obtain a more
equal ferrite/austenite balance, apart
from a beneficial effect on corrosion resistance (Refs. 8, 41 ).
The two phases in duplex steels can
differ appreciably in composition, Cr and
Mo partitioning to ferrite and nitrogen
and nickel to austenite. Accordingly, a
thermodynamic approach to alloy design
has been pursued by Berhardsson, so
that, by appropriate control ofCr, Mo and
N levels, the compositions of both phases
following conventional annealing are
likely to lead to similar corrosion resistance (Fig. 21) (Ref. 15). This does not,
however, avoid problems at welded
joints. Equilibrium partitioning of elements between the phases during welding is most unlikely to be achieved, and,
even under the range of welding conditions appropriate to normal industrial
practice, the compositions of ferrite and
austenite can differ markedly.
The ferrite-austenite transformation
during welding has been modeled by
Hertzman, et al. (Ref. 42), using a thermodynamic approach based on the Thermocalc system, with diffusional growth.
Viewed more simply, the mean diffusion
distance, x, of any element from random
walk theory is roughly (Ref. 43)
x = 2~/Dt
(3)
where D is the diffusion coefficient and t
is the time. Integrating the diffusion dis-
/200
I
1000
J
I
I
I
I
------
17%Cr: 0.05%C
-.,-~
.-.--.
.....
18%Cr: 8%Ni: 0 . 0 5 % C
18%Cr: 9%Ni: 0.06%C: 0 . 6 4 % N b
18%Cr: 8%Ni: 0.02%C
0,01
0.1
1,0
10
Ageing time,hrs
6OO
.,
,q
"4/
400
0
100
1000
Fig. 2 7 - - Effect of steel composition on sensitization behavior. Strauss testing, sensitized within bounding lines (Ref. 54).
lOOp, m
{
•0 T 2 5 8 2
I
Fig. 26 - - Intergranu/ar po/ythionic acid
cracking in sensitized 304 stainless steel,
125X.
tance (Ref. 43) over a cooling cycle, for
say a 1kJ/mm (25 kJ/in.) GTA weld run in
10-mm plate (0.4 in.), substitutional elements Cr and Ni would diffuse in ferrite
(Ref. 44) during cooling some 0.5 and 2
pm, respectively. As an interstitial element, nitrogen would diffuse roughly
50-100 pm. The resultant austenite laths
are perhaps 4 lam thick, and, in this example, austenite formation could be due
to substitutional element diffusion. Certainly, at higher arc energy and slower
cooling, appreciable redistribution of Cr
and Ni between the ferrite and austenite
would be expected. Conversely, with
lower heat input, austenite reformation
would be controlled by nitrogen movement, as was assumed by Hertzman (Ref.
42). Such differences in ferrite and
austenite composition have been observed over the range of cooling rate appropriate to normal welding practice
(Refs. 45, 46), and are illustrated in Table
2. Consequently, the relative corrosion
resistance of the two phases can vary
considerably, depending on the specific
steel and welding conditions entailed - Fig. 22.
The situation is further complicated in
that a single weld thermal cycle is most
unlikely to result in the equilibrium
austenite content being obtained, so that
additional transformation will occur on
deposition of subsequent weld runs. [n
part, this will involve growth of existing
transformed units, but it is possible also
for secondary austenite to be nucleated
and grow within the prior ferrite grains on
reheating (Fig. 23) (Ref. 47). The nitrogen
content of the ferrite matrix is fairly low
and, following diffusion of the nitrogen
Table 2 - - E D X Phase Analysis (wt-%) on SMA Welds in $31803 Steel Made with Overalloyed
Filler Metal under Varying Arc Energy Conditions
Element, wt-%
Weld region
Phase
Cr
Mo
Ni
As-deposited
root(a)
Reheated
rootTM
Reheated
root(b)
Base
metal
ferrite
austenite
ferrite
austenite
ferrite
austenite
ferrite
austenite
23.9
23.7
23.7
23.4
25.1
22.7
23.2
19.5
3.0
2.8
3.0
2.7
3.7
2.5
3.3
2.4
7.5
7.8
7.2
7.7
6.2
8.6
4.1
7.0
(a) Root and ill[ passes at 0.7 kJ/mm.
(b) Root at 0.5 kJ/mm, fill passes at 3.2 kJ/mm.
and substitution elements, the final PRE
of the secondary austenite will be below
the surrounding material (Table 3), leading to appreciable reduction in corrosion
resistance (Ref. 48).
From the metallurgical viewpoint, the
prime objective in formulating a welding
procedure for duplex steels should be the
attainment of a satisfactory phase balance in the weld locale. It is generally accepted that austenite contents between
about 30 and 70% are preferred for optimum overall properties. In particular,
high ferrite contents can lead to reduced
resistance to pitting corrosion (Fig. 24)
(Ref. 49), and to stresscorrosion, whether
in hot chloride media (Ref. 50) or under
sour, H2S conditions (Ref. 51). Such an
effect of high ferrite levels may be exacerbated by intragranular precipitation of
chromium nitrides, as a result of nitrogen
supersaturation in the ferrite at peak temperatures during welding (Ref. 52).
