Influence of Alloying on Hydrogen-Assisted Cracking

Sadhana Vol. 28, Parts 3 & 4, June/August 2003, pp. 383393. Printed in India

Influence of alloying on hydrogen-assisted cracking anddiffusible hydrogen content in CrMo steel welds

S K ALBERT, V RAMASUBBU, N PARVATHAVARTHINI andT P S GILL

Materials Technology Division, and Corrosion Science and TechnologyDivision, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, Indiae-mail: [email protected]

Abstract. Study of hydrogen-assisted cracking and measurement of diffusiblehydrogen content in different CrMo steel welds shows that under identical condi-tions, susceptibility to cracking increased and diffusible hydrogen content decreasewith increase in alloy content. Hydrogen permeation studies show that hydrogendiffusivity decreases and solubility increases with increase in alloy content. Thusdecrease in diffusible hydrogen content with increase in alloying is attributed toincrease in apparent solubility and decrease in apparent diffusivity of hydrogen.Analysis of the results indicates that variation of diffusible hydrogen content andapparent diffusivity of hydrogen with alloy content can be represented as a func-tion of carbon equivalent CE1 originally proposed to predict the hardness in theheat-affected zone of alloy steel welds.

Keywords. Hydrogen-assisted cracking; diffusible hydrogen content; hydrogendiffusivity; CrMo steels; welding; carbon equivalent.

1. Introduction

Carbon and alloy steel welds are susceptible to hydrogen-assisted cracking (HAC).As the alloycontent increases, susceptibility to HAC also increases due to increase in the hardenabilityof the steel. Preheating and post heating of the weldments and use of low hydrogen weldingconsumables are employed to avoid this form of cracking. The main source of hydrogen is themoisture in the welding consumables which dissociate in the welding arc to form hydrogen andoxygen. The susceptibility of the weldment to HAC is assessed from the hydrogen diffused outfrom the weld after the welding is over. Hydrogen thus diffused out is referred to as diffusiblehydrogen (HD) and is estimated from the measurement of hydrogen evolved from a weld beadmade by the welding consumable on a standard steel plate by employing specified weldingparameters. During this measurement it is assumed that HD content does not depend on thecomposition of the weld metal. Thus, it is accepted that the higher the HD content, the higher isthe susceptibility of the weldment to HAC. This assumption is true in the case of carbon steelwelding consumables, which contain no significant alloying additions and the compositionis controlled in such a way that the weld metal microstructure is predominantly ferritic.However, in the case of welding of alloy steels like CrMo steels, welding consumables with

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384 S K Albert et al

composition matching that of the base metals are employed and the as-welded microstructureof both weld metal and the heat affected zone (HAZ) can vary from predominantly ferritic tofully martensitic depending on the composition. Data on hydrogen diffusivity and solubilityhave 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 below100C) vary with alloy composition. Further, it is also reported that the Dapp of hydrogenis 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 identicalconditions of welding, less hydrogen should diffuse out from a weld metal with high alloycontent than from that with low alloy content as the hardenability of the steels increases withalloy content. In other words, HD content measured from a high alloy steel weldment wouldbe less than that measured from a low alloy steel, even if the same amount of hydrogen hasgone into the weld metal during welding for both the steels. This in turn means that the generalassumption, that the higher the HD content, the higher is the susceptibility to HAC, which istrue for carbon steel, need not be true in the case of alloy steels. In fact, it is expected thatas the alloy content increases, susceptibility of the steel to HAC should increase, while themeasured HD content should decrease.

Results of a systematic study involving estimation of HAC susceptibility, measurementof HD content and hydrogen diffusivity of different CrMo steels, conducted to investigatethis hypothesis are presented in this paper. Susceptibility to HAC was estimated using theUT-modified hydrogen sensitivity test (UT-modified HST) and HD measurements were madeusing a gas chromatograph.Apparent hydrogen diffusivity and solubility were estimated usingDevanathans electrochemical permeation technique (ASTM 1997).

2. Experimental

The materials used in the present study are 05Cr05Mo, 225Cr1Mo and 9Cr1Mo steelswith 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 CrMo steels) in the as-welded condition.

