Influence of welding process on Type IV cracking behavior of P91 steel

M. Divya a, C.R. Das a,n, S.K. Albert a, Sunil Goyal a, P. Ganesh b, R. Kaul b, J. Swaminathan c,B.S. Murty d, L.M. Kukreja b, A.K. Bhaduri a

a Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, Indiab Laser Materials Processing Division, Raja Ramanna Centre for Advanced Technology, Indore 452013, Indiac National Metallurgical Laboratory, Jamshedpur 831007, Indiad Department of Metallurgical and Materials Engineering, Indian Institute of Technology Madras, Chennai 600036, India

a r t i c l e i n f o

Article history:Received 15 October 2013Received in revised form24 June 2014Accepted 24 June 2014Available online 1 July 2014

Keywords:Laser weldingShielded metal arc weldingP91 steelType IV crackingLath martensiteFinite element analysis

a b s t r a c t

Influence of laser welding (LW) and shielded metal arc welding (SMAW) processes on Type IV crackingbehavior of modified 9Cr–1Mo (P91) steel has been investigated in this paper. The study involvedcomparison of stress rupture lives of modified 9Cr–1Mo steel weldments prepared by SMAW andcontinuous wave CO2 laser welding processes. Width of the heat affected zone (HAZ) in laser weldmentwas found to be 1.0 mm, whereas it was �2.5 mm in SMAW weldment. The rupture lives of laserweldment were found to be higher than SMAWweldment at higher stress level and comparable at lowerstress level. Under similar stress levels, the creep rupture lives of 875 1C heat treated specimens werefound to be lower than that of the base metal and cross weld specimen. These results clearly suggest thatthe instability of microstructure in the intercritical heat affected zone (ICHAZ) is responsible for lowercreep rupture lives of P91 steel weldment than the base metal. The experimentally observed variationsin creep cavitation have been corroborated with the results of finite element (FE) analysis.

& 2014 Elsevier B.V. All rights reserved.

1. Introduction

Creep strength of the high chromium ferritic steels weldmentsis inferior to that of the base metal. These weldments fail in theweakest regions of the heat affected zone (HAZ), which issandwiched between the base metal on one side and the weldmetal on the other side. Creep cavitation in these regions becomesthe life limiting factor for the welded components during hightemperature service. Failure of the weldment in these regionsunder creep conditions is known as “Type IV” cracking. Type IVcracking is associated with the microstructural degeneration in theHAZ of the weldment during welding and long-term serviceexposure [1–8]. The microstructure in the HAZ is influenced bythe peak temperatures experienced during weld thermal cycles.Based on these peak temperatures, the phase transformation occurlocally which causes variations in microstructure in the HAZ.Therefore, the HAZ can be divided into three zones, viz., coarsegrain (CGHAZ), fine grain (FGHAZ) and intercritical (ICHAZ) heataffected zone. Among these zones, FGHAZ and ICHAZ are theweakest regions. These regions experience peak temperatures forextremely short duration during welding which are expected to be

just above Ac3 (FGHAZ) and between Ac1 and Ac3 (ICHAZ) tem-peratures. This causes solid state phase transformation and thetransformation products are distinctly different than those of thebase metal and weld metal. Lath martensite forms in the FGHAZhave less carbon than those forms in the CGHAZ. In addition tothat, the precipitates do not dissolve completely in this region. As aresult, the strength of this region is less compared to the CGHAZ.But in the ICHAZ, the partial transformation of tempered lathmartensite to austenite as well as partial dissolution of precipitatestakes place during heating cycle which results in formation of lowcarbon martensite during cooling. During post weld heat treat-ment (PWHT) less precipitates form in these regions comparedto the CGHAZ. However, the coarsening of partially dissolvedprecipitates occurs rapidly resulting in the drop in hardness andstrength in the FGHAZ and ICHAZ. Higher strength of CGHAZ/weldmetal and the base metal than the weakest regions imposes amicrostructural constraint to the deformation behavior in theseregions which leads to the development of multi-axial loadingconditions under creep condition [1–5]. This constraint effect leadsto the creep cavitation induced intergranular failure during long-term service. To mitigate the “Type IV” cracking in 9Cr steels,advanced 9–12Cr ferritic steels have been developed in recentyears [6–8]. Stable microstructure of these steels at high tempera-ture improves the creep strength. On the other hand, use ofadvanced welding processes, like laser and electron beam welding