However, even with control of material composition and weld cooling rate to
obtain 50/50 phase balance, corrosion
resistance will not necessarily equal the
base metal - - Fig. 25. The adverse effects
of nonequilibrium partitioning tend to be
most marked in weld metal, and, since
they cannot yet be predicted reliably, the
use of overalloyed filler metals has been
propounded (Ref. 20). Commercial consumables for the established 22%Cr
grades are commonly slightly enriched in
Cr, Mo or N relative to base metal, while
even more highly alloyed 25%Cr superduplex filler metals may be used for
root runs exposed to the service environment. Experience has shown that analogous overalloying is problematic with
superduplex base metal, because the
weld metal may be unacceptably sensitive to precipitation of intermetallic
phases, which reduce corrosion resistance and toughness (Ref. 53).
The last decade has seen considerable
attention paid to the chloride pitting resistance of stainless steels, both
Table 3 - - EDX and WDX Analyses on
26Cr/9Ni/4Mo/0.26N Superduplex Weld
Metal (Ref. 47) (PRE = Cr + 3.3Mo + 16N)
Element, wt-%
Region
Cr
Mo
N
PRE
Primary austenite
26.6 3.3 0.52 45.8
Ferrite
27.4 4.0 0.07 41.7
Secondary austenite 24.3 3.4 0.24 39.4
WELDING RESEARCH SUPPLEMENT I 145-s
Table 4 - - PRE Relationships Derived from Polarization Tests on Duplex Weld Metals in
CO2-Saturated 3%NaCI: (FeCI3 data in brackets) (ref. 20)
Pitting Criterion
Specific PRE Equation
Pitting potential at 80°C
CPT at 700mVscE
(CPT in FeCI3
Cr + 1.9Mo + 0.55N - 0.27Ni
Cr + 2.9Mo + 21N - 0.25Ni
Cr + 2.1Mo + 40N - 1.2Ni)
austenitic and duplex grades, for marine
and oil and gas service. The PRE concept
of Equation 1 was developed originally
from laboratory pitting tests, especially in
ferric chloride. Clearly, from Tables 2 and
3, the PRE values of ferrite and austenite
can differ relatively, and the net pitting
resistance of a weld will depend on the
minimum PRE in any one region. At the
same time, studies on duplex steels have
shown that caution is necessary in applying the PRE approach directly to practice. As indicated in Table 4, the relationship between resistance to chloride
pitting and material composition depends critically on the environmental criterion adopted to describe corrosion behavior (Ref. 20).
Precipitation of Second Phases
Carbide Formation
Austenitic Steels
All grades of stainless steel are subject
to precipitation of a range of secondphase particles at intermediate temperatures. The effect is particularly associated
with austenitic alloys, in which "weld
decay" constitutes the classic case of the
corrosion resistance of a material being
adversely affected by fusion welding
(Ref. 1). Grain boundary precipitation of
chromium-rich M23C6 carbides takes
place at temperatures between about
500°-900°C (932°-1652°F), as inevitably experienced in part of the heataffected zone of a fusion weld. Because
the chromium diffusion rate is appreciably less than that of carbon, chromium is
removed from the surrounding matrix as
the particles form, leaving a depleted
area of lower passivity than the grain centers - - Fig. 26. Historically, the problem
was overcome (Refs. 1,54) by addition of
"stabilizing" elements, principally niobium or titanium to the material to form
stable carbides and getter the carbon
from the matrix, or by reduction in material carbon content via the use of lowcarbon ferro-chrome in s t e e l m a k i n g Fig. 27. With the introduction of AOD or
VOD processing in steelmaking, carbon
levels in commercial products have
fallen dramatically, typically to 0.05% or
below, so that the practical problem is
greatly reduced, and weld decay is now
rarely encountered in plant used in the
as-welded condition (Ref. 55). This is the
case even without going to extra-lowcarbon grades (i.e., less than 0.03% carbon), although postweld heat treatment
(PWHT) or service within the "sensitizing" temperature range can still induce
susceptibility to corrosion.
