Copper-Contamination Cracking
in the W e l d Heat-Affected Zone
Alloy systems susceptible to embrittlement by Cu-induced hot cracking
in the weld heat-affected zone are identified and classified
BY W. F. SAVAGE, E. P. NIPPES A N D M. C MUSHALA
ABSTRACT. Hot cracking in the weld
heat-affected zone has traditionally
been associated with microsegregation and liquation of low-melting,
second-phase particles w i t h i n the base
metal. Recently, hot cracking has been
reported at positions in the weld heataffected zone remote from the fusion
boundary in Haynes 188, a Co-base
superalloy. 1 The hot-cracking was induced by surface contamination of the
base metal by Cu abraded from the Cu
welding fixtures prior to welding.
Therefore, this investigation was undertaken to determine the mechanism
responsible for the Cu-induced hotcracking which cannot be explained
by contemporary theory.
A variety of Fe-, Ni-, and Co-base
alloys were tested with the Tigamajig
and the Varestraint weldability tests,
both with and without Cu surfacecontamination, to determine the structure sensitivity of the Cu-induced hotcracking. The Fe- and Co-base alloys
that have a FCC structure at the melting point of Cu, 1982.3 F (1083 C) were
found to be susceptible, whereas the
Ni-base superalloys were insensitive to
the Cu-contamination hot cracking. In
general, the use of unplated Cu weld
fixtures is likely to result in sporadic
Cu-contamination hot-cracking as a
consequence of Cu pickup on the
work surface while being loaded into
the fixture.
Introduction
The problem of hot-cracking has
long been recognized in a w i d e variety
of weldments. This problem has traditionally been associated w i t h microsegregation
in
the
heat-affected
zone.
Recently a curious form of sporadic,
though troublesome, cracking has
been observed by several fabricators of
FCC Co-base superalloys 1 and BCC Febase alloys which experience a transformation to FCC austenite at elevated
temperatures. The cracking appears to
be similar to hot-cracking except
that:
1. The cracking is usually observed
at positions in the heat-affected zone
slightly removed from the fusion
boundary.
2. The
completely
intergranular
crack morphology is typical of hotcracking; but since the cracks are not
contiguous w i t h the fusion boundary,
their morphology can not be explained
by any of the existing models for hotcracking.
3. The cracking propensity does not
follow the usual pattern, whereby
certain heats are consistently more
susceptible than others; rather the
incidence of cracking seems to be
almost statistically random, being extremely severe at isolated locations
and virtually non-existent elsewhere in
the same weldment.
Preliminary work indicates that the
hot-cracking results from Cu abraded
onto the surface. Verification of this
analysis w o u l d be significant because
a large proportion of the fixtures used
by welding fabricators is constructed
of either Cu or Cu-base alloys to
provide high thermal and electrical
conductivity.
The mechanism responsible for this
form of hot-cracking is believed to be
liquid-metal embrittlement involving
molten Cu. It is postulated that the
surface of the weldment is contaminated by Cu which is accidentally
picked up from the weld tooling by
abrasion. The heat of welding then
melts this minute amount of Cu
which, under the action of thermal
stresses, penetrates the grain boundaries by a liquid-metal-embrittlement
mechanism to cause the observed
cracking.
Heretofore, the role of liquid-metal
embrittlement has only been considered in systems involving large quantities of liquid metal in contact w i t h a
solid base metal. This investigation
will demonstrate that minute quantities of certain metals w i t h melting
temperatures below that of the base
metal may contaminate the surface of
a base metal and induce significant
cracking by liquid-metal embrittlement.
Throughout the remainder of this
paper this form of hot-cracking shall
be referred to as copper-contamination cracking. Accordingly, the symbol
"CCC" will be used to refer to coppercontamination cracking, to minimize
the repetition of this lengthy phrase.
The first recognition that Cu could
be responsible for hot shortness was
associated w i t h the production of hotrolled sheet of low-C and mild steels.
The majority of such hot-shortness
investigations reached t w o similar
conclusions:
1. There is no penetration of the
grain boundaries in the base metal at
temperatures below the melting point
W. F. SAVAGE is Professor of Metallurgical
of Cu; the significance of the fact that
Engineering and Director of Welding
liquid Cu is required for cracking was
Research, and E. F. NIPPES is Professor of
finally
realized by Cox and W i n n 2 w h o
Metallurgical Engineering, Rensselaer Polymentioned
that the mechanism of
technic Institute, Troy, N. Y.; M. C.
embrittlement might be that proposed
MUSHALA, former Graduate Assistant at
RPI, is now a Captain in the U.S. Air Force, by Eborall and Gregory'—namely, liqRandolph Air Force Base, San Antonio,
uid-metal embrittlement.
Texas.
2. It is also widely accepted that the
W E L D I N G RESEARCH S U P P L E M E N T I 145-s
surface of the material must be
stressed in tension above the yield
stress if there is to be measurable
grain-boundary penetration by the
liquid Cu.
Smith predicted that the penetration
by a liquid phase w o u l d be greater as
the dihedral angle between the solid
and liquid phase decreased. 4 This
concept further explains why the
attack by liquid Cu should be so
severe. According to Nicholson and
Murray, 8 the dihedral angle between
Cu and mild steel is a m i n i m u m at the
melting point of Cu as shown in Fig. 1.
Therefore in mild steel, at a temperature only slightly above the melting
point of Cu, the amount of undissolved Cu is a maximum and the dihedral angle between the liquid Cu and
an austenite grain boundary is a minimum. These observations can account
for the many experimental observations which have indicated the penetration by liquid Cu into mild steel to
be a maximum at a temperature slightly above the melting point of Cu.