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Materials Science & Engineering A

http://dx.doi.org/10.1016/j.msea.2014.06.0890921-5093/& 2014 Elsevier B.V. All rights reserved.

n Corresponding author. Tel.: þ91 44 27480118.E-mail address: chitta@igcar.gov.in (C.R. Das).

Materials Science & Engineering A 613 (2014) 148–158

for joining the steels, reduces the width of the HAZ significantly.This is due to the localized heating/cooling during thermal cyclingin advanced welding process compared to SMAW or TIG process.This has a beneficial effect in improving the creep life of weld-ments [9].

Attempts have been made to estimate the stress and straindistributions developed across the weldment and its consequenceson type IV creep cavitation [5,10–13] using finite element (FE)analysis. Hayhurst et al. have extensively used physically basedcontinuum damage mechanics (CDM) coupled with FE analysis toestimate the creep damage evolution and failure life of the butt-welded components [14–17]. Hall and Hayhurst [16] and Perrinand Hayhurst [17] had developed creep constitutive equations ofweldments considering the physical processes of dislocation sub-structural evolution, carbide coarsening and creep cavity nuclea-tion and growth under multiaxial state of stress. These creepconstitutive equations were incorporated in the finite elementmodel to study the stress and strain distribution, creep damageevolution and failure life of the weldments [14–17].

Though, there are a large number of publications on creepdeformation of SMAW, TIG and submerged arc weld joints, similaremphasize is not given to the effect of welding techniques oncreep deformation behavior of weldment of P91 steels prepared byLW and SMAW processes. Therefore, the objective of the presentstudy is to understand the creep deformation behavior of P91 steelweldments, prepared using LW and SMAW processes and under-standing there deformation behavior using FE analysis.

2. Experimental details

Normalized and tempered (N&T) P91 steel was received in tubeand plate forms. Normalizing and tempering temperatures were1050 1C for 30 min and 760 1C for 60 min, respectively. Plates of250�150�12 mm3 were used for preparation of SMAW weld-ments and tubes having outer diameter of 34 mm and wallthickness of 2.5 mm were used for laser weld joints fabrication.E9016 B9 electrodes were used for preparation of the weld padsusing SMAW process. Preheating and interpass temperatures weremaintained in the range of 200–250 1C. Welding parameters usedfor preparation of weld pads are given in Table 1. In order toremove the diffusible hydrogen, weld pads were subjected to postweld heating at 250 1C for 30 min.

LW was performed with indigenously developed 5 kW contin-uous wave (CW) CO2 laser machine. The raw laser beam, emanat-ing from the laser system was folded by a 451 plane mirror andsubsequently focused with a 100 mm focal length ZnSe lens.Focusing lens was mounted on a water-cooled copper nozzle.Autogenous laser welding was performed by rotating the tubewhich was to be joined with the focused laser beam. During thecourse of laser welding, argon gas was passed through the nozzlethat served the dual purpose of shielding the materials to bewelded as well as protecting the ZnSe lens from damage due topossible particulate emissions from the material. Experimentalparameters used for laser welding of P91 tubes were laser power:4.15 kW; specimen rotation speed: 12 rpm (corresponding linearspeed¼21.4 mm/s). Weldments were further subjected to

radiographic examination and found to be free from defects. Theseweldments were subjected to PWHT at 760 1C for 3 h in a mufflefurnace.