Increasing time in the sensitizing temperature range leads to greater width of
the depleted zone and lower local matrix
chromium content. The situation is illustrated by the unidirectional analysis of
Stawstr6m and Hillert (Ref. 56): assuming
a thin continuous carbide film,
I Kinetic parameters
diffusivities
Temperature (t)
Strain (t)
Input information
Alloy composition
Alloy condition
Initial DOS
Prior work
TM history
TA history
Strain history
I
Fig. 28 - - Flow diagram for prediction of sensitization in austenitic
steels (Re£ 59)
146-s I M A Y 1996
Ef
Effective
Cr
diffusivity
d
Nucleation
N
kinetics
Cr depletion
width
Minimum Cr
at interface
Thermodynamic parameters
Activity coefficients
Composite Cr
C solubility
Equilibrium constant
Carbide concentration
I
[
I
Degree of
sensitisation
I
IGSCC
susceptibility
~ - ( Cr~ - Cri I
m= 2~/Dtl ~
)
(4)
where m is the depleted zone width,
while Cr o, Cr i and Crp are the chromium
levels in the bulk material, at the carbide
interface, and necessary for passive film
stability, respectively. Since this study,
the development of depletion has been
well modeled, both under isothermal
conditions (Refs. 57, 58) and during a
weld thermal cycle (Fig. 28) (Refs. 59,
60). Using a short-term electrochemical
potentiokinetic reactivation (EPR) test,
Bruemmer was able to obtain good correlation between the degree of sensitization and material variables, composition,
heat treatment conditions, and thermomechanical history (Fig. 29) (Ref. 57). It
has been assumed that Crp in Equation 4
is about 13% (Refs. 56, 57), i.e., the
chromium content necessary to render
iron alloys capable of forming a passive
film. In fact, any reduction in chromium
below the matrix content will expand the
potential range over which the material
is in the active state (Fig. 2) (Ref. 3), and
will thus lead to some susceptibility to
preferential intercrystalline corrosion
(Ref. 61 ). Whether or not an attack occurs
then depends on the service redox potential relative to the grain boundary active range, i.e., the risk of weld decay depends on the specific environmental
conditions, as well as on the steel composition and welding procedure. A small
reduction in grain boundary chromium
content will be significant only in media
close to the active/passive transition,
whereas extensive carbide formation can
lead to intercrystalline attack over the entire potential range giving a stable passive film on the grain centers (Fig. 30)
(Ref. 61 ). Further to this thermodynamic
effect on passive film stability, the carbide particles also act to accelerate passive film breakdown (Fig. 31 ), increasing
the chance of attack under transient plant
conditions associated, for example, with
changes in oxygen activity.
Given that carbide precipitation is
time dependent, as illustrated by Equation 4, reduction in arc energy, and the
time spent in the sensitizing temperature
range, is advisable (Ref. 55). This is commonly recognized in welding procedures
for austenitic steels which stipulate maximum levels of arc energy and interpass
temperature. The effect of varying cooling time was modeled by Solomon (Ref.
60), who demonstrated the relationship
between increased cooling time or carbon level and enhanced susceptibility to
intercrystalline corrosion. The work further showed the environmental dependence of attack, and much more rapid
cooling was necessary to avoid attack in
the electrolytic oxalic acid test (ASTM
A262A) than in acid CuSO 4 (ASTM
A262E). Two caveats to this approach
must be added. First, in an actual weld,
concern attaches to the maximum sensitizing time experienced, and this will
occur at some distance from the fusion
boundary where the peak temperature is
high enough to enter the sensitizing
range, but not so high that the region is
in effect solution treated, i.e., sensitization takes place on heating as well as on
cooling. Second, the conjoint strain from
local expansion and contraction during
welding can act to accelerate the sensitisation process, as illustrated by Atteridge
and his coworkers (Fig. 32) (Ref. 59),
Although the general reduction in carbon content from AOD/VOD steelmaking has diminished the practical problem
of weld decay, experience in pipework
for light water reactor power generation
plants has indicated that the consequences of precipitation and sensitization can be greatly increased by residual
welding stresses, and intercrystalline
cracking has been experienced in a number of units. A considerable amount of
work has been carried out, therefore, to
address residual stress patterns at welds.
Techniques such as "heat sink" welding,
with or without local induction heating,
have been employed to induce compressive stresses on the inner wall of pipe,
thereby avoiding cracking (Ref. 62). At
the same time, the problem is especially
associated with the environmental conditions encountered in some light water
systems, and these measures would not
normally be employed for the typical
welded chemical plant, for example.
Again, in the context of light water reactor power plants, it is recognized that
initiation of grain boundary carbide precipitation in the HAZ can take place very
rapidly. Hence, although a controlled
weld thermal cycle may not induce sufficient precipitation to cause sensitization directly, the nucleated carbides
might grow in subsequent long-term service at lower temperatures, in this case
roughly 300°C (572°F). Indeed, from
evaluation of the activation energy as determined by slow strain rate tests, Povich
and Rao demonstrated that some measure of "low-temperature sensitization"
of welded Type 304 pipe could be anticipated after, say, 10 years exposure to
300°C (Ref. 63).
Most environmental conditions in
which stainless steels are used correspond to a moderately low redox potential. Under highly oxidizing conditions,
exemplified by a nitric acid plant, the effect of chromium depletion can be very
marked, while it is possible also for the
carbides to be dissolved. As a result,
these media are very searching and either
!
!