CEI D C C Si/24 C Mn/6 C Cu/15 C Ni/12 C fCr(1 016pCr)g/8 C Mo/4.(1)

For both UT-modified HST and HD measurement, these steels were used in the normalisedand tempered condition, while for hydrogen diffusivity measurement (carried out only for225Cr1Mo and 9Cr1Mo steels), they were austenitised at 940C for 30 min and waterquenched to produce a microstructure that may be present in the weld metal and HAZ afterwelding.

Table 1. Chemical composition of the steels (wt.%).

Element C Cr Mn Si S Mo CEI

05Cr05Mo steel 022 05 030 0009 043225Cr1Mo steel 012 218 046 025 0001 10 0689Cr1Mo steel 0072 824 036 021 0001 10 093

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effect of Boron

Hydrogen-assisted cracking susceptibility of CrMo steel welds 385

Figure 1. Straining jig for UT-modified HST.

2.1 UT-modified HSTThis test is a modified version of the RPI augmented strain test (Savage et al 1976) andwas originally developed at the University of Tennessee (Lundin et al 1986, 1990). For thistest, small specimens of dimensions 3 15 40 mm were used and welding was carriedout using the autogenous gas tungsten arc welding (GTAW) process with current D 90A,voltage D 12V, speed D 12 mm/s and Ar gas flow rate D 10 litres/minute. Hydrogen wasintroduced into the weld pool by mixing it with shielding gas. During welding, the specimenwas 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 iscooled to the preheat temperature) and strained in a fixture as shown in figure 1. The nominalaugmented strain on the surface is given approximately as

t/2R, (2)where D nominal augmented strain, t D specimen thickness andR D bending radius of thedie. In the present study, the R and t are chosen such that a strain of 4% is felt by thespecimen. The susceptibility to cracking is determined by observing crack formation on thespecimen face strained in tension for 24 h. Hydrogen content in the shielding gas was variedfrom 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 thathydrogen level.

2.2 Diffusible hydrogen (HD) measurementFor 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 fixtureand transferred to a stainless steel chamber provided with an inlet and outlet connection for gasflow (the chamber was He leak-tested and the leak rate was found to be below 109 std cc/minat a pressure of 108 kg/cm2). Immediately after transferring the specimen, the chamber wasfirst flushed and then filled with Ar at 2 kg/cm2 and transferred to an oven maintained at45C. Hydrogen evolved from the specimen was collected inside the chamber for 72 h andthe 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

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Figure 2. Schematic of HD mea-surement set-up.

of hydrogen evolved from the specimen was estimated. A schematic of the measurement setup is shown in figure 2.

For estimating HD in ml/100g of fused metal, it was necessary to determine the mass ofthe fused weld metal. For this purpose, the volume of the weld metal was estimated from thelength of the weld bead and the average area of cross-section of the weld. Density of the weldmetal was taken as 79 g/cm3. For a given vol.% of hydrogen in the shielding gas, HD contentwas estimated for a minimum of four specimens.

2.3 Measurement of hydrogen permeabilityCircular specimens (diameter D 25 mm and thickness D 12 mm) were machined out fromwater-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 theanodic compartment. The electrolyte in anodic compartment is 01 M NaOH solution while

Figure 3. Schematic of electrochemicalpermeation cell.

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that in cathodic compartment is 05 M H2SO4 with 200 ppm As2O3. Both the compartmentsare 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 01 M NaOH, 40 mV (SCE) applied tothe 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 specimensto 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 ofthe output flux of hydrogen. The variation in Pt with respect to time is monitored until asteady state is reached. The rate of hydrogen permeation rises after a certain break throughtime (tb) and then approaches asymptotically to the steady state permeation rate, P1. Fromthe steady state permeation current density, P1, permeability (p) is calculated using theexpression,

p D (P1 L)/(Z F), (3)where L is the thickness of the specimen, Z is the number of electrons participating in thereaction and F is Faradays constant.

From the breakthrough time, apparent diffusion coefficient (Dapp) is calculated using thefollowing equation.