For microstructural observations, the weldment with thedimensions of 10�5�3 mm3 and 10�5�2.5 mm3 were takenfrom both the weld pads and then polished metallographically upto diamond (1 mm) finish followed by chemical etching usingVillella's reagent. The microstructures of etched specimens wereobserved under an optical electron microscope and a scanningelectron microscope (SEM). These specimens were also used forhardness measurements across the weld interface. The Vickersmicrohardness tests were carried out using a Shimadzu HMV2000machine with 50 and 500g load. Creep tests on SMAW weldmentswere carried out at stress levels of 80, 100 and 120 MPa onspecimens having 50 mm gauge length, 6 mm gauge diameterand total specimen length of 100 mm. On the other hand, thecreep test on LW specimens was carried out having the dimen-sions of 25 mm gauge length, 5.5 mmwidth and 2.5 mm thicknesswith total length of 60 mm. To compare the results, the creep testsof base metal were also carried out under similar testingconditions.

In order to understand the microstructural features of ICHAZand CGHAZ, the base metal was heat-treated at temperatures of8751 and 1175 1C for 5 min, in a muffle furnace, which areexperienced in the HAZ of the weldment during welding. Thesetemperatures were chosen considering the transformation start(Ac1) and transformation finish (Ac3) temperature values for thissteel, which are 836 1C and 900 1C respectively, at a heating rate of50 1C min�1. Though the transformation temperatures increasewith heating rate, the equilibrium transformation temperaturewould be lower than that of the non-equilibrium transformationtemperatures. It may be noted that the heating and cooling ratesexperienced during welding are higher substantially than that isregularly used in a furnace. Therefore, the heating rate commonlyused in a furnace cannot simulate the actual heating ratesexperienced in arc welding. Gleeble thermal cycle is commonlyused for microstructural simulation. However it does not produceuniform microstructure with sufficient dimensions and so thestandard tensile creep tests cannot be carried out. Consideringthis limitation associated with actual microstructure simulation inGleeble, the furnace heating has been considered in the presentinvestigation. Specimens heat-treated at 875 1C were subjected tocreep under similar testing conditions for base metal and weld-ments. Though Type IV cracking is reported to occur in the FGHAZor ICHAZ of the weldment, for easiness ICHAZ was considered inthis study. Electron Back Scattered Diffraction (EBSD) and trans-mission electron microscopic techniques were used for the micro-structural characterizations.

To understand the creep deformation and fracture behavior ofweldment, analysis of the stress and strain distribution across itduring creep exposure was carried out by FE analysis. Norton'screep law relating the steady state creep rate with applied stress(_εs ¼ Aσn, where _εs is the strain rate (h�1), σ is the stress (MPa) andA and n are the constants) was used as a constitutive equation.The FE simulations were carried out using ABAQUS 6.10 commer-cial software. Specimens were modeled using 3D 8 noded brickelements for cylindrical specimens and C3D8R elements for platetype specimens. Half of the specimen length was modeled inanalysis because of geometrical and loading symmetry. In order tounderstand the effect of HAZ width on triaxiality, FE analysis wascarried out on cylindrical specimens considering width of theICHAZ of 0.5 and 1.5 mm. Young's modulus and Poison's ratio ofthe different constituents of the ferritic steel weldment werechosen as 109 GPa and 0.3 respectively [18]. Analysis was carriedout at a stress of 100 MPa up to 5000 h of creep exposure as thestress distribution saturated beyond 5000 h.

Table 1Welding parameters used for welding of P91 steel plate using SMAW process.

Welding process Current (A) Voltage (V) Heat input (kJ/mm)

SMAW 80–90 12 1.0

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Table 2Chemical composition of plate and tube in wt%.

Elements C Cr Mo Si Mn V S P Ni Al N Nb Fe

Plate 0.100 9.500 1.00 0.480 0.390 0.250 0.009 0.021 0.14 0.024 0.065 0.100 Bal.Tube 0.09 9.00 0.89 0.25 0.40 0.20 0.008 0.020 0.15 0.02 0.050 0.065 Bal.SMAW weld metal 0.100 9.00 1.00 0.24 0.700 0.170 0.012 0.009 0.500 0.03 0.055 0.06

Fig. 1. (a) Secondary electron image of N&T P91 base metal, (b) bright field TEM image of N&T P91 steel base metal, (c) EDS spectra obtained from a large precipitate markedwith a black arrow and (d) EDS spectra obtained from a smaller precipitate marked with a white arrow.