I
°
100
/6)
'
,3
•0
m
1.0
,4
L~. "'1"
=m|
°
°
0.1
•
0.01
001
01
Fig. 29 - - Measured
and predicted sensitization times for
304 and 316 steels.
Isothermal heat
treatment at
600°-700°C.
00
I
I
1
10
l'fe.asured hme, h
!
0.50
•
•
100
!
~ O25
./ /
0.15%C /
o
i
"-0.25
-0.50
l ) l
i
~ m ,, i
m
).Ol
l•
im
l
0.osO/oi
i
l
~,, qim ° i
i m, , ~ m
i
N
/0
0.1
1
Ageing hme,hrs
Fig. 30 - - Relationship between
potential necessary
Ior grain boundary
passivation and
minimum aging
time between 500 °
and 850°C, 18%Cr10%Ni austenitic
steel 20% H2SO4
with 0. l g/L
NH4CNS at room
temperature, general passive/active
transition for grain
centers below-0.48
VSHE (Re£ 61).
i
!
HAZ of electron beam weld (0.16kJ/mm
"me=
....
-----~---
0.6
H A Z of friction weld
H A Z of S M A weld (1 k J / m m )
H A Z of G M A weld (0.6 k J / m m )
Solution treated material
0.4,
,
02
,4##amntFl~e
peter,~l
0
-02
I
~I
-
" ~ .
ApparenfFlade
"
"" - 0 4
-0.6
-0.8
""
!
0
0.1
,I
•
1
Time,rain
#
10
100
Fi~¢. 31 - - Decay of
HAZ passive film
tbllowing
anodic
polarization
at
0.4VsH E tbr 10 rain,
20% HjSO 4 with
0.1 g/L NH4CNS,
18%Cr- 10%Ni0.15%C austenitic
steel.
W E L D I N G RESEARCH SUPPLEMENT[ 147-s
an extremely low
carbon content or
40
effective stabilizaThermal only
tion is essential for
Pass by pass strain
satisfactory service
•
Cumulative
strain
•
30
to be obtained. Further, boron added to
the steel for improved hot workability can stabilize
kj
the carbides to
higher
tempera10
tures, and this has
led to the developi 1 ~..~pv
. , , , , . ~ 1 0 ....
I
I
ment of low-boron
0
"nitric acid grade"
0
0.02
0.03
0.04
0.05
0.06
(NAG) alloys, in
Carbon contenf, vf %
which silicon and
phosphorus are also
Fig. 32 - - Predicted sensitization in a 304 steel pipe weld, assessed by controlled to miniEPR test. Effect of local strain (Ref. 59).
mize adverse effects
of grain boundary
segregation
(Ref.
64). Even with stabilized steels, there is
an advantage in reducing carbon for
nitric acid duty.
When making a
weld, the material
immediately adjacent to the fusion
boundary is effectively
solution
treated at high temperature, and some
/
carbon is thus released into solid soi~
,,
lution
to
form
0
M23C 6 on cooling.
Intercrystalline cor,0.2mm,
AF1035
rosion in this region
in a nitric acid plant
is known as "knifeFig. 33 - - Intercrystalline attack at weld in $40900 ferritic steel, 50X.
line attack"(Refs. 54,
65), and experience
has indicated that
niobium is more effective at gettering
carbon in the hightemperature HAZ
than titanium.
The effect of
carbide formation
depends upon the
density of precipitation at a single point,
and thus on the grain
size or total boundary area. With castings and weld metals
that contain some
ferrite, precipitation
occurs initially on
the ferrite/austenite
boundaries (Refs.
Fig. 34 - - Fracture face of intercrystalline attack in low-carbon 12% Cr 58, 66, 67). The initial development of
steel. Corrosion took place at ferrite/ferrite and ferrite/martensite
boundaries, leaving martensite units unattacked, 200X.
depletion depends
!
I
I
I
,/"
I
148-s I M A Y 1996
on the supply of carbon from the austenite and of chromium from the ferrite, but
overall the effects of precipitation are reduced, both by the high phase boundary
area available and also by the higher
chromium content in the ferrite phase.
Thus, materials containing appreciable
amounts of ferrite are resistant to sensitization (Ref. 66), although they are by no
means immune to the problem given prolonged exposure to elevated temperatures for postweld heat treatment or in
service (Ref. 67). As a consequence, weld
metals can contain higher carbon contents than base metal (Ref. 55), and this
permits the use of shielding gases containing some CO 2, provided that a sufficient ferrite level is maintained. For gas
metal arc welding, an addition of CO 2
may improve metal transfer characteristics across the arc, and reduce porosity
from shielding gas entrapment.
