17-4 PH Stainless Steel Weldments.pdf

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Effect of Applied Potentialson Environmental Cracking Behaviorof 17-4 PH Stainless Steel Weldments

K.S. Raja and K.P. Rao*

ABSTRACT

The effects of anodic, cathodic, and open-circuit potentials(OCP) on the environmental cracking behavior of 17% Cr-4%Ni (17-4 [UNS S17400]) precipitation-hardenable (PH) stain-less steel (SS) welds subjected to different thermaltreatments were studied. Sheets of 17-4 PH SS 1.5 mm(0.059 in.) thick and in solution-treated condition were full-penetration welded autogenously using the gas tungsten arcwelding process (GTAW). Weldments were given one of twopostweld heat treatments: direct aging and solution treat-ment + aging. Samples were aged at 480°C for 1 h , 510°Cfor 4 h, and 600°C for 4 h. Environmental cracking testswere conducted using U-bend samples. Samples were testedin 3.5% sodium chloride (NaCl) solution (pH = 2.0) undervarious applied potentials. The failure time at each potentialwas taken as the criterion for cracking resistance of thesamples. At OCP and at anodic potentials, cracking wasfound to occur by an active path dissolution mechanism inthe heat-affected zone (HAZ). Although the weld metal hard-ness was more than that of the HAZ, no cracking took placein the weld metal. Solution treating the welds improved theircracking resistance. In contrast, failure occurred within theweld metal at applied cathodic potentials. Therefore, thehardness criterion (the harder the structure, the higher thesusceptibility to cracking) was found to be applicable underapplied cathodic potentials. The hardest structure obtainedby peak aging showed the least cracking resistance underthese potentials. To increase the cracking resistance of weldmetals under cathodic potentials, solution treating followedby overaging (at 600°C for 4 h) was found to be the best

4980010-9312/95/0001

© 1995, NACE

Submitted for publication December 1993; in revised form,December 1994.

* Department of Metallurgical Engineering, Indian Institute ofTechnology, Madras, 600 036, India.

method. However, the same treatment was found to behighly detrimental under anodic and OCP conditions.

KEYWORDS: aging, cracking, hardness, open-circuitpotential, potential, stainless steel, 17-4 PH stainless steel,weldments

INTRODUCTION

The need for steels for use in moderate operatingtemperatures (< 400°C) and with high strength-to-weight ratios led to development of materials such asprecipitation-hardenable (PH) stainless steels (SS).

PH SS are easy to fabricate because of their softand ductile nature. After fabrication, they can behardened by simple heat treatments that do not in-volve high temperatures.

There are three types of PH SS available:—␣ The transformable or martensitic type,—␣ The controlled transformation or

semiaustenitic type, and—␣ The nontransformable or austenitic type.The martensitic type is used most widely in engi-

neering applications because of its simplicity in heattreatments. In the case of martensitic PH steels, com-plete martensitic transformation is achieved by aircooling the alloy from its solution-treated tempera-ture (~ 1,050°C) to room temperature. The chemicalcomposition of the alloy is controlled so that thestarting and completion temperatures of the marten-sitic transformation (Ms and Mf, respectively) are wellabove room temperature.

Only the mechanical properties have drawn themaximum attention of researchers investigating

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(a)

(b)

(c)

FIGURE 1. (a) Initial shape of the block (scale 1:2), (b) weldedcoupon, (c) final shape of the coupon (weld bead flushed, with heattreatment at this stage, and (d) U-bend sample. (Dimensions in mm,radius (R) = 10 mm).

(d)

postweld heat treatments (PWHT). Even though PHSS are used for their better corrosion properties andexcellent mechanical properties, very little informa-tion is available in the open literature on thecorrosion properties of PH SS weldments. Goochcompared the stress corrosion cracking (SCC) behav-ior of a semiaustenitic SS weld metal with otherhigh-strength alloys, such as maraging steel welds.1

He concluded that the retained austenite provided forbetter SCC resistance in the case of electron beamwelded, semiaustenitic PH SS weld metal.