Smith 4 theorized that the dihedral
angle could be predicted from the
balance of surface free energy. If the
surface free energy of the austeniteaustenite (y-y) grain boundary is
denoted by cr Ty and the surface free
energy of the austenite-liquid (y-L)
phase boundary is denoted by o"v,., as
shown in Fig. 2, then the dihedral
angle can be computed by the f o l l o w ing relationship:
CTyy = 2cr yl . cos (6/2)
This dihedral angle can have any
value from 0 to 180 deg, depending on
the value of cr Yy /o~ YI ,. If the dihedral
angle is 0 deg, annealing above the
melting point of the second phase will
result in continuous films of the liquid
second phase at the grain boundaries
and in disintegration of the alloy along
these grain boundaries. Penetration
along the grain boundaries is still
possible for dihedral angles greater
than 0 but less than 60 deg. In such
600.
r
40
TEMPERATURE (°F)
I800
2000
2200
r
SH0WING ESTIMATED
ACCURACY OF * 5 %
Ixl
^30
900
IOOO
I IOO
I200
TEMPERATURE (°C)
Fig. 1—Dihedral angle-temperature
relationship for the copper-mild
steel system
cases, the liquid w o u l d penetrate
along the grain edge until the
predicted equilibrium dihedral angle is
attained.
According to Smith, capillarity is the
mechanism of liquid-metal transport
along the grain boundary. This penetration is accelerated by stress and it is
possible that a tensile stress could
cause continued penetration and subsequent material failure for dihedral
angles between 0 and 60 deg.
W i t h a knowledge of the role of
surface tension in the penetration of
grain boundaries by liquid metals,
many investigators sought to predict
the embrittlement of alloy systems
from a knowledge of the surface
tensions involved. Van Vlack 6 reported
the interfacial energy for an austeniteaustenite grain boundary in Fe to be
850 ergs/cm 2 and that of a liquid Cuaustenite boundary to be 430 ergs/cm 2
at 2021 F (1105 C). If it is assumed that
the surface tension is numerically
equal to the interfacial surface energy,
then these data predict a ratio of the
surface energies, cr s l - /cr s s , to be 0.506.
As shown in Fig. 3, this corresponds to
a dihedral angle of approximately 18
B IS THE DIHEDRAL ANGLE
SCALE A
deg. Therefore, one w o u l d expect
rapid penetration of the austenite
grain boundaries in Fe by liquid Cu at
2021 F (1105 C). This is consistent w i t h
the previously mentioned investigations of hot shortness in mild steel.
However, subsequent attempts to
predict the incidence of hot shortness
from the surface tensions involved
have met with varying degrees of
success. From a review of surfacetension literature, it can be seen that
the surface tension is dependent upon
composition,
temperature,
grain
boundaries, surface condition, and
trace element additions. As a result,
different investigators have quoted
values for surface tensions in the same
system, under similar test conditions,
which may vary as much as 20%.
Obviously, such variations make the
prediction of the dihedral angle by
Smith's formulation at best uncertain.
As can be seen in Fig. 3, a 10% error in a
surface-tension measurement w o u l d
change the ratio of o" S I /o- s s from 0.50
to 0.55, w o u l d change the calculated
dihedral angle from 0 to 45 deg.
The real crux of the problem lies in
the critical assumption that the surface
tension is equal to the interfacial energy (surface free energy). Michaels 7 has
shown that, because of the anisotropy
of the crystal, the surface free energy
and the surface tension are rarely, if
ever, equal. Hence, the assumption
that the surface tension is numerically
equal to the surface energy can introduce a significant error in the aforementioned computations.
Bredzs and Schwartzbart 8 have
shown that a volume diffusion mechanism for the penetration of Cu-brazing
filler metals into Fe-base alloys w o u l d
be entirely too slow to account for the
observed penetration. They calculated
that the penetration by volume diffusion w o u l d only be 0.0008 in. (0.02
mm) after four months at 1990 F (1088
C). Since the rate of grain-boundary
penetration in an Fe-base by solid Cu
SCALE B
X
°rr^
r
Fig. 2—Equilibrium between a grain boundary
and two equal interphase boundaries
Fig. 3 (right)—Ratio of interphase boundary tension and grain
boundary tension as a function of dihedral angle of second
phase
146-s | M A Y 1978
0
30
60
90
I20
I50
ISO
8 = Angle between faces of grain of second phase
Table 1—Summary of the Alloys and Their Nominal Compositions Used in This Investigation
Material
AISI Type 304
AISI Type 316
AISI Type 321
AISI Type 347
AISI Type 410
AISI Type 430
AISI Type 446
A-286
Inconel 750
Inconel 718
Rene 41
Waspaloy
Hastelloy-X
L-605
Haynes 188
HY-80
AISI 4130
AISI 4340
Autobody Steel
6061 T-6
Fe
.067
.15max
.08max
.12max
.05-.15
.12max
.35max
.05max
.05
.08
.05-.12
.05-.1
.09
.10
.10
.18
.28-.33
.38-43
.055
—
BAL
BAL
BAL
BAL
BAL
BAL
BAL
BAL
6.69
18.30
5.0max
2.0m ax
18.05
1.4
BAL
BAL
BAL
BAL
.33
may be 10,000 times faster than
volume
diffusion,
Bredzs
and
Schwartzbart have shown that any
diffusion process involving solid Cu
could not possibly account for the
observed penetration of Cu into mild
steel. They have concluded that the
grain-boundary penetration must have
resulted from
liquid-metal
grainboundary penetration. They reported
that the Cu-penetration depth was
dependent upon temperature, austenite grain size, composition of the steel,
stress distribution in the steel, and
composition of the brazing alloy.