Martensitic and Ferritic Steels
Carbide formation and associated
chromium depletion as a result of welding is by no means confined to austenitic
grades but can develop also with martensitic and ferritic alloys (Ref. 54). With
martensitic steels, the problem is primarily associated with medium- to high-carbon contents (e.g., 0.2-0.5%) as may be
employed for surfacing applications. The
loss in corrosion resistance takes place
not only at prior austenite grain boundaries but also within the martensitic
structure of the matrix (Ref. 68). The effect is again controlled by chromium diffusion, and, since this is more rapid in a
martensitic or ferritic matrix than in fcc
austenite, depletion can occur in regions
heated to a peak temperature rather
lower than the austenitic case, typically
450°-500°C (842°-932°F). Low arc energy is preferred, but the situation is complicated by the need for fairly high preheat levels to avoid hydrogen cracking
on cooling to room temperature. In most
applications, martensitic stainless steels
will receive a postweld heat treatment to
temper the material. This will lead to
healing of any chromium-depleted areas,
and from the practical standpoint, the
problem is encountered only at weldments put into service in the as-welded
condition or tempered at very low temperatures to maintain high material hardness.
Because of the high chromium diffusion rate and low carbon solubility, sensitization develops rapidly in ferritic
steels (Refs. 54, 68, 69). The critical temperature range is similar to that for
austenitic alloys, but sensitization in ferritic grades requires only a fairly short exposure time even at temperatures below
say 500°C (Fig. 27). It can therefore be
difficult or impossible to avoid sensitization conlpletely by use of low arc: energy,
especially if multipass welding is involved. This is the case for virtually all
ferritic stainless steels, from lean alloy
11-12%Cr grades (Ref. 70) up to "superferritic" materials containing 25 30%Cr
(Ref. 71). In ferritic materials, nitrogen
has an analogous effect to carbon, although probably less severe on a wt-%
basis. Attempts have been made to avoid
the problem by steel processing procedures giving exceptionally low total
(C+N) levels, but in general practice, it
has been found more satisfactory to employ stabilization by titanium or niobium. It is remarked that nitrogen is not
recognized in standards such as UNS
$40900, yet preferential intercrystalline
attack has been observed in practice at
welded joints in such stainless steel complying with specification limits on carbon and titanium but having relatively
high nitrogen levels - - Fig. 33.
Time risk of sensitization of low-carbon
12% Cr alloys depends also on the weld
a r e a phase balance. Fully ferritic materials are appreciably more prone to the
problem than ferritic/martensitic alloys,
the adverse effect of precipitation being
increased by the local grain growth inherent in welding these steels. It has been
suggested that the formation of grain
boundary austenite, and hence martensite (Refs. 54, 69), during a weld thermal
cycle has a deleterious effect either by
leading to increased strain in the surrounding lattice, or because the austenite has a lower chromium content than
the matrix. The latter mechanism is not
valid, since, with normal welding conditions and cooling rates, EDX analysis at
TWl has shown the final martensite and
ferrite to have virtually the same
chromium level. Moreover, welds which
have suffered intergranular attack display
no evidence of preferential corrosion of
time martensite - - Fig. 34. Overall, there
is no doubt that high niartensite contents
are beneficial from the corrosion viewpoint, and apart from the high interphase
boundary area, this is probably associated with the higher solubility of carbon
and nitrogen in austenite than in ferrite,
together with the fact that, when nlartensite formation occurs, it takes place at a
temperature below the sensitizing range.
Further, if precipitation then occurs on
reheating in multipass welds, the "density" will be reduced by the high number
of nucleation sites within the martensite
and at the martensite/ferrite boundaries.
The extent of transformation to martensite at welds is therefore important, and it
must be recognized that, even though a
/00
80
50
4,0
20
0
5
10
15
20
"2F
Keltmt~user ferrite ~ d o r
12%Cr steel may be two phase (ferrite
and tempered nlartensite) in the annealed condition (Ref. 72), rapid cooling
during welding can lead to retention of
the ferrite formed at high temperature
(Fig. 35) (Ref. 73). Direct analogy can be
drawn with duplex ferrite/austenitic material (Ref. 40), and, in both cases, development of a predominantly ferritic structure at welds increases the risk of
intergranular corrosion in service.
The coarse grain structure normally
produced in ferritic stainless steel weld
metals can lead to unacceptable loss of
toughness, as well as sensitization, and
the materials are therefore sometimes
welded with austenitic consumables.
The activity of carbon and nitrogen will
depend upon the levels of chromiun] and
other elements, but if carbon and nitrogen contents in an austenitic weld metal
are appreciably higher than in the base
metal, these elements will diffuse into the
high-temperature HAZ (Ref. 74). This can
promote precipitation of chromium carbides or nitrides on cooling and thus loss
of corrosion resistance in this area, and
must be recognized by minimizing heat
input or ensuring that the base metal is
"overstabilized."
