2. 384 S K Albert et al
composition matching that of the base metals are employed and the as-welded microstructure
of both weld metal and the heat affected zone (HAZ) can vary from predominantly ferritic to
fully martensitic depending on the composition. Data on hydrogen diffusivity and solubility
have been reported by many researchers (Lange & Knoig 1977; Sakamoto & Handa 1977;
Sakamoto & Takao 1977) and a comparison of these results indicate that apparent diffusivity
(Dapp) and apparent solubility of hydrogen in steels at low temperature (typically below
100◦
C) vary with alloy composition. Further, it is also reported that the Dapp of hydrogen
is high in fully ferritic microstructure while it is low in fully martensitic microstructure
(Sakamoto & Takao 1977; Chan et al 1990). Both these facts imply that under identical
conditions of welding, less hydrogen should diffuse out from a weld metal with high alloy
content than from that with low alloy content as the hardenability of the steels increases with
alloy content. In other words, HD content measured from a high alloy steel weldment would
be less than that measured from a low alloy steel, even if the same amount of hydrogen has
gone into the weld metal during welding for both the steels. This in turn means that the general
assumption, that the higher the HD content, the higher is the susceptibility to HAC, which is
true for carbon steel, need not be true in the case of alloy steels. In fact, it is expected that
as the alloy content increases, susceptibility of the steel to HAC should increase, while the
measured HD content should decrease.
Results of a systematic study involving estimation of HAC susceptibility, measurement
of HD content and hydrogen diffusivity of different Cr–Mo steels, conducted to investigate
this hypothesis are presented in this paper. Susceptibility to HAC was estimated using the
UT-modified hydrogen sensitivity test (UT-modified HST) and HD measurements were made
using a gas chromatograph.Apparent hydrogen diffusivity and solubility were estimated using
Devanathan’s electrochemical permeation technique (ASTM 1997).
2. Experimental
The materials used in the present study are 0·5Cr–0·5Mo, 2·25Cr–1Mo and 9Cr–1Mo steels
with their composition given in table 1. CEI given in the table is the carbon equivalent (CE)
proposed by Yurioka et al (1987) to predict maximum hardness of the heat affected zone
(HAZ) of alloy steels (including Cr–Mo steels) in the as-welded condition.
CEI = C + Si/24 + Mn/6 + Cu/15 + Ni/12 + {Cr(1 − 0·16
√
Cr)}/8 + Mo/4.
(1)
For both UT-modified HST and HD measurement, these steels were used in the normalised
and tempered condition, while for hydrogen diffusivity measurement (carried out only for
2·25Cr–1Mo and 9Cr–1Mo steels), they were austenitised at 940◦
C for 30 min and water
quenched to produce a microstructure that may be present in the weld metal and HAZ after
welding.
Table 1. Chemical composition of the steels (wt.%).
Element C Cr Mn Si S Mo CEI
0·5Cr–0·5Mo steel 0·22 0·5 0·30 – 0·009 – 0·43
2·25Cr–1Mo steel 0·12 2·18 0·46 0·25 0·001 1·0 0·68
9Cr–1Mo steel 0·072 8·24 0·36 0·21 0·001 1·0 0·93
3. Hydrogen-assisted cracking susceptibility of Cr–Mo steel welds 385
Figure 1. Straining jig for UT-modified HST.
2.1 UT-modified HST
This test is a modified version of the RPI augmented strain test (Savage et al 1976) and
was originally developed at the University of Tennessee (Lundin et al 1986, 1990). For this
test, small specimens of dimensions 3 × 15 × 40 mm were used and welding was carried
out using the autogenous gas tungsten arc welding (GTAW) process with current = 90A,
voltage = 12V, speed = 12 mm/s and Ar gas flow rate = 10 litres/minute. Hydrogen was
introduced into the weld pool by mixing it with shielding gas. During welding, the specimen
was held in a copper fixture, which can be preheated to the desired temperature.After welding,
the specimen was allowed to cool to room temperature (in the case of preheating, specimen is
cooled to the preheat temperature) and strained in a fixture as shown in figure 1. The nominal
augmented strain on the surface is given approximately as
ε ≈ t/2R, (2)
whereε = nominal augmented strain,t = specimen thicknessandR = bending radius of the
die. In the present study, the R and t are chosen such that a strain of ∼ 4% is felt by the
specimen. The susceptibility to cracking is determined by observing crack formation on the
specimen face strained in tension for 24h. Hydrogen content in the shielding gas was varied
from 1 to 5 vol.%. For a given hydrogen content in the shielding gas, the preheat tempera-
ture above which no cracking occurred is taken as the critical preheat temperature for that
hydrogen level.