Other investigations of brazing and
soldering have also demonstrated the
following:
1. The liquid Cu or Cu-alloy penetrates the austenite grain boundaries in
Fe and, by pinning those grain b o u n d aries, prohibits grain growth as long as
the liquid metal is present.
2. The extent of penetration by
liquid Cu increases with increase in:
the C content of the base metal, the
hardness of the base metal, and the
grain size of the base metal.
3. For liquid metals to penetrate
grain boundaries, the liquid metal on
the surface of the base metal must
contact an area w h i c h has experienced
some localized plastic deformation.
4. Penetration is enhanced by wetting of the base-metal substrate by the
liquid metal.
Borland 9 was the first to suggest that
liquid Cu could cause cracks in the
weld heat-affected zone. This was
predicted by his "Generalized Theory
of Super-Solidus Cracking" w h i c h is
based, in part, on the dihedral angle
dependence of liquid-metal penetration.
Asnis1" has shown that liquid Cu on
the work surface can produce significant heat-affected zone cracking.
However, the most comprehensive
Cr
18.3
16-18
17-19
18-1914
11'/2-13'/2
14-18
23-27
15.0
15.18
18.93
18-20
18-21
21.8
20.0
21.82
1.68
.8-1.1
V-.9
.01
.22
Ni
8.65
12-14
8-11
10-14
.5max
-
26.0
73.10
BAL
BAL
BAL
BAL
10.0
21.87
2.99
-
1.65-2
.05
—
Other
Co
.11
10-12
12-15
1.58
BAL
BAL
—
treatment to date has been that by
Matthews 11 . He investigated the infiltration of the weld heat-affected zone
by dissimilar weld metals. He has
shown that Cu-rich weld deposits can
produce
significant
heat-affectedzone cracking when the base metal is
either FCC or experiences an allotropic
transformation to an austenitic phase
during welding. His experiments utilized a Cu-Ni alloy which was introduced as a filler metal during a G M A W
process.
Such a system can be envisioned as a
reservoir of a liquid Cu-rich alloy (the
weld pool) in contact w i t h a base
metal which is stressed in tension from
thermal stresses normally associated
with welding. Therefore, because his
system corresponded to the classic
model for liquid-metal embrittlement,
Matthews
hypothesized
that
the
mechanism for the observed hotcracking was liquid-metal embrittlement. The results of his investigation
showed that the hot-cracking appeared to obey most of the rules
quoted as governing the incidence of
liquid-metal embrittlement. In particular, the penetration was more severe
when the liquid filler metal and the
base metal exhibited low mutual solubility and did not form intermetallic
compounds w i t h one another.
Following Matthews' work, Stanton 1 2 has shown that Cu deposited as a
weld overlay could induce significant
heat-affected-zone cracking in a variety of steels, including AISI 4340 steel.
His results and conclusions closely
paralleled those of Matthews 11 .
Matthews 1 was the first to report the
embrittlement
of
the
weld-heataffected zone by Cu w h i c h had been
abraded onto the surface of a Co-base
superalloy prior to welding. Furthermore, he demonstrated that Cu concentrations in solid solution (up to
1.68Mn, .51 Si
2.0Mn max, 1.5-2.5Mo
Ti = 5C
2.0Mn max, Cb + Ta = 10C, 1.5max
1.0Mn max, .5Mo max
l.OMn max
-
1.35Mn, 1.25MO
.58Mn, 2.36Ti, .68AI, .96Cb + Ta
3.02Mo, .62AI, .88Ti, 5.16Cb + Ta
1.5AI, 3.15Ti, 9-10mo
1.25AI, 3.0Ti, 3.5-5.0MO
8.65MO
15.0W
13.83W
.3Mn, .41Mo
4-.6Mn, .2-.3MO
.6-.8Mn, .15-.25MO
.33Mn, .01 M o
.2Cu, 52Si, .88Mg, .032Mn, AI-BAL
0.96% in L-605) do not increase the
susceptibility to hot-cracking—that is
to say, Cu abraded on the surface can
induce hot-cracking but Cu in solid
solution should not be a factor in this
type of hot cracking. Matthews suggests that liquid-metal embrittlement
may be the mechanism responsible for
the Cu-contamination-induced hotcracking.
Liquid-metal
embrittlement
has
been described as the reduction of
fracture stress and loss of ductility
caused by liquid metal in contact w i t h
the surface of a solid metal. 13 For the
purposes of this investigation, two
forms of liquid-metal embrittlement
are applicable:
1. Near instantaneous failure when
a liquid metal is applied to a solid
metal subject to a tensile stress.
2. Grain-boundary penetration of
the solid by the liquid such that the
structure eventually disintegrates.
The incidence of liquid-metal embrittlement is restricted to specific
combinations of liquid and solid
metals. Because of the similarity to
stress-corrosion cracking, it was first
thought that liquid-metal embrittlement arose from basically the same
mechanism. However, there are several experimental observations which
indicate that different mechanisms are
involved. 13 Specifically:
1. Several pure metals (Cu, Sn, Zn,
and Al) are embrittled by Hg, but do
not appear to crack under conventional stress-corrosion tests.
2. Liquid-metal embrittlement can
occur nearly instantaneously following
contact with the liquid metal while
stress-corrosion cracking invariably exhibits a finite incubation time.
Phenomenologically, three conditions must be satisfied, for liquidmetal embrittlement to occur between
a specific solid-liquid metal couple.
W E L D I N G RESEARCH SUPPLEME NT I 147-s
The three criteria, w h i c h have been
derived from experimental observation, are the following: 1 3
1. There exists a low mutual solubility between the liquid and solid
metals.
2. There exists no
intermetallic
compound formations between the
solid-liquid metal couple.
3. There exists a barrier to plastic
flow in the base metal, w h i c h is also in
contact with the liquid.