Intermetallic Phases
In the past, intermetallic formation in
conventional austenitic stainless steels
has been associated primarily with extended cycles at between perhaps
500 ° 900°C, during PWHT or service
(Ref. 54). More recently, it has become
evident that significant precipitation of
30
Fig. 3 5 - - Eftect o f
composition on
.structure o f 1 2 % C r
territic/martensitic
steels, x = base
m e t a l (Ret: 27);
h a t c h e d area = H A Z
data o b t a i n e d at
TWI. Ferrite l~ctor =
Cr + 6%Si+ *%Ti +
4%Mo + 2%AI +
4%Nb - 2%Mn
4 % N i - 4 0 ( C + N).
intermetallics can occur in the course of
welding, especially ill high-alloy
austenitic and duplex steels, leading to
loss of corrosion resistance. A number of
intermetallic types has been identified,
and, since they are highly alloyed and
presumably resistant to direct corrosion
in a wide range of media, it would seem
that the precipitation has an adverse effect by a depletion mechanism, analogous to that with carbide formation.
Intermetallic
precipitation
in
austenitic grades can be enhanced in
weld metal by alloy element segregation,
possibly contributing to the decreased
corrosion resistance of autogenous weld
metals (Ref. 32). Intermetallics can form
also when high-molybdenum nickelbased filler metals are used, but joint behavior remains controlled by the UMZ.
Grain boundary intermetallic precipitation has been further observed in the
HAZ of superaustenitic weldments (Ref.
75), although reduced corrosion resistance in this region is apparent only
when solidification conditions have
been such that the UMZ is m i n i m i z e d - Fig. 36. Nonetheless, the effect renders it
difficult to achieve weldment corrosion
properties matching those of the base
steel unless some form of PWHT is carried out for homogenization. Nitrogen
retards intermetallic formation (Refs. 21,
76), and the problem may be ameliorated
in new superaustenitic steels containing
very high levels of nitrogen (circa
0.4-0.5%) for improved resistance to pitting and crevice corrosion in chloride
media (Ref. 77).
The consequences of intermetallic
W E L D I N G RESEARCH SUPPLEMENT I 149-s
1
Fig. 40. Indeed there
is debate as to
whether any intermetallic formation at
welds in superduplex
steels is tolerable, yet
it is difficult to see
how it can be
avoided
entirely.
Certainly, the rapidity of intermetallic
formation (Fig. 37)
makes it problematic
to design welding
consumables for superduplex steels that
are overalloyed to a
significant degree to
Fig. 36 - - FIAZ attack at weld in $31254 steel, 5X.
mitigate the effects of
partitioning during
the
ferrite/austenite
formation are very much more severe in
transformation (Ref. 53). Accordingly,
duplex and especially superduplex
nonmatching nickel-based fillers are
grades (Ref. 78). The rapid diffusion in
being explored (Ref. 81), and, provided
ferrite means that precipitate nucleation
that adequate joint strength can be
and growth can be extremely rapid (Figs.
achieved, these may well prove of prac37 and 38), and the concomitant depletical advantage for superduplex steels.
tion in chromium and molybdenum is
The effect of intermetallic formation
sufficient to cause drastic loss in pitting
when
welding duplex/superduplex steels
resistance (Fig. 39), although with somehas received particular attention with rewhat less severe effect on general acid
spect to pitting attack in chloride media.
corrosion resistance (Ref. 79).
From data such as in Fig. 40, the need has
Intermetallic formation has been exbeen expressed for definition of a maxitensively studied in the context of weldmum tolerable intermetallic volume fracing superduplex stainless steels since it is
tion. This, however, is not the critical paseen as a major hindrance in obtaining
rameter. The form of Equation 4 is of
optimum joint properties. Its reliable
general
application to describe the deavoidance necessitates use of low arc enpleted zone round a precipitate particle.
ergy and very low interpass temperaBoth the width of a depleted region and
tures, which severely restricts overall
the
local reduction in alloy content will
joint completion rate (Ref. 46). Only a
increase with increasing time in the presmall amount of intermetallic phase can
cipitation range, and pitting resistance
diminish pitting resistance (Ref. 80) - will thus depend on
the maximum size of
an intermetallic par1
ticle, and not the vol• UNS $32304
|
ume fraction.
[] UNS $31803
UNS $32550
Intermetallic for[] Modified UNS $32550: 3.8%Mo, 0.26%N
mation in HAZs or
weld metals of su1000
perduplex
grades
may readily reduce
.,800
the CPT in a ferric
chloride test by some
S 600
20°C (36°F), equivalent to a drop in PRE
D
of about 7 units in the
I
adjacent
material
2O0
(Fig. 6). Typical compositions of inter0
metallic phases are
10 "100
1000
10000
given in Table 5
Time,see
(Refs. 82-85), and
enrichment of a
Fig. 3 7 - - Time-temperature transformation curves for duplex stainless sigma particle in Cr
steels (Ref. 78). Intermetallic formation, 550 ° - 1000°C; (z', 300°-550°C.