Generally, a solution treating + aging PWHT isgiven to 17-4 PH SS (UNS S17400)(1) joints to preventsevere corrosion of the HAZ.2 For service in hot con-centrated chloride media, an aging PWHT at 550°C to620°C generally is carried out.3-4 Available reportsindicate the HAZ in martensitic PH steel (such as17-4 PH SS) is more prone to severe corrosion in theas-welded condition.2,5-6 PH SS are susceptible tocracking under tensile loading in several corrosiveenvironments. Truman found PH SS prone to SCC,hydrogen-assisted cracking (HAC), hydrogen em-brittlement (HE), and sulfide stress cracking (SSC).7

Different environmental conditions can be simu-lated by using a potentiostat and applying differentpotentials to the specimens. The objective of thepresent work was to study the environmental crack-ing behavior of 17-4 PH SS welds under differentpotential conditions using a potentiostat.

EXPERIMENTAL

The steel was in the form of 17-4 PH SS sheets1.5 mm (0.059 in.) thick. Its chemical compositionwas (wt%): 16.8% Cr, 4.2% Ni, 3.8% Cu, 0.054% C,0.3% Si, 0.6% Mn, 0.22% Nb, 0.002% S, 0.02% P,and balance Fe. The specimens were solution treatedand autogenous full-penetration welded by a gastungsten arc welding (GTAW) process (100 A, 20 V,and 20 cm/min [7.87 in./min] speed with a directcurrent electrode negative). After welding, these cou-pons were ground on abrasive belts on both sides to1 mm (0.039 in.) thick so that the uneven surfacescaused by welding were removed and so that the finalsample size of 90 mm by 10 mm by 1 mm (3.54 in.by 0.39 in. by 0.039 in.) was obtained.

The welded plates were directly aged and solu-tion treated + aged. The solution treatment was at1,050°C for 15 min, followed by air cooling. Agingwas at 480°C for 1 h, 510°C for 4 h, and 600°C for4 h. Optical metallographic studies were made afteretching with Fry’s reagent (5 g [0.176 oz] cupric chlo-ride + 40 mL [1.353 oz] hydrochloric acid (HCl) +30 mL [1.015 oz] water + 25 mL [0.845 oz] methylalcohol).

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(1) UNS numbers are listed in Metals and Alloys in the UnifiedNumbering System, published by the Society of AutomotiveEngineers (SAE) and cosponsored by ASTM.

Transmission electron microscopy (TEM) alsowas used for microstructural studies after electro-chemical thinning in a 5% perchloric acid and aceticacid mixture (maintained at –5°C) using a double-jetpolishing unit. A Vickers hardness survey also wasmade. For environmental cracking studies, U-bendspecimens were made according to ASTM G 30-79(Figure 1).8

The test environment was 3.5% sodium chloride(NaCl) solution (pH adjusted to 2.0 by the addition of2 mL [0.068 oz] to 2.5 mL [0.085 oz] concentrated

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FIGURE 3. d-ferrite in martensitic matrix in the weld metal (as-welded) (5,000x).

FIGURE 2. Schematic of the environmental cracking test setup: (1) U-bend sample (working electrode), (2) level of 3.5%NaCl solution, (3) salt bridge, (4) SCE reference electrode, (5) graphite rod (counter electrode), (6) potentiostat, (7) insulatedcopper wire, (8) stand, and (9) glass beakers.

HCl). Figure 2 is a schematic of the test setup. Thepotentials (from–1,200 mVSCE to 200 mVSCE in steps of 100 mV, wereapplied using a Wenking Model TS 1† potentiostatwith reference to a saturated calomel electrode (SCE).Tests were done at open-circuit potential (OCP) also.Only the strained half-circular portion of the U-bendsample was immersed into the electrolyte.

The time to final fracture was noted for eachapplied potential. The specimens were subjected tothe applied potential until fracture or until 500 h,whichever occurred first. If no fracture or crack ini-

500

† Trade name.

tiation occurred within 500 h, the specimen wastaken out and considered to be a no-failure speci-men. Three specimens were tested for each condition,and an average failure time was reported. Crack mor-phology studies of failed specimens were done usingboth optical microscopy (OM, on polished and etchedspecimens) and scanning electron microscopy (SEM).

RESULTS

MicrostructureThe as-welded mitrostructure contained a

d-ferrite network in the martensitic matrix. d-ferritewas solidified at the interdendrites. TEM observationof the as-welded sample shown in Figure 3 indicatedthe presence of d-ferrite.