Object
A peculiar form of heat-affectedzone hot-cracking has been attributed
to the presence of minute amounts of
metallic Cu on the surface of the w e l d ment. In this case, rubbing contact
between Cu weld tooling and the
material being welded transfers sufficient Cu to the surface of the work
pieces to cause cracking to occur
when welding materials sensitive to
this form of cracking.
To date, no quantitative investigation of this CCC phenomenon has
been reported in the literature. Thus, it
is not certain what type of alloy
systems are sensitive to this form of
cracking, nor is the amount of Cu
required k n o w n . Therefore, it seemed
appropriate first to report what alloy
systems are susceptible to this phenomenon and then in a later paper
report the responsible mechanism.
Materials and Apparatus
Materials
Table 1 lists the nominal composition of the alloys studied and the
specific tests used to evaluate the
sensitivity of each alloy to CCC. First,
the alloys can be classified as either
transformable or
nontransformable
within the temperature range of room
temperature to slightly above the
melting point of Cu (1982.3 F, 1083 C).
Attention is drawn first to the n o n transformable alloys w h i c h can be
subdivided as follows:
1. Fe-base stainless steels w h i c h
have a BCC structure throughout the
temperature range of interest—Types
430 and 446 stainless steels were
chosen as representative of this subdivision. These alloys have considerable differences in the Cr and Ni
concentrations used to maintain their
BCC structure.
2. Ni-Base superalloys which are
precipitation hardened by a Ni:1 (Al,
Ti), y', precipitate and have a FCC
structure throughout the temperature
range of interest—Inconel 750, Inconel
718, Rene 41, Waspaloy, and HastelloyX were chosen as representative of this
subdivision.
3. Co-base superalloys which are
precipitation hardened w i t h various
carbide phases and have a FCC structure throughout the temperature range
of interest—Haynes 188 and L-605 were
chosen as representative of this subdivision.
4. An Al alloy, 6061 T-6, was
selected as a representative example of
an Al structural material w h i c h has a
FCC structure.
Secondly, note the transformable
alloys w h i c h can be subdivided as
follows:
1. Fe-base stainless steels having a
FCC austenitic structure from room
temperature to well above the melting
point of Cu—the alloys selected represent normal austenitic stainless steels
(Types 304 and 316), and the Ti- and
Cb-stabilized austenitic steels (Types
321 and 347, respectively). However,
these alloys do experience at least
partial transformation to BCC ferrite at
temperatures near the solidus.
2. Fe-base
martensitic
stainless
steels which have a FCC structure at
the melting point of Cu and transform
to BCC ferrite and/or martensite on
cooling—Type 410 stainless steel was
chosen as representative of this subdivision.
3. Fe-base
precipitation-hardened
stainless steels which have a FCC
structure at the melting point of Cu
and are precipitation hardened by
intermetallic compounds containing
Ti which dissociate and dissolve in the
matrix at the melting point of Cu;
Table 2—Composition of Electrolytic
Copper Plating Bath
Fig. 4—Photograph of the Tigamajig testing device showing: A—GTA torch, B—specimen,
C— constant-radius die block, D-clamping fixtures, E—stop blocks, and F—loading ram
148-s | M A Y 1978
CuSO.-SH.O
H,SO, cone.
H,0
1355 gm
268 cc
5680 cc
A - 2 8 6 w a s c h o s e n as r e p r e s e n t a t i v e o f
this subdivision.
4. P l a i n - C a n d l o w - a l l o y c o n s t r u c t u ral steels w h i c h h a v e a FCC s t r u c t u r e at
the melting p o i n t of Cu and transform
t o a BCC o r a BCT s t r u c t u r e o n c o o l i n g — t h e s e a l l o y s are r e p r e s e n t e d b y a
l o w - C , p l a i n - C steel ( a u t o b o d y s t o c k ) ,
a
quench-and-tempered
low-alloy
steel ( H Y - 8 0 ) , a n d t w o
medium-C,
h i g h - a l l o y steels (AISI 4130 a n d AISI
4340).
Table 3—Summary of the Standard Welding Conditions for the Tigamajig tests
Electrode extension
Arc length
Electrode
Material
T h e T i g a m a j i g test w a s d e v e l o p e d at
RPI 1S t o c o m p l e m e n t t h e V a r e s t r a i n t
testing device by p r o v i d i n g a capability for testing sheet metal specimens.
A l t h o u g h t h e T i g a m a j i g test is s u b s t a n tially different in design f r o m the
V a r e s t r a i n t test, its p r i m a r y f u n c t i o n is
still t o e v a l u a t e t h e w e l d a b i l i t y o f a
structural material.
Figure 4 is a p h o t o g r a p h o f t h e a p p a ratus. T h e G T A t o r c h (A) is u s e d t o
p r o d u c e an arc s p o t w e l d in t h e c e n t e r
of t h e s p e c i m e n (B). A f t e r e s t a b l i s h i n g
a d y n a m i c t h e r m a l e q u i l i b r i u m , t h e arc
is i n t e r r u p t e d a n d t h e l o a d i n g r a m is
accelerated u p w a r d to force the specim e n to c o n f o r m to the radius die
b l o c k (C). This causes a r e p r o d u c i b l e
t e n s i l e s t r a i n , w h i c h is i n v e r s e l y p r o p o r t i o n a l t o t h e radius of t h e die
b l o c k , to b e a p p l i e d o n t h e t o p s u r f a c e
of t h e w e l d e d s p e c i m e n . T h e c r a c k i n g
p r o p e n s i t y o f t h e test s p e c i m e n can
t h e n be r e l a t e d t o t h e m a g n i t u d e o f
t h e strain necessary t o c a u s e c r a c k ing.