2.2 Diffusible hydrogen (HD) measurement
For HD measurement, specimens were prepared in exactly the same way as for the UT-
modified HST. Immediately after welding, specimens were removed from the copper fixture
and transferred to a stainless steel chamber provided with an inlet and outlet connection for gas
flow (the chamber was He leak-tested and the leak rate was found to be below 10−9
std cc/min
at a pressure of 10−8
kg/cm2
). Immediately after transferring the specimen, the chamber was
first flushed and then filled with Ar at 2 kg/cm2
and transferred to an oven maintained at
45◦
C. Hydrogen evolved from the specimen was collected inside the chamber for 72h and
the concentration of hydrogen in the chamber was measured using a gas chromatograph.
Knowing the volume of the chamber, and pressure of the gas inside the chamber, total volume
4. 386 S K Albert et al
Figure 2. Schematic of HD mea-
surement set-up.
of hydrogen evolved from the specimen was estimated. A schematic of the measurement set
up is shown in figure 2.
For estimating HD in ml/100g of fused metal, it was necessary to determine the mass of
the fused weld metal. For this purpose, the volume of the weld metal was estimated from the
length of the weld bead and the average area of cross-section of the weld. Density of the weld
metal was taken as 7·9 g/cm3
. For a given vol.% of hydrogen in the shielding gas, HD content
was estimated for a minimum of four specimens.
2.3 Measurement of hydrogen permeability
Circular specimens (diameter = 25 mm and thickness = 1·2 mm) were machined out from
water-quenched steels for this measurement. One side of the specimen was coated with pal-
ladium before introducing into the electrochemical permeation cell. A schematic of the per-
meation cell is shown in figure 3. It essentially consists of two polarisation cells, one oper-
ated at cathodic potential and the other at anodic. The palladium-coated surface faces the
anodic compartment. The electrolyte in anodic compartment is 0·1M NaOH solution while
Figure 3. Schematic of electrochemical
permeation cell.
5. Hydrogen-assisted cracking susceptibility of Cr–Mo steel welds 387
that in cathodic compartment is 0·5M H2SO4 with 200ppm As2O3. Both the compartments
are purged with Ar gas. More details of the experimental set-up are given elsewhere (Par-
vathavarthini et al 1999).
Initially the anodic compartment was filled with 0·1M NaOH, −40 mV (SCE) applied to
the specimen, and the anodic current monitored. After the anodic current stabilised, the cath-
ode compartment was also filled with electrolyte. Using a potentiostat in galvanostatic mode,
a cathodic current of 0.05 mA/cm2
was applied. Hydrogen permeated through the specimens
to the anodic side where it was instantaneously oxidised and turned into an equivalent cur-
rent. Therefore, the permeating current density (Pt ) at the exit side is a direct measure of
the output flux of hydrogen. The variation in Pt with respect to time is monitored until a
steady state is reached. The rate of hydrogen permeation rises after a certain break through
time (tb) and then approaches asymptotically to the steady state permeation rate, P∞. From
the steady state permeation current density, P∞, permeability (p) is calculated using the
expression,
p = (P∞ ∗ L)/(Z ∗ F), (3)
where L is the thickness of the specimen, Z is the number of electrons participating in the
reaction and F is Faraday’s constant.
From the breakthrough time, apparent diffusion coefficient (Dapp) is calculated using the
following equation.
D = L2
/15·3tb. (4)
From Dapp and p, solubility (S) is calculated, since
p = Dapp ∗ S. (5)
3. Results
Results of the UT-modified HST are shown in figure 4. Here critical preheat temperature
determined from the cracking tests is plotted against volume % of hydrogen in the shielding
gas. Cracking susceptibility is highest for 9Cr–1Mo steel for which critical preheat tempera-
ture increases from 175 to 225◦
C as the of hydrogen in Ar increases from 1 to 5 vol.%. For
the other two steels, the corresponding change in the critical preheat temperature is from 100
to 175◦
C.
Variation in the HD content in the weld for different volume % of hydrogen in the shielding
gas for these steels is shown in figure 5. The HD content increases with increase in volume
% of hydrogen in the shielding gas and this variation is approximately linear for all the three
steels. Further, for a given volume % of hydrogen, HD content is lower for 9Cr–1Mo steel
than for 2·25Cr–1Mo steel. For 9Cr–1Mo steel it increases from 1 to 2ml/100g of weld metal
when hydrogen in the shielding gas increases from 1 to 5 vol.%, while the corresponding
increases for 2·25Cr–1Mo and 0·5Cr–0.5Mo steels are from 2 to 5ml/100g and 2 to 6ml/100g
respectively. Thus the results clearly show that for identical experimental conditions, HD
content measured in 9Cr–1Mo steel is substantially lower than that measured in the other two
steels.