No change was noted in the opticalmetallographs of weld metals aged directly (withoutsolution treatment) at 480°C or 510°C compared totheir as-welded counterparts. However, precipitationwas observed during TEM for all the heat treatmentsinvolving aging. The d-ferrite network was not seendistinctly, as in the case of the as-welded condition.This may have been a result of overaging of the mar-tensite. At this temperature, a higher degree ofcarbide precipitation was found to occur. d-ferrite-martensite interfaces were found to be preferentialsites for the carbides to precipitate. TEM observationclearly showed the same findings.

When the weld metal was solution treated, thedendrites vanished. The d-ferrite network that wasobserved in the as-welded condition mostly was dis-solved after solution treatment at 1,050°C. A smallamount of d-ferrite was found as globules.

Aging the weld metal at 480°C and at 510°C aftersolution treatment did not change the microstructure

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TABLE 1Failure Times vs Applied Potentials

of 17-4 PH SS Weldments with Directly Aged Conditions

AppliedPotential As-Welded 480 °C 510°C 600°C(mVSCE) (Not aged) (1 h) (4 h) (4 h)

+200 2.0 4.0 – ––100 5.0 5.0 3.0 1.5–200 33.4 6.7 4.0 3.2–300 – 15.0 5.8 4.5OCP 60.0 23.5 35.0 86.6–500 44.0 10.0 13.3 57.0–600 16.7 6.2 10.0 35.0–700 20.0 5.25 12.0 60.0–800 11.5 4.0 11.7 50.0–900 – 2.0 8.5 21.0

–1,000 9.25 1.8 8.0 25.5–1,200 7.2 1.0 2.0 8.3

Average Failure Time of Directly Aged Weldmentswith Directly Aged Conditions

compared to the solution-treated weld metal micro-structure. After solution treating + aging, the weldmetal microstructure was similar to that of the basemetal. The HAZ was etched very dark where higheramounts of carbides precipitated. Aging the HAZ at480°C and 510°C increased the amount of carbideprecipitation.

Solution treatment at 1,050°C caused all thecarbides and overaged precipitates to dissolve intothe solid solution. Upon air cooling, the microstruc-ture was uniformly martensitic. The as-welded metalVickers’ hardness number (VHN) was found to be340 VHN. Aging at 480°C for 1 h (peak aging) andaging at 510°C for 4 h increased the hardness to460 VHN and 420 VHN, respectively. Aging at 600°Cfor 4 h (overaging) was found to result in a hardnessvalue of 340 VHN. Even though solution treatment ofthe weld metal caused changes in microstructure,the hardness results were not affected. Both directlyaging the weld metal and solution treating + agingthe weld metal gave almost the same hardness value(~ 460 VHN).

Environmental Crackingof the As-Welded Samples

OCP — The times to failure for U-bend sampleswith various PWHT at open-circuit condition aregiven in Table 1. As-welded samples took ~ 60 hfor complete fracture at OCP. The failure occurred atthe HAZ, and no cracks were observed at the weldmetal. The fracture was completely intergranular.This could be attributed to selective dissolution ofgrain boundaries at the chromium-depleted regionsdue to grain-boundary carbide precipitation at theHAZ.

Cathodic Potentials — ␣ Figure 4 illustrates therelation between applied potentials and times to fail-


ure of the as-welded samples. Application of–1,200 mVSCE caused failures within 10 h. Final frac-ture occurred at the center of the weld metal, whichexhibited a transgranular mode of cracking (Figure5). Secondary cracks also were observed on the frac-ture surface. Solute segregation at the interface ofthe two meeting dendrites could have caused hydro-gen ingress, which could have led to decohesionbetween the dendrites and caused secondary cracks.

The application of lower cathodic potentials in-creased times to failure in the weld metal.Application of more negative cathodic potentialscaused failures at the weld metal, and less negativecathodic potentials caused failures at the HAZ.

Anodic Potentials — Application of anodic poten-tials caused faster failures when compared to failuresat OCP (Figure 4). Under anodic potentials, failureoccurred at the HAZ only and was found to be inter-granular in nature. At anodic potentials, multiplecracks were observed in the HAZ of the U-bend weldsamples. Application of potentials nobler (more posi-tive) than –100 mVSCE caused failure by generalcorrosion.