Procedure
The experimental investigations e m ployed to determine the mechanism
responsible for the CCC of the w e l d
heat-affected zone were designed on
the assumption that liquid-metal e m brittlement was responsible. T w e n t y
Arc time, s
Inconel 750
Type 446
Rene 41
Type 410
Inconel 718
A-286
L-605
Hastelloy-X
Waspaloy
Type 430
Type 304
Type 316
Type 347
6061 T-6<"
Description of the Varestraint Testing
Device
Description of the Tigamajig Testing
Device
Vs in. (3.2 mm)
Vs in. (3.2 mm) W-2% T h 0 2 ,
90 deg incl. angle
Ar, 60 cfh
30 s
12 V
40 A
None
3-phase rectifier
Shielding gas
Pre-purge time
Arc voltage
Arc current
Post arc delay time
Power supply
The a b o v e alloys represent a w i d e
v a r i e t y o f Fe-, N i - , a n d C o - b a s e a l l o y s
and represent a cross-section of t h e
s t r u c t u r a l a l l o y s in c o m m o n
usage
today. Thus by studying this g r o u p of
m a t e r i a l s it w a s p o s s i b l e t o d e t e r m i n e
t h e alloys m o s t s u s c e p t i b l e t o t h i s
f o r m of w e l d h e a t - a f f e c t e d - z o n e e m brittlement.
T h e V a r e s t r a i n t test w a s d e v e l o p e d
at RPI t o e v a l u a t e t h e w e l d a b i l i t y o f
structural alloys. A c o m p l e t e descript i o n of t h e test has b e e n p u b l i s h e d
p r e v i o u s l y 1 4 . T h e V a r e s t r a i n t test w a s
used in t h i s i n v e s t i g a t i o n t o e v a l u a t e
t h e s u s c e p t i b i l i t y o f H Y - 8 0 , AISI 4130,
a n d AISI 4340 steel t o
liquid-Cu
embrittlement.
VA in. (6.4 mm)
2.0
6.0
6.0
6.0
6.0
6.0
8.0
12.0
24.0
24.0
24.0
24.0
24.0
5.0
"Note: 12 V, 70 A.
materials representing a w i d e variety
of a l l o y a n d s t r u c t u r e c l a s s i f i c a t i o n s
w e r e subjected to either Tigamajig or
Varestraint testing t o evaluate their
susceptibility to CCC. Both the Tigam a j i g a n d V a r e s t r a i n t test e v a l u a t e t h e
hot-cracking susceptibility of structural w e l d m e n t s , w h i c h are G T A w e l d e d ,
by c o m p a r i n g t h e a m o u n t o f h o t c r a c k i n g p r o d u c e d b y d i f f e r e n t levels
o f an e x t e r n a l l y a p p l i e d a u g m e n t e d
strain. N o t o n l y d i d t h e s e tests i n d i c a t e
w h i c h m e t a l l u r g i c a l s y s t e m s are susc e p t i b l e t o C C C , b u t t h e y also p e r m i t ted a systematic
analysis
of
the
m o r p h o l o g y and related characteristics o f C C C in real w e l d s .
Copper Plating Procedure
C o p p e r w a s i n t r o d u c e d as a c o n t a m i n a n t o n t h e s u r f a c e o f t h e test s p e c i -
mens by electrodeposition. A n acid
Cu-plating solution,1" w i t h the c o m p o s i t i o n l i s t e d i n T a b l e 2, w a s u s e d as t h e
plating electrolyte. This solution was
stable a n d o p e r a t e d w i t h a c a t h o d e
efficiency experimentally d e t e r m i n e d
t o b e 100%. T o assure u n i f o r m p l a t i n g
quality, the electrolyte was m o n i t o r e d
by standard specific-gravity and titration measurements.
A constant-current source was built
t o s u p p l y t h e p l a t i n g c u r r e n t . Because
t h e source maintains a constant current f l o w b e t w e e n t h e Cu anodes and
t h e s p e c i m e n , t h e d e p o s i t i o n rate o f
m e t a l i o n s o n t h e s p e c i m e n s u r f a c e is
c o n s t a n t . T h i s assures a p l a t i n g c o n s i s t e n c y b e t w e e n s p e c i m e n s b e c a u s e it
eliminates the p r o b l e m s associated
w i t h the non-linear voltage-current
characteristics d u r i n g e l e c t r o d e p o s i tion.16
Table 4—Summary of the Externally Applied Strains Used for the Tigamajig Tests
Material
Inconel 750
Type 446
Rene 41
Type 410
L-605
Inconel 718
A-286
Type 430
Hastelloy-X
Waspaloy
Type 347
Type 304
Type 316
6061 T-6
Thickness,
mils
applied strains
62
56
65
74
1.0
0.9
1.1
0.90
1.0
0.9
1.0
0.8
1.1
1.0
0.9
0.9
1.1
87
75
62
105
65
86
108
107
137
128
WELDING
3.5
3.2
3.7
4.2
3.5
4.3
3.5
4.2
3.7
3.5
4.3
4.3
3.4
R E S E A R C H S U P P L E M E N T 1149-s
Table 5—Summary of Characteristics of Alloys Not Sensitive to Copper-Contamination
Cracking
Material
Principal
component
Inconel 718
Inconel 750
Rene 41
Waspaloy
Hastelloy-X
AISI Type 430
AISI Type 446
6061 T-6
Ni
Ni
Ni
Ni
Ni
Fe
Fe
Al
Solubility
of Cu in
principal
component,
Varestraint Testing Procedure
%
Crystal
structure
at MP
of Cu
100
100
100
100
100
8.9
8.9
4.1
FCC
FCC
FCC
FCC
FCC
BCC
BCC
FCC
Table 6—Summary of Characteristics of Alloys Sensitive to Copper-Contamination
Cracking
Material
Principal
component
AISI Type 304
AISI Type 316
AISI Type 321
AISI Type 347
AISI Type 410
A-286
L-605
Haynes 188
HY-80
AISI 4130
AISI 4340
Auto-body stock
Fe
Fe
Fe
Fe
Fe
Fe
Co
Co
Fe
Fe
Fe
Fe
To eliminate the problem of preferential plating of the specimen edges, a
plating tank was designed and constructed for "shadowing"" 1 the specimen edges. In this design, insulating
panels prohibit the current travel other
than in straight lines in the space
between the panels. This causes a
uniform current distribution across the
specimen surface and robs the specimen edges of plating current. Therefore, a uniform plate thickness is
produced across the surface of the
specimens w i t h minimal plating of the
specimen edges.