D D L2/153tb. (4)From Dapp and p, solubility (S) is calculated, since

p D Dapp S. (5)

3. Results

Results of the UT-modified HST are shown in figure 4. Here critical preheat temperaturedetermined from the cracking tests is plotted against volume % of hydrogen in the shieldinggas. Cracking susceptibility is highest for 9Cr1Mo steel for which critical preheat tempera-ture increases from 175 to 225C as the of hydrogen in Ar increases from 1 to 5 vol.%. Forthe other two steels, the corresponding change in the critical preheat temperature is from 100to 175C.

Variation in the HD content in the weld for different volume % of hydrogen in the shieldinggas 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 threesteels. Further, for a given volume % of hydrogen, HD content is lower for 9Cr1Mo steelthan for 225Cr1Mo steel. For 9Cr1Mo steel it increases from 1 to 2 ml/100 g of weld metalwhen hydrogen in the shielding gas increases from 1 to 5 vol.%, while the correspondingincreases for 225Cr1Mo and 05Cr0.5Mo steels are from 2 to 5 ml/100 g and 2 to 6 ml/100 grespectively. Thus the results clearly show that for identical experimental conditions, HDcontent measured in 9Cr1Mo steel is substantially lower than that measured in the other twosteels.

In actual welding conditions, unlike in the present study, the main source of hydrogen is themoisture content in the electrode coating. It is difficult to quantify the hydrogen that enters

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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, notall the hydrogen, but only the diffusible hydrogen that is present in the weld contributes tocracking. Due to these reasons, it is more appropriate to represent critical preheat temperatureas 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 preheattemperature increases with HD content. Further it also shows that under identical conditionsof welding, susceptibility is higher and HD content is lower for 9Cr1Mo steel than for theother 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 solubilityis estimated from diffusivity and permeability using (5). For each steel, four separate mea-

Figure 5. Variation of diffusible hydrogencontent in the weld with vol.% of hydrogenin the shielding gas.

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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. Resultsclearly show that apparent diffusivity of hydrogen in 9Cr1Mo steel is an order of magnitudelower than in 225Cr1Mo steel. Further apparent solubility of hydrogen in the former steelis significantly higher than in the latter.

4. Discussion

4.1 Effect of alloy composition on HD content of CrMo steel weldsIt is known that susceptibility HAC would increase with increase in alloy content as thehardenability of the steel increases with alloy content. However, the effect of compositionon HD is not widely known. Results indicate that under identical condition of measurement,diffusible hydrogen content decreases with increase in alloy content. A regression analysisof the results on HD measurements was carried out to represent the HD as a function of bothcomposition and hydrogen content in the shielding gas. For this, hydrogen in the shieldinggas was represented as partial pressure of hydrogen, PH2, instead of vol.% (PH2 for 5 vol.% istaken as 0.05). To represent composition, CEI, carbon equivalent, which has been found to beapplicable for a wide range of composition and used to predict the hardness fairly accuratelyfor alloy steels including CrMo steels (Yurioka et al 1987; Albert et al 1996) was chosen.Analysis provided the following relation with an R2 value of 08951 and standard error ofdetermination of 05,

HD D 508PH2 58CEI C 57. (6)

Table 2. Results of hydrogen permeation studies.

Steel Dapp 108(cm2/s) P 1012(moles/cm.s) Sapp 106(moles/cm3)

225Cr1Mo steel 2331 254 10279Cr1Mo steel 128 219 1752

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Figure 7. Comparison of measured andestimated 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 canbe represented as a function of composition, and CEI, the carbon equivalent used to predictHAZ hardness of the alloy steels, is suited to represent composition.

4.2 Effect of alloy composition on hydrogen diffusivity and solubilityResults 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 anddiffusible hydrogen measurements are carried out under identical conditions, it is reasonableto assume that the physical basis for this variation could be the differences in the apparentdiffusivity and solubility of hydrogen in steels of different composition. The steel that giveshigh HD values should have high diffusivity and low solubility. Diffusivity and solubility datafor 225Cr1Mo and 9CrMo steels from permeability studies given in table 2 support thisassumption.