Welded + Directly Peak-Aged SamplesOCP — Table 1 shows the cracking test results of

directly peak-aged welded specimens. Aging the weldmetal samples at 480°C for 1 h caused quicker fail-ures at open-circuit condition. These samples failedat the end of the HAZ. No secondary cracks wereobserved on the failed surface. The samples aged at480°C for 1 h took ~ 24 h, and the samples aged at510°C for 4 h took ~ 35 h for complete fracture.

Cathodic Potentials — Figure 6 shows the effectof applied potentials on times to failure of the U-bendsamples aged at 480°C for 1 h and at 510°C for 4 h.Application of –1,200 mVSCE to the weld samples aged

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FIGURE 4. Failure time vs applied potential of the as-welded samples.

FIGURE 5. Transgranular cracking in weld metal (as-welded) failedat applied potential of –1,200 mVSCE (200x).

at 480°C for 1 h caused very quick failure (within1 h). The samples failed at the center of the weldmetal. Secondary cracks were seen in the HAZ also.Increases in applied cathodic potentials in the noble(less negative) direction increased the times of failureas in other cases. At cathodic potentials, fractureoccurred predominantly in the weld metal.

Application of –1,200 mVSCE to the weldedsamples aged at 510°C for 4 h caused complete frac-ture in ~ 2 h. Compared to peak-aged (480°C for 1 h)weld metal, the samples aged at 510°C for 4 hshowed a little higher resistance to cracking at morecathodic potentials. The mode was transgranular. Atall cathodic potentials (Figure 6), the samples aged at510°C for 4 h showed better resistance to cracking.In these samples also, the fracture occurred mainlyat the weld metal from application of cathodic poten-tials.

Anodic Potentials — Figure 6 illustrates failuretimes under anodic potentials. Application of anodic

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potentials caused quicker failures in both the agedconditions (480°C for 1 h and 510°C for 4 h). How-ever, the samples aged at 510°C for 4 h failed fasterthan the samples aged at 480°C for 1 h. This con-trasted with results for cathodic potentials. Anodicpotentials mainly resulted in HAZ failure. Applicationof potentials more noble than –100 mVSCE causedfailure by general corrosion in both aged conditions,where the sample thinned down gradually and brokeinto two pieces.

Welded + Directly OveragedOCP — The U-bend weld samples that were

overaged at 600°C for 4 h failed at OCP after 85 h.When compared to the as-welded or peak-aged condi-tions, the overaged samples had longer failure timesat OCP. In this case, failures also took place only atthe HAZ.

Cathodic Potentials — Figure 7 illustrates therelation between applied potentials and times to fail-ure of overaged weld metal samples. As was evidentfrom Table 1, at cathodic potentials, overagedsamples exhibited better resistance to cracking thantheir as-welded or directly peak-aged counterparts.Application of –1,200 mVSCE caused failures at ~ 9 h.Even though the failures occurred quickly, the timeto failure was greater in this case than for that of theas-welded or peak-aged weldments. In general, in-creasing the cathodic potentials toward the nobleside (less negative potentials) increased the times tofailure. However, there were some anomalous resultsat some potentials. As in the previous cases, theoveraged samples also failed at weld metal regions(mostly at the center of the weld metal) from the ap-plication of cathodic potentials. Figure 8 shows thetransgranular mode of fracture.

Anodic Potentials — Figure 7 illustrates the fail-ure times under applied anodic potentials of overagedsamples. Application of anodic potentials drasticallyreduced the times to failure of overaged weldsamples. When compared to as-welded or peak-agedweld metal samples, the overaged weld samples failedsooner. Application of –200 mVSCE, caused failure bygeneral corrosion, where the sample thinned downgradually and broke into two pieces. In the as-weldedor peak-aged samples, general corrosion occurredonly at potentials less negative to –100 mVSCE. Asfound in other cases, overaged samples also failed atthe HAZ from the application of anodic potentials.