All specimens were cleaned before
plating by wire brushing to remove
any oxide or other contaminate layer
on the surface, followed by degreasing
with acetone, and rinsing w i t h ethanol
in an ultrasonic cleaner. The areas not
to be plated were covered w i t h masking tapes.
The thickness of the Cu electrodeposited on the surface of the specimens was controlled by varying the
plating time. As an example, the plate
thickness for a specimen w i t h 4.5 in. 2
(2900 mm 2 ) area and 400 M A plating
current w o u l d be 0.0119 mils/min
(0.302 / m i / m i n ) plating time. After
plating, the specimens were cleaned
150-s | M A Y 1978
graphic examination at higher magnification.
Solubil ty
of Cu n
princip al
c o m p o n ent,
%
Crystal
structure
at MP
of Cu
8.9
8.9
8.9
8.9
8.9
8.9
9.0
9.0
8.9
8.9
8.9
8.9
FCC
FCC
FCC
FCC
FCC
FCC
FCC
FCC
FCC
FCC
FCC
FCC
with acetone to remove any electrolyte "dragged-out" on the surface and
then stored in a desiccator.
Tigamajig Testing Procedure
The Tigamajig test was used to
examine all the alloys listed in Table 1
with the exception of HY-80, AISI 4130
and AISI 4340. Three specimens of
each alloy plated w i t h a uniform layer
of copper 0.0714 mils (1.81 fim) thick
were tested. Three unplated specimens were then tested as control specimens.
The welding current, voltage, and
arc time necessary to achieve a stable
GTA spot weld were determined using
a blank specimen of each material.
Table 3 summarizes these welding
parameters for all Tigamajig tests. For
each value of augmented strain listed
in Table 4, both an unplated control
specimen and a Cu-plated specimen
of each material were tested.
The propensity for CCC was determined by comparing the amount of
cracking observed in the plated and
control specimens. Photomacrographs
of the as-tested specimens were taken
at approximately X17, and selected
specimens were subjected to metallo-
The Varestraint test was used to
determine the susceptibility to CCC of
materials too thick to be tested w i t h
the Tigamajig testing device. The steels
were HY-80, AISI 4130, and AISI 4340.
As w i t h the Tigamajig test, the crack
susceptibility of these steels was
assessed by comparing the cracking in
as-tested, Cu-plated specimens w i t h
that of unplated control specimens.
The Varestraint test utilized a longitudinal GTA weld on specimens w h i c h
were subjected to augmented strains
as high as 4% during the test. The GTA
weld was operated at 300A dcrp and
12.5 V, a travel speed of 6.0 ipm (152.4
m m / m i n ) , an Ar shielding gas flow
rate of 45 cfh (21.1 liters/min), an Ve in.
(3.2 mm) diam. W-2% T h 0 2 electrode
with a 90 deg ground incl. angle at the
tip, and a Vt in. (3.2 mm) tip-to-work
distance.
The specimens were examined metallographically
to determine
the
amount and location of any cracks
induced by the augmented strain in
the Cu-plated and unplated control
specimens.
Results and Discussion
Twenty alloys were subjected to
screening tests to determine which
were sensitive to copper-contamination cracking (CCC). The Tigamajig
test and Varestraint test were employed for this purpose. The eight
alloys listed in Table 5 did not exhibit
CCC, while the 12 listed in Table 6
were all subject to this phenomenon.
The following generalizations can
be drawn from the results of these
screening tests and from examination
of binary phase diagrams for the major
components of the individual alloys
with Cu:
1. The major component of each of
the 12 alloys subject to CCC invariably
exhibits a limited solid solubility for
Cu (refer to Column 3, Tables 5 and
6).
2. At temperatures in the vicinity of
the melting point of Cu, all alloys
subject to CCC exhibit a FCC structure
(refer to Column 4, Table 6).
3. Neither of the alloys (Types 430
and 446) exhibiting a BCC structure at
temperatures in the vicinity of the
melting point of Cu were subject to
CCC (refer to Column 4, Table 5).
4. In no case does the major c o m p o nent of an alloy subject to CCC form a
stable intermetallic c o m p o u n d w i t h
Cu.
All of the above generalizations are
consistent w i t h the requisites for
liquid-metal embrittlement reported
in the literature. Furthermore, none of
the alloys subject to CCC exhibits
characteristics contrary to these requisites.
The Types 430 and 446 stainless steel
are not sensitive to CCC because the
liquid Cu can not be transported
through the crack of the instantaneous
crack tip by capillarity. According to
Asnis,10 the contact angle (0) for liquid
Cu on ferritic steels above the melting
point of Cu is 92-100 deg. Thus, the
capillary-transport velocity, V, approaches zero because of the cosine 0
term in the velocity expression:
V =
*W- V X \. i
Pr)
Tr cos 0
6,i L
where T is the surface tension of the
liquid, r is the distance between the
crack surfaces, /J, is the coefficient of
viscosity, and L is the crack length.