As already mentioned, permeation measurements were carried out for both 225Cr1Moand 9Cr1Mo steels which were austenitised at 940C for 30 minutes followed by waterquenching. Under these conditions, the structure of both these steels would be martensitic.Hence, the major difference between the specimens used for permeation studies from thesetwo steels would be in their Cr content. It is interesting to find out how the diffusivity ofhydrogen varies with CE1 for various CrMo steels. For this hydrogen diffusivity data forvarious wrought CrMo steels, including that from the present study, and CrMo steel weldmetals reported in literature (Moreton et al 1971; Sakamoto & Handa 1977; Sakamoto &Takao 1977; Kushida & Kudo 1991; Valentini & Solina 1994; Parvathavarthini 1995) wereplotted against the composition. This is shown in figure 8 (numbers given in brackets are thereferences for the source of the data). In the case of wrought steels, diffusivity of hydrogen inthe as-quenched condition and in the case of weld metal the data in the as-welded conditionwere taken. Diffusivity decreases by two orders of magnitude as the CEI varies from 025(mild steel) to 11(9Cr1Mo steel). The trend is similar even for steels with major alloying

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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 increasewith alloy content as observed in the present study. It is reported that for a 1Cr05Mo steelwith CEI D 067, it is around 4 5 105 mol/cm3Fe (Sakamoto & Handa 1977) and for12Cr steel (CEI D 094) it is around 3 103 mol/cm3 Fe (Sakamoto & Takao 1977).

From the above discussion it appears that the apparent diffusivity and solubility of hydrogenin steel have significant influence on the measured diffusible hydrogen content. It is alsoknown that the diffusivity and solubility of hydrogen in a steel strongly depend on trapping ofhydrogen by various defects like dislocation, grain boundaries, matrix-particle interface etc.(Gibala & Kummick 1985). Depending on the energy of traps, hydrogen traps are broadlyclassified as weak and strong traps. Hydrogen atoms trapped by strong traps are released onlyif heated to very high temperatures and do not contribute to diffusivity at ambient temperatureand hence HAC. Such traps include, interface betweenAlN, Fe3C andTiC (Gibala & Kummick1985). In the case of alloy steels with as quenched or as-welded microstructure, strong trapslike particle-matrix interface would be less as most of the precipitates would have gone intosolution during heating. Hence, most of the traps present in these conditions would be of lowbinding energy like solute atoms, dislocations, lath and prior austenite boundaries etc. (Aosaka1982; Lacombe et al 1985). Density of such traps, especially that of dislocations would bevery high in an as-welded structure. Further, there would be significant differences in the trapdensity of 9Cr1Mo and 225Cr1Mo steels due to variation in the solute content. At ambienttemperatures, these traps cannot retain hydrogen permanently. The net result of trapping ofhydrogen atoms by these defects would be to decrease the diffusion rate of hydrogen, whichin turn results in an apparent increase in the solubility at low temperature. This decreasein diffusivity and increase in solubility of hydrogen would increase with increase in alloycontent. As hydrogen-assisted cracking occurs as a result of interaction of hydrogen withthe defects in the steel, high apparent solubility and low apparent diffusivity of hydrogen atambient temperatures should make it more susceptible to cracking. This explains why 9Cr1Mo steel with higher apparent solubility and lower diffusivity for hydrogen shows higher

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susceptibility to HAC and gives lower HD content than 225Cr1Mo steels under identicalconditions of testing and measurement.

5. Conclusions

From the results discussed above, the following conclusions can be drawn regarding thesusceptibility of CrMo steels to HAC:

(1) Susceptibility of 9Cr1Mo steel to HAC is higher than that of 225Cr1Mo and 05Cr05Mo steels.

(2) Under identical conditions of testing and measurement, HD content in 9Cr1Mo steel issignificantly lower than that of the other two steels.

(3) Significant differences in the apparent solubility and diffusivity of hydrogen between9Cr1Mo and 225Cr1Mo 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 caseof 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 carryingout the permeation studies is also gratefully acknowledged.

References

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Influence of Alloying on Hydrogen-Assisted Cracking

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