Solution-Treated + Peak-Aged WeldmentsOCP — The weld samples solution treated and

peak aged at 480°C for 1 h failed at ~ 33 h at OCP. Incontrast, the weldment that was solution treated butnot aged did not fail at OCP until after 500 h. Sincesolution treatment caused the disappearance of thedendritic weld structure and the formation of mar-tensite on the equiaxed grains, the cracking behavior


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FIGURE 7. Failure time vs applied potential of directly overagedweldments.

FIGURE 6. Failure time vs applied potential of directly peak-agedweldments.

FIGURE 8. Transgranular fracture in directly overaged weld metalfailed at –1,000 mVSCE.

could be compared to that of the base metal, whichhad a similar microstructure. It was interesting tonote that both peak-aged base metal samples8 andweld samples after solution treating + peak agingexhibited similar failure times at OCP.

The optical microstructure of the same sampleillustrated the secondary cracking at the weld metaland nearby base metal. The cracks at the weld metalwere very sharp and discontinuous. The crackingpredominantly occurred at the fusion line and in thenearby base metal. It was found that the directlypeak-aged weld metal region (which was peak-agedwithout solution treatment) did not contain anycracks when tested at OCP. However, solution-treated and peak-aged weld metal had multiple sharpcracks even though the failure took place at thenearby fusion line/base metal region.

Cathodic Potentials — Figure 9 illustrates theeffect of applied potentials on failure times for solu-tion-treated + peak-aged weld samples. Application of–1,200 mVSCE caused failure within 1 h. As the poten-tials moved to less negative values, the failure timesalso increased. The samples failed at the weld metalregion at all the cathodic potentials. The solution-treated and peak-aged weld samples that failedshowed heavy crack branching.

Anodic Potentials — Figure 9 shows how applica-tion of anodic potentials reduced the times to failure.Breaking of the material into two pieces was observedat less noble potentials because of general corrosion.Anodic potentials caused severe cracking throughoutthe immersed portion of the samples. The crackswere not confined to either the HAZ (as in the case ofdirectly aged samples) or the weld metal region, Thepeak-aged base metal samples and the solution-treated + peak-aged weld samples behaved similarlyat anodic potentials.

Solution-Treated + Overaged WeldmentsThe solution-treated + overaged samples failed at

~ 420 h for the open-circuit condition, When com-


pared with overaged base metal, the solution-treated+ overaged weld metal was more prone to cracking.But when compared to that of directly overagedweldments without solution treatment, the resistanceto cracking improved many times over at OCP and atboth anodic and cathodic applied potentials. At–1,200 mVSCE, the solution-treated + overagedweldment showed very high resistance to cracking,better than that of the overaged base metal. As theapplied cathodic potentials moved in the noble (lessnegative) direction, failure times also increased. Atanodic potentials, the failure times decreased whencompared to that at the OCP. However, the resis-tance to fracture was better than any other PWHTcondition.

Figure 10 shows the optical microstructure ofthe solution-treated + overaged weld sample thatfailed at OCP. Severe intergranular corrosion wasevident. Failure occurred mainly because of attack

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FIGURE 9. Failure time vs applied potential of solution-treated andaged samples.

FIGURE 10. Intergranular cracking in the HAZ of solution-treated +overaged weld failed at OCP (200x).

on grain boundaries, which caused selective dissolu-tion rather than SCC. Even though multiple crackswere seen throughout the exposed portion of thesample, final fracture occurred at the base metal(HAZ) region.

The sample that failed at –300 mVSCE exhibitedsevere grain boundary attack. The final fractureoccurred at the weld metal. For anodic potentials, thefinal failures generally took place at the HAZ for thesamples that were not solution treated. But aftersolution treatment, there was no preferential crack-ing at the HAZ.

DISCUSSION

Table 2 summarizes the overall observations ofenvironmental cracking of the weld samples. Thecomposite nature of the weld samples did not followthe relationship between strength and environmental

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cracking resistance. The hardness values were notuniform throughout the weldment (weld + HAZ +base). The weld samples failed at the HAZ underopen-circuit conditions in all the PWHT conditions(without solution treatment). For example, the trans-verse-weld, peak-aged (480°C for 1 h withoutsolution treatment) U-bend sample failed at the HAZat OCP.