and Mo would re-
1
150-s I MAY 1996
Table 5 - - Reported Analyses (wt-%) for
Intermetallic Phases in Duplex Stainless
Steels (Refs. 80-82)
Phase
Cr
Mo
Ni
Sigma
30
34
30
27
28
20
26
6
7
10
12
18
39
25
5
4
4
5
3
3
4
Chi
R
quire diffusion from a radius roughly two
times that of the particle, to decrease the
PRE by 7 units, i.e., a reduction of 2.3%
Cr and 1.4% Mo in the ferrite. This assumes that sigma formation is controlled
by diffusion in ferrite and the ferrite composition. The fall in CPT in Fig. 40 is associated with intermetallic particles
some 0.5 to 1.0 IJm in diameter. The associated volume of alloy depleted material of 1-2 pm radius is certainly large
enough to constitute a stable pit nucleus
(Ref. 9) and the observed effect on pitting
resistance would be expected.
This represents an extreme case, since
such a particle size requires a fairly extended time in the precipitating temperature range. Diffusion of Mo and Cr (Ref.
44) over 1pm would involve several minutes exposure at, say, 800°C (1472°F),
equivalent to an arc energy of well over 2
kJ/mm (51 kJ/in.) on 12-ram (0.47-in.)
plate, given room temperature "preheat."
This arc energy would now be regarded
as excessive, following procedural development trials over the last few years. From
TTT curves such as in Fig. 37, intermetallic formation may be more rapid than indicated by bulk diffusion data, and it is
likely that growth is influenced by faster
phase boundary diffusion, which will
modify the elemental depletion pattern
and the ease of initiation of stable pitting.
Nevertheless, a particle size of 0.1-0.2
lam could be attained with welding conditions deemed entirely acceptable for
22% Cr alloys, the concomitant alloy depletion being sufficient to reduce the CPT
by perhaps 10°C (18°F). This particle
width is at the limit of resolution by optical microscopy, while, assuming continuous phase boundary precipitation and a
ferrite/austenite lamellar width in the subcritical HAZ of say 20-40 pm, the volume
fraction would be less than 0.5%. Clearly,
a weld procedure qualification specification requiring intermetallic phases to be
below, say, 1% would not offer a safeguard against loss of pitting resistance due
to intermetallic precipitation during
welding, and it is better to place reliance
on actual corrosion testing.
Discussion
The need to consider weldability is
well recognized by steelmakers and has
appreciably influenced alloy design in
recent years, for example in the development of NAG austenitic steels and of current production duplex alloys. The industrial importance of welding has
further led to significant study of the fundamental changes that take place during
a welding cycle, with regard to solidification, partitioning and precipitation.
Equally, factors contributing to passive
film stability and breakdown have been
increasingly researched, and in combination the two lines of investigation have
greatly clarified the practical effects of fusion welding on corrosion behavior of
stainless steels.
Perhaps inevitably, the greater part of
research directly aimed at corrosion
properties of weldments has been in response to industrial problems. This has
normally entailed identification of controlling metallurgical or process factors,
followed by trials to optimize the key
variables. This is well illustrated by the
extensive work in the last two decades on
the problem of sensitization of stainless
steels and nickel alloys for nuclear power
plants, which has led to a good understanding of the interplay between material, welding and environmental factors
and to demonstration of the effects of applied stress upon passive film breakdown
at chromium depleted areas (Ref. 62).
The work has heightened awareness of
the importance of the metal/environment
potential in service relative to a passive/active transition, and this concept is
being increasingly employed in other industrial regimes (Ref. 86),
Much of the recent work on corrosion
of welds in stainless steels has entailed
newly developed alloys, both superaustenitic and duplex grades. In the former context, the practical significance of
weld metal segregation and the fusion
boundary UMZ is now well recognized.
It remains to be seen if the latter can be
controlled by design of filler metal to obtain a preferred freezing temperature
range or by welding techniques giving
more rapid pool stirring. It may be also
that new high-nitrogen steels (circa
0.5%N) (Refs. 76, 77) in which segregation of alloy elements may be rather different (Ref. 15) will be less sensitive to the
problem. If the unmixed zone can be
negated, then, in high-alloy austenitic
steels at least, weldment corrosion resistance is likely to be limited by intermetallic formation. As a precipitation
process, this should be controllable by
use of suitable low arc energy and total
Fig. 38 - - SEM backscattered electron image
showing intermetallic formation in a superduplex
steel (Ref. 80).
thermal cycle, albeit at some cost in productivity. One area which does not seem
to have received attention is the consequence of the base metal solution treatment procedure. If the final anneal is set
at too low a temperature, residual intermetallic nuclei may remain, which could
act to form particles during a weld thermal cycle, analogous to the problem of
low-temperature sensitization of 300 series alloys.
Turning to the duplex grades, the advances in modeling alloy element partitioning (Refs. 42, 48) are extremely encouraging. With current welding
procedures, it would seem that the consequences of partitioning are most
marked in consequence of secondary
austenite formation (Ref. 47), and in
many cases, the net corrosion resistance
is dictated by the formation of this constituent. The general guidelines for welding are understood in that secondary
austenite is particularly likely to develop
when a low arc energy root pass is reheated by a high arc energy second run.