In actual welding conditions, unlike in the present study, the main source of hydrogen is the
moisture content in the electrode coating. It is difficult to quantify the hydrogen that enters
6. 388 S K Albert et al
Figure 4. Variation of critical preheat tem-
perature with vol.% of hydrogen in the shield-
ing gas.
the weld metal from the moisture that may be present in the electrode coatings. Further, not
all the hydrogen, but only the diffusible hydrogen that is present in the weld contributes to
cracking. Due to these reasons, it is more appropriate to represent critical preheat temperature
as a function of HD content rather than of concentration of hydrogen in the shielding gas.
Such a diagram is shown in figure 6. It may be seen that for all the three steels critical preheat
temperature increases with HD content. Further it also shows that under identical conditions
of welding, susceptibility is higher and HD content is lower for 9Cr–1Mo steel than for the
other two steels.
Results obtained from permeation studies are shown in table 2. Apparent hydrogen diffu-
sivity and permeability are obtained directly from the measurements while apparent solubility
is estimated from diffusivity and permeability using (5). For each steel, four separate mea-
Figure 5. Variation of diffusible hydrogen
content in the weld with vol.% of hydrogen
in the shielding gas.
7. Hydrogen-assisted cracking susceptibility of Cr–Mo steel welds 389
Figure 6. Variation of critical preheat tem-
perature with diffusible hydrogen content.
surements are carried out and data shown here are averages of these measurements. Results
clearly show that apparent diffusivity of hydrogen in 9Cr–1Mo steel is an order of magnitude
lower than in 2·25Cr–1Mo steel. Further apparent solubility of hydrogen in the former steel
is significantly higher than in the latter.
4. Discussion
4.1 Effect of alloy composition on HD content of Cr–Mo steel welds
It is known that susceptibility HAC would increase with increase in alloy content as the
hardenability of the steel increases with alloy content. However, the effect of composition
on HD is not widely known. Results indicate that under identical condition of measurement,
diffusible hydrogen content decreases with increase in alloy content. A regression analysis
of the results on HD measurements was carried out to represent the HD as a function of both
composition and hydrogen content in the shielding gas. For this, hydrogen in the shielding
gas was represented as partial pressure of hydrogen, PH2, instead of vol.% (PH2 for 5vol.% is
taken as 0.05). To represent composition, CEI, carbon equivalent, which has been found to be
applicable for a wide range of composition and used to predict the hardness fairly accurately
for alloy steels including Cr–Mo steels (Yurioka et al 1987; Albert et al 1996) was chosen.
Analysis provided the following relation with an R2
value of 0·8951 and standard error of
determination of 0·5,
HD = 50·8PH2 − 5·8CEI + 5·7. (6)
Table 2. Results of hydrogen permeation studies.
Steel Dapp × 108
(cm2
/s) P × 1012
(moles/cm.s) Sapp × 106
(moles/cm3
)
2·25Cr–1Mo steel 23·31 2·54 10·27
9Cr–1Mo steel 1·28 2·19 175·2
8. 390 S K Albert et al
Figure 7. Comparison of measured and
estimated value of diffusible hydrogen con-
tent.
Figure 7 shows the comparison of the measured values of HD and those estimated from (6).
It is seen that the variation in HD for different alloys tested under identical conditions can
be represented as a function of composition, and CEI, the carbon equivalent used to predict
HAZ hardness of the alloy steels, is suited to represent composition.
4.2 Effect of alloy composition on hydrogen diffusivity and solubility
Results of the HD measurements and the discussion above clearly indicated there is a system-
atic decrease in the HD content with increase in alloy content. As specimen preparation and
diffusible hydrogen measurements are carried out under identical conditions, it is reasonable
to assume that the physical basis for this variation could be the differences in the apparent
diffusivity and solubility of hydrogen in steels of different composition. The steel that gives
high HD values should have high diffusivity and low solubility. Diffusivity and solubility data
for 2·25Cr–1Mo and 9Cr–Mo steels from permeability studies given in table 2 support this
assumption.
As already mentioned, permeation measurements were carried out for both 2·25Cr–1Mo
and 9Cr–1Mo steels which were austenitised at 940◦
C for 30 minutes followed by water
quenching. Under these conditions, the structure of both these steels would be martensitic.