If the relationship between strength and suscep-tibility were followed, the sample should have failedat the weld metal, where the hardness was higherthan at the HAZ. Both weld and HAZ containeddifferent microstructures. This difference in micro-structure probably caused different rest potentials(OCP) at the weld metal and the HAZ. The rest poten-tial or OCP of the HAZ was found to be more activethan the weld metal because of carbide precipitationand the resultant chromium-depleted regions. Thiscaused preferential corrosive attack at the HAZ atOCP under stress when exposed to 3.5% NaCl solu-tion (pH = 2.0). It could be said that the weld metalwas protected cathodically by the HAZ, which madethe weld metal free of cracks. Failure occurred at theHAZ mostly by a single crack, predominantly byanodic path dissolution of grain boundaries. Underanodic potentials, fracture also occurred at the HAZ,which could be attributed to preferential attack atgrain boundaries due to carbide (such as M23C6) pre-cipitation.

If the difference between the applied cathodicpotentials and the individual OCP were considered,the weld metal regions had higher cathodic polariza-tion than their HAZ counterparts. This higher degreeof cathodic polarization along with the higher hard-ness of weld metal could be the reason for theincreased susceptibility to hydrogen cracking at morenegative cathodic potentials.

When compared to base metal, the weld metalthat was directly aged (without solution treatment)exhibited lower resistance to cracking at all poten-tials.9 Further, the as-welded samples took relativelyless time to fail at all potentials when compared totheir solution-treated counterparts. It was more truefor cathodic potentials, where the fracture occurredat the weld metal regions.

Local compositional variations due to segregationduring solidification of weld metal also would haveaffected the cracking behavior. This could be ob-served from the solidification structures of thefracture surfaces of the weld samples. Moreover,segregation of alloying elements such as copperwould have affected the aging characteristics of theweld metal. Crack nuclei were reported to haveformed by separation at the interface of the precipi-tated phases or segregated elements from theplastically deformed matrix under cathodic poten-tials.10 The crack could nucleate at the d-ferrite-martensite interfaces also. The d-ferrite morphology


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TABLE 2Summary and Grading of Environmental Cracking Behavior of 17-4 PH SS Weldments

17-4 PH SS Failure Fracture (B) More Less Failure Fracture Failure FractureWeldments Rating (A) Location Mode Cathodic Cathodic Location Mode Rating Location Mode

As-welded C HAZ I D C Weld T C HAZ I(more)HAZ(less)

PWHT: Directly agedPeak-aged (480°C, 1 h) D HAZ I P E Weld T D HAZ IPeak-aged (510°C, 4 h) C HAZ I P D Weld T D HAZ I

Overaged Weld TOveraged (600°C, 4 h) B HAZ I C C (more) I E HAZ I

HAZ(less)

PWHT:Solution Treated + AgedPeak-aged (480°C, 1 h) D F/L I P E Weld + I D Base I

Base basemetal

Overaged (600°C, 4 h) S Base I A A Weld I E Weld/ Imetal base

(A) Relative gradings: S = excellent (10 points), A = very good (9 points), B = good (8 points), C = mediocre (7 points), D = poor (6 points), E = verypoor (5 points), and P = worst (0 points).

(B) Fracture mode: T = transgranular, I = intergranular.

AnodicRatingOCP

Cathodic

in weld metal was different from that in the basemetal. In weld metal, it was found to solidify at thedendrite interspaces. Segregation of solute atoms atthe d-ferrite martensite interfaces (in the case of as-welded metal or precipitation of carbides in the caseof the aged weld metals) could have promoted hydro-gen segregation and reduced the cohesive strength ofthe lattice, which might have led to fracture. d-ferritein the base metal was considered to be disadvanta-geous from the fracture toughness and fatigue crackgrowth points of view because the d-ferrite-marten-site interfaces act as preferential sites for voidnucleation and growth.11 As the weld metal was agedat higher temperatures, some amount of d-ferrite wasfound to dissolve, and the morphology of the d-ferritechanged.