Therefore, the ferritic steels w h i c h are
not notch sensitive at 2000 F (1093 C)
could not be embrittled because the
liquid Cu-atoms can not be continually supplied to the crack tip.
It is also of interest to note that 19 of
the 20 alloys tested had melting points
above that of Cu. The 6061T-6 Al alloy,
the only exception, was not sensitive
to CCC. In addition, it should be noted
that Al forms a stable intermetallic
compound w i t h Cu. Thus, it was not
certain at this point whether the fact
that 6061T-6 is not subject to CCC is a
result of its lower melting point or its
tendency to form stable intermetallic
compounds w i t h Cu.
Tigamajig Cracking Morphologies
The as-tested specimens from the
Tigamajig screening tests described
earlier revealed a unique facet of CCC
morphology. Figure 5 shows the appearance of Cu-plated Type 304 stainless steel specimens tested at 0.0 and
0.9% augmented strain, respectively.
Note that the CCC is circumferential
and is located in the weld heataffected zone but does not extend to
the weld fusion boundary.
The preferential location of the CCC
in the weld heat-affected zone remote
from the fusion boundary indicates
that the ductility of the Cu-contaminated Type 304 stainless steel is a
minimum just above the melting point
of Cu. Furthermore, the principal
shrinkage strains in the GTA spot weld
are strictly radial in the absence of
augmented strain and, thus, w o u l d be
expected to produce circumferential
cracking. Thus, the circumferential
cracks in Fig. 5 are located where the
peak temperature present renders the
material most susceptible to CCC (just
above 1982 F, 1083 C) and are absent
where the temperature is near or
above the brittle-to-ductile transition
temperature.
\
fVs V\
,4,'S
h
* .
V
Si
vs'O ' \
Fig. 5-As-tested, copper-plated Type 304
stainless steel Tigamajig specimens: A
(top)—0.0% augmented strain; B (bottom)—0.9% augmented strain. X17 (reduced 45% on reproduction)
This is confirmed by visual observation (Fig. 5) which indicates that the
cracking in as-tested Tigamajig specimens was located adjacent to the line
where the surface luster of the specimen had a distinct change in appearance. The change in surface luster is
taken as an indication of where the
temperature on the surface of the
specimen exceeded the melting point
of Cu.
However, the CCC morphology
changed in the Cu-plated Type 304
stainless steel specimens w h e n the
level of augmented strain in the Tigamajig testing was increased to 4.3%.
Figure 6 consists of photomacro-
?1
; \
• *v^B
-SiM
•'
s
YvS Jl
A
' d§9
Fig. 6—As-tested, copper-plated Type 304
stainless steel Tigamajig specimen tested at
4.3% augmented strain: A (top)-X'l5; B
(bottom)-x60 (reduced 53% on reproduction)
graphs, taken at x 1 5 and X60 respectively, illustrating this change in CCC
morphology.
Note that the cracks are parallel to
one another and normal to the
augmented strain produced by bending the specimen over the die block.
This results from the fact that the 4.3%
augmented strain is so large as to
overpower the thermally
induced
shrinkage strains. Thus the principal
strain is oriented parallel to the
augmented strain around the entire
circumference of the GTA spot weld.
As a generalization, the Cu-contamination cracks will be oriented normal
to the principal strain. If the principal
strain is parallel to the fusion boundary
the cracks will be normal to the fusion
line (the usual case in butt and fillet
welds). O n the other hand, if the
Table 7-Summary of Alloys Predicted to be Insensitive to Copper-Contamination
Cracking
Have
extensive
solubility
w i t h Cu
Mn
Ni
Pd
Pt
Form
stable
intermetallic
compounds
w i t h Cu
Al
Be
Cd
Ga
Mg
Th
Sn
Ti
U
Zn
Have
melting
points below
that of Cu
Lack of
wetting
or solubility
w i t h Cu
Al
Cd
Ga
Ferritic steel
W
Au
Ph
Hg
Ag
Sn
Zn
W E L D I N G RESEARCH S U P P L E M E N T I 151-s
principal strain is normal to the fusion
boundary, the cracks w i l l be parallel to
the fusion line (the case for the
unstrained GTA spot weld). Where the
orientation of the principal strain
varies from one extreme to the other,
cracks of mixed orientation may
occur.
Note that the cracks in Fig. 6 extend
through the heat-affected zone to a
location much closer to the weld
fusion line than was the case for the
specimens tested at the lower strain
levels. The embrittlement of the weld
heat-affected zone nearer the weld
fusion line, w h i c h was not susceptible
to the CCC at the lower strain levels,
can be explained in terms of a strainrate effect.
W i t h the loading mechanism used
to induce the externally applied strain
in the Tigamajig specimen, the strain
rate is proportionate to the amount of
the augmented strain. Therefore, the
strain rate for tests w i t h 4.3% augmented strain w o u l d be 5 times that
for tests made with 0.9% augmented
strain. W i t h specimens of 304 stainless
steel tested at 0.0 and 0.9% strain, the
strain rate in the outer fibers was slow
enough for a nearly complete brittleto-ductile transition to occur in the
heat-affected zone which extended
nearly 0.03 in. (0.76 mm) from the
fusion boundaries. O n the other hand
at the higher strain rates associated
with the 4.3% strain, the heat-affected
zone nearer the weld fusion line was
subject to the copper-contamination
cracking.
Note, however, that even w i t h the
higher strain rate associated w i t h a
4.3% strain, the cracks invariably stop
short of the fusion boundary. This is
believed to be caused by the presence
of a nearly complete network of delta
ferrite at the austenite grain boundaries at peak temperatures above 2512
(1378 C). It should be recalled that
neither of the ferritic stainless steels
were subject to CCC in the screening
tests and, therefore, it seems reasonable that ferrite in the austenite grain
boundaries w o u l d prevent CCC.