Hence, the major difference between the specimens used for permeation studies from these
two steels would be in their Cr content. It is interesting to find out how the diffusivity of
hydrogen varies with CE1 for various Cr–Mo steels. For this hydrogen diffusivity data for
various wrought Cr–Mo steels, including that from the present study, and Cr–Mo steel weld
metals reported in literature (Moreton et al 1971; Sakamoto & Handa 1977; Sakamoto &
Takao 1977; Kushida & Kudo 1991; Valentini & Solina 1994; Parvathavarthini 1995) were
plotted against the composition. This is shown in figure 8 (numbers given in brackets are the
references for the source of the data). In the case of wrought steels, diffusivity of hydrogen in
the as-quenched condition and in the case of weld metal the data in the as-welded condition
were taken. Diffusivity decreases by two orders of magnitude as the CEI varies from 0·25
(mild steel) to 1·1(9Cr–1Mo steel). The trend is similar even for steels with major alloying
9. Hydrogen-assisted cracking susceptibility of Cr–Mo steel welds 391
Figure 8. Variation of hydrogen diffusivity at room temperature with composition.
elements other than Cr (Bollinghaus et al 1996).Apparent solubility also was found to increase
with alloy content as observed in the present study. It is reported that for a 1Cr–0·5Mo steel
with CEI = 0·67, it is around 4 − 5 × 10−5
mol/cm3
Fe (Sakamoto & Handa 1977) and for
12Cr steel (CEI = 0·94) it is around 3 × 10−3
mol/cm3
Fe (Sakamoto & Takao 1977).
From the above discussion it appears that the apparent diffusivity and solubility of hydrogen
in steel have significant influence on the measured diffusible hydrogen content. It is also
known that the diffusivity and solubility of hydrogen in a steel strongly depend on trapping of
hydrogen by various defects like dislocation, grain boundaries, matrix-particle interface etc.
(Gibala & Kummick 1985). Depending on the energy of traps, hydrogen traps are broadly
classified as weak and strong traps. Hydrogen atoms trapped by strong traps are released only
if heated to very high temperatures and do not contribute to diffusivity at ambient temperature
andhenceHAC.Suchtrapsinclude,interfacebetweenAlN,Fe3CandTiC(Gibala&Kummick
1985). In the case of alloy steels with as quenched or as-welded microstructure, strong traps
like particle-matrix interface would be less as most of the precipitates would have gone into
solution during heating. Hence, most of the traps present in these conditions would be of low
binding energy like solute atoms, dislocations, lath and prior austenite boundaries etc. (Aosaka
1982; Lacombe et al 1985). Density of such traps, especially that of dislocations would be
very high in an as-welded structure. Further, there would be significant differences in the trap
density of 9Cr–1Mo and 2·25Cr–1Mo steels due to variation in the solute content.At ambient
temperatures, these traps cannot retain hydrogen permanently. The net result of trapping of
hydrogen atoms by these defects would be to decrease the diffusion rate of hydrogen, which
in turn results in an apparent increase in the solubility at low temperature. This decrease
in diffusivity and increase in solubility of hydrogen would increase with increase in alloy
content. As hydrogen-assisted cracking occurs as a result of interaction of hydrogen with
the defects in the steel, high apparent solubility and low apparent diffusivity of hydrogen at
ambient temperatures should make it more susceptible to cracking. This explains why 9Cr–
1Mo steel with higher apparent solubility and lower diffusivity for hydrogen shows higher
10. 392 S K Albert et al
susceptibility to HAC and gives lower HD content than 2·25Cr–1Mo steels under identical
conditions of testing and measurement.
5. Conclusions
From the results discussed above, the following conclusions can be drawn regarding the
susceptibility of Cr–Mo steels to HAC:
(1) Susceptibility of 9Cr–1Mo steel to HAC is higher than that of 2·25Cr–1Mo and 0·5Cr–
0·5Mo steels.
(2) Under identical conditions of testing and measurement, HD content in 9Cr–1Mo steel is
significantly lower than that of the other two steels.
(3) Significant differences in the apparent solubility and diffusivity of hydrogen between
9Cr–1Mo and 2·25Cr–1Mo steels explain the large differences in their HD contents.
(4) Variations in HD content and hydrogen diffusivity can be represented as functions of CE1,
carbon equivalent proposed to predict hardenability in alloy steels.
(5) Assuming that the higher the HD content, the higher is the susceptibility to HAC, nor-
mally true in the case of carbon steel consumables, does not seem to be valid in the case
of alloy steels.
Authors acknowledge the support and encouragement given by Dr Baldev Raj, S L Mannan,
S K Ray and R K Dayal during the course of this work. Support by Ms K Parimala in carrying
out the permeation studies is also gratefully acknowledged.
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