For weldments, the microstructural aspectswere important in determining the susceptibility toenvironmentally assisted cracking. The hardnessfactor alone could not predict the cracking behaviorat all potentials. For example, comparing the basemetal9 and weldment samples identically overaged at600°C for 4 h for an applied cathodic potential of–1,200 mVSCE, it was observed that the base metalshowed higher resistance than the weldment sample.The weldment sample failed at the weld metal region.The cast, dendritic structure of the weld metal, which


would have experienced more microcracks and weakspots, may have been the reason for the inferior re-sistance of the weldment at more cathodic potentials.

Overall, the following observations could bedrawn in the case of weldments. Under applied ca-thodic potentials, the weld metal region was the weakspot. Under OCP and applied anodic potentials, theHAZ region was weak. As in the case of base metals,the failures under cathodic potentials for weldmentscould be related to hardness to some extent. UnderOCP and anodic potentials, since cracking was foundto occur by an active path dissolution mechanism,the level of carbides and other intermetallic precipita-tion determined the susceptibility to cracking.

For better resistance under cathodic potentials,solution treating + overaging was found to be thebest treatment. Directly overaging gave relativelybetter results under cathodic potentials. But, thesetwo treatments were poor when the resistance toapplied anodic potentials was considered. Both di-rectly peak-aging the weldments and solutiontreating + peak aging were found to behave similar tothe peak-aged base metal.9 They showed the leastresistance to cracking under cathodic conditions andintermediate resistance under anodic conditions. Theas-welded samples, whose preweld condition wassolution-treated, showed intermediate resistance

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under cathodic conditions and better resistance un-der anodic conditions when compared to the othertwo directly aged conditions. Solution-treated but notaged weldments behaved almost the same as thesolution-treated base metal.

CONCLUSIONS

❖ Irrespective of their thermal histories, samplescracked in the HAZ at the OCP and anodic potentials.This was found to be due to the active path dissolu-tion mechanism in the HAZ.❖ Though the weld metal hardness was more thanthe HAZ, no cracking took place in the weld metal.❖ Solution treating the welds improved their crackingresistance at OCP and anodic potentials. In contrast,at applied cathodic potentials, the failure was withinthe weld metal.❖ The hardness criterion (the harder the structure,the more the susceptibility to cracking) was found tobe applicable mainly for applied cathodic potentials.Peak aging resulted in the least cracking resistanceunder these potentials.❖ To increase the cracking resistance under cathodicpotentials, solution treating (1,050°C) followed by

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overaging (at 600° C for 4 h) was found to be the besttreatment. However, the same treatment was foundto be highly detrimental for anodic potentials andOCP.

REFERENCES

1. T.G. Gooch, Weld. J. 53 (1974): p. 287s-298s.2. Aerospace Structural Metals Handbook, vol.12 (Mechanical,

Properties Data Centre, Batelle Columbus Laboratories, 1980).3. Welding Handbook, vol. 4, 7th ed. (Miami, FL: American Weld-

ing Society), p. 129.4. C.R. Weymuller, Weld. Design Fabrication May (1980): p. 68.5. K.S. Raja, R.K. Prasad, “Effect of Thermal Treatments on Elec-

trochemical Behavior of Two PH Stainless Steels,” ed. A.D.Tiruchirappalli (Tamil, India: Indian Institute of Welding, 1991),p. 419-428.

6. K.S. Raja, R.K. Prasad, “Effect of Thermal Treatments on Me-chanical and Corrosion Properties of PH 15-7 Mo (Cu) StainlessSteel Weldments,” Proc. Natl. Welding Seminar, Madras, India,October 1991 (Tamil, India: Indian Institute of Welding, 1991),p. 58-70.

7. J.E. Truman, Int. Met. Rev. 26 (1981): p. 301-349.8. ASTM G 30-79, “Standard Recommended Practice for Making

and Using U-Bend Stress Corrosion Test Specimens,” in AnnualBook of Standards (Philadelphia, PA: ASTM, 1990).

9. K.S. Raja, “Effect of Thermal Treatments on Corrosion Behaviorof Type 316, 17-4 PH, and PH 15-7 Mo (Cu) Stainless SteelWelds (Ph.D. diss., Indian Institute of Technology, 1993).

10. G. Burnell, D. Hardie, R.N. Parkins, Brit. Corros. J. 22 (1987):p. 229-237.

11. H.J. Rack, D. Kalish, Metall. Trans. 5 (1974): p. 595-1,605.


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