Another possible explanation involves volatilization of the Cu by the
welding arc in the region immediately
adjacent to the weld fusion line. If the
change in surface luster present at this
location is indicative of removal of the
Cu by volatilization, the lack of
embrittlement w o u l d be expected.
The results of this experiment can be
summarized by the f o l l o w i n g :
1. Cu-contamination cracks are intergranular.
2. Cu-contamination
cracks
are
aligned perpendicular to the direction
of the principal strain.
3. Cu-contamination cracks usually
are arrested short of the fusion b o u n d ary.
4. CCC occurs preferentially in regions of the weld heat-affected zone
where the instantaneous temperature
renders the material involved most
susceptible to CCC.
Conclusion
In the final analysis, it is possible to
generalize the susceptibility of all
structural alloys to this form of weld
heat-affected embrittlement. The occurrence of CCC has been observed in
FCC Fe-base alloys and FCC Co-base
superalloys. However, the liquid-metal
embrittlement mechanism responsible
for CCC w o u l d not be possible in
alloys with the f o l l o w i n g characteristics:
1. Extensive solubility w i t h Cu.
2. A tendency to form intermetallic
compounds w i t h Cu.
3. A melting point below that of
Cu.
4. Inability to be wet by molten
Cu.
Therefore, the alloy systems that
should be insensitive to CCC have
been listed in Table 7 under headings
indicating the reason for the hypothesized insensitivity.
References
1. Matthews, S. )., Maddock, M. O., and
Savage, W. F., "How Copper Surface
Contamination Affects Weldability of Cobalt Superalloys," We/d/ng lournal, 51 (5),
May 1972, Research Suppl., pp. 326-s to
328-s.
2. Cox, A. R., and Winn, J. M., "Scaling of
Plain and Complex Copper Steels," JISI, 203
(1965), pp. 175-179.
3. Eborall, R., and Gregory, P., "The
Mechanism of Embrittlement by a Copper
Phase," /. Inst. Metals, 84 (1955), pp.
88-90.
4. Smith, C. S., "Grains, Phases and Interfaces: An Interpretation of Microstructure,"
Trans. AIME, 175 (1948), pp. 15-51.
5. Nicholson, A., and Murray, J. D., "Surface Hot Shortness in Low-Carbon Steel,"
IISI, 203, (1956), pp. 1007-1018.
6. Van Vlack, L. H., "Intergranular Energy
of Iron and Some Iron Alloys," Trans. AIME,
191 (1951), pp. 251-259.
7. Michaels, A. S., "Fundamentals of
Surface Chemistry and Surface Physics,"
Symposium on Properties of Surfaces,
ASTM Pub. #340, Philadelphia, Pa.,
(1962).
8. Bredzs, N., and Schwartzbart, H.,
"Metallurgy
of
Bonding in Brazed
Joints-Part II, Migration of Filler Metal Into
the Base Metal," Welding lournal, 38, (8),
Aug. 1959, Research Suppl., pp. 305-s to
314-s.
9. Borland, |. C, "Generalized Theory of
Super-Solidus Cracking in Welds (and Castings)," British Welding lournal, 7 (1960), pp.
508-512.
10. Asnis, E. A., and Prokhorenko, V. M.,
"Mechanism of Cracking During Welding
or Deposition of Copper onto Steel," Welding Production, 12 (1965), pp. 15-17.
11. Matthews, S. )., and Savage, W. F.,
"Heat-Affected Zone Infiltration by Dissimilar Liquid Weld Metal," Welding lournal,
50 (4), April 1971, Research Suppl., pp. 174-s
to 182-s.
12. Savage, W. F., Nippes, E. F., and Stanton, R. P., "Intergranular Attack of Alloy
Steels by Molten Copper," We/d/ng Journal,
57 (1) Jan. 1978, Research SuDpl., pp. 9-s to
16-s.
13. Stoloff, N. S„ "Liquid Metal Embrittlement," Surfaces & Interlaces, Vol. II,
Syracuse University Press, Syracuse, N. Y.,
(1968).
14. Savage, W.F., and Lundin, C. D., "The
VARESTRAINT Test," Welding lournal, 44
(10), Oct. 1965, Research Suppl., pp. 433-s to
442-s.
15. Savage, W. F., Nippes, E. F., and
Goodwin, G. M., "Effect of Minor Elements
on the Hot-Cracking Tendencies of Inconel
600," We/d/ng lournal, Research Suppl., to
be published.
16. Mohler, J. B., Electroplating, Chemical Pub. Co., New York, N. Y., (1951).
17. Perrone, N., and Liebowitz, H., "Effect of Capillary Action on Fracture Due to
Liquid Metal Embrittlement," Proceedings
of the First International Conference on
Fracture, Sendai, lapan, (1965) p. 2065.
AWS A2.4-76—Symbols for Welding and Nondestructive Testing
Provides the means for placing complete welding information on drawings. Symbols in this publication
are intended to be used to facilitate communications among designers and shop and fabrication
personnel. Many changes were made from the previous edition in recognition of the increasingly
international use of welding symbols. Illustrations showing brazing and its symbols have been added to
more clearly define usage.
Part B. Nondestructive Testing Symbols, establishes the symbols for use on drawings to specify
nondestructive examination for determining the soundness of materials.
The price of A2.4-76 Symbols for Welding and Nondestructive Testing is $5.00. Discounts: 25% to A and
B members: 20% to bookstores, public libraries and schools; 15% to C and D members. Send your orders to
the American Welding Society, 2501 N.W. 7th St., Miami, FL 33125. Florida residents add 4 % sales tax.
152-s I M A Y 1978