This situation can be avoided by appropriate design of the welding procedure.
However, the propensity for secondary
austenite formation depends also on the
bulk material composition, for example
the presence or absence of tungsten (Ref.
47), and further effort is required to clarify this effect. In superduplex steels especially, the susceptibility to intermetallic
formation is evident, and with present alloys, it must be accepted that low arc energy/interpass temperature conditions
are required, although again with loss of
productivity. A predictive model, as derived for sensitization of 300 series steels
(Refs. 56-60), would be valuable, but
would necessitate careful consideration
of the range of intermetallic types and
compositions formed in superduplex alloys. More detailed quantification of the
effects of welding is needed, since the
recommendations for avoiding secondary austenite and intermetallic
phases are not entirely compatible. Development of the equilibrium austenite
content in a weld run so that there is little tendency for secondary austenite requires observance of a minimum arc energy,
whereas
suppression
of
intermetallics will entail a restriction to
the usable heat input.
The increased use of duplex steels has
had a major effect on weldment testing
requirements. In the past, the principal
problem stemming from welding was
weld decay, and this was controlled by
Fig. 39 - - Chloride pitting due
to phase boundary intermetallic
formation in superduplex steel
(Re£ 80), 200X.
WELDING RESEARCH SUPPLEMFNTI 151-~
601
Fig. 40 - - Effect
of intermetallic
precipitation in
superduplex
stainless steels
on reduction in
FeCI ~pitting
resistance
(Re£ 80).
I
I
a
m
0
specifying that the base steel should pass
on intercrystalline corrosion test as in
ASTM A262, following a simple isothermal sensitizing heat treatment to simulate a welding operation. This approach
is not viable for avoidance of weld area
attack in duplex steels, since both high
and low heat input may be damaging,
while it does not allow for loss of weld
metal properties. Hence, it has now become the norm to include corrosion testing in weld procedure qualification, usually involving ferric chloride solution at
one or more temperatures. The tendency
has been to specify conservative acceptance criteria, and the need to pass such
tests in practice had stimulated much of
the basic work carried out on the welding behavior of duplex steels.
The philosophy of incorporating a
corrosion test in stainless steel weld procedure qualification requirements is
gaining wider acceptance. However, it
means that study is needed of the practical consequences of a welding cycle on
corrosion behavior, in terms of the relationship between short tests and longer
term operation. To illustrate the point,
while intermetallic precipitation in superduplex steels has an adverse effect in
chloride media at fairly high potential
(exemplified by the ferric chloride test
method), the critical size of the pit nucleus to develop stable growth may well
be higher in service media of lower oxidizing power, such as CO2-containing
brines. If this is the case, then intermetallic formation will be of less practical import than generally considered, providing of course that it does not
unacceptably reduce other service properties, notably toughness. Clearly, it must
be recognized that imposition of unrealistically demanding test procedures may
severely restrict the application of
welded stainless steels and incur appreciable economic penalty.
Concluding Remarks
A review has been carried out of a re-
1 5 2 - s l M A Y 1996
I
1
4
5
ln~lic.%
cent study on the effects of fusion welding on the corrosion resistance of stainless steels. Reference has been made to
ferritic, martensitic and austenitic alloys
and to dual-phase grades. In all cases,
concern attaches primarily to the stability of the passive film and the effects have
been addressed of 1) segregation during
solidification, 2) partitioning during
phase changes, and 3) precipitation of
second-phase particles. Advances have
stemmed from fundamental metallurgical and corrosion studies, and from practical procedural development trials, especially on recently developed alloys.
Service behavior will be determined
also by factors such as the degree of
residual oxide present on the joint on entering service (Ref. 46), and study remains necessary so that base metal corrosion resistance can be reliably
obtained at welded joints and unforeseen
service failures are avoided. Nonetheless, a good level of understanding has
been established, and in most cases
quantitative recommendations on welding procedures can be made to obtain
optimum corrosion resistance at fusion
weld joints.
Acknowledgments
The author thanks colleagues at TWl
and associates throughout the world for
their detailed study of stainless steels, for
their published papers and for many informative discussions. Their work is
gratefully acknowledged as the primary
input of this paper.
Reteren( es
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N.Y., Wiley.
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Electrochemical evaluation of resistance of
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Physical Metallurgy of Welding. London. Butter~:orths.
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Effect of composition on corrosion resistance
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$31254 high alloy austenitic stainless steel.
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]1. Woollin, P. 1995. laser beam surface
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34. Baeslack, W. A., Lippold, J. C., and
Savage, W. F. 1979. Unmixed zone formation
in austenitic stainless steel weldments. Welding Journal 58(6): 168-s to 176-s.
35. Flasche, L. H., and Ahluwalia, H. S.
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