Fig. 2. Macrostructure of weldments of P91 steel prepared using (a) SMAW and (b) laser welding processes.

Fig. 3. Secondary electron image of weld metal (a) SMA weld metal and (b) laser weld metal.

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3. Results and discussion

3.1. Microstructure of base metal and weld metal

Chemical composition (wt%) of tube and plate is given inTable 2. Secondary electron image of base metal in temperedcondition is shown in Fig. 1(a). The microstructure of P91 steelconsists of tempered lath martensite, where M23C6 precipitates arepresent along the lath, packet, block and prior austenite grainboundaries. For these steels, two types of precipitates are reportedin literature [5,19]. Bright field transmission electron microscopic(TEM) image of normalized and tempered steel is shown in Fig. 1(b).The EDS spectra obtained from the bigger and fine precipitatelocated at the boundaries and those inside the lath martensite are

shown in Fig. 1(c) and (d) respectively. EDS analysis confirmed thatthe bigger precipitates present at the boundaries were enriched inchromium, iron, molybdenum and carbon and its stoichiometry issimilar to that of M23C6 precipitate [20,21]. On the other hand, thefine precipitates present inside the lath martensite were carboni-trides either rich in Nb or V. These observations are in agreementwith those reported in literature [20,21]. Niobium and vanadiumrich carbides, nitrides and carbo-nitrides (MX) are reported toimprove creep strength of this steel [21].

Macrostructure of SMA and laser weldments is shown in Fig. 2(a)and (b). It is clearly evident from the macrographs that the widthof the HAZ is higher in SMA weldment than that of the laserweldment. Microstructures of SMA and laser weld metals areshown in Fig. 3(a) and (b) respectively. From these secondary

Fig. 4. Microstructures of HAZs of SMAW weldment (a) CGHAZ, (b) FGHAZ and (c) ICHAZ; microstructures of HAZs of laser weldment (d) CGHAZ, (e) FGHAZ and (f) ICHAZ.(For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)

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electron images it is clearly evident that precipitates are coarser inSMA weld metal than that of the laser weld metal. This differenceis attributed to single and multiple thermal cycles experienced inlaser and SMA weld metals respectively.

3.1.1. Microstructure of weldmentSEM photomicrographs of HAZ for SMA weldment taken at a

distance of 0.5, 1.0 and 1.8 mm from the weld interface are shownin Fig. 4(a)–(c). Similarly, SEM photomicrographs for laser weld-ment are shown in Fig. 4(d)–(f) which are taken at distances of 0.1,0.3 and 0.75 mm from the weld interface. Generally, thermalcycles experienced in the HAZ are characterized by the peaktemperatures which lead to the phase transformation in HAZ.The prior austenite grains observed at a distance of 0.5 mm fromthe weld interface are relatively coarser as the peak temperature ismuch higher than Ac3 transformation temperature. This causescomplete dissolution of precipitates resulting in increased carboncontent in solid solution. As a result marteniste formed in thisregion has higher tetragonality and prominent straight edges asmarked in Fig. 4(a) [3].

Peak temperatures experienced at 1.0 mm and 0.3 mm dis-tances from the weld interface of SMA and laser weldments werejust above the Ac3 temperature. In this temperature range, chro-mium rich carbides do not dissolve completely on heating duringwelding [22,23]. These undissolved precipitates act as pinningsites and restrict the coarsening of austenite grains at lowertemperature in this region.

Therefore, fine prior austenite grains (3–8 μm) were observedin this region, which is known as FGHAZ [22]. On the other hand,the peak temperature experienced at 1.8 mm and 0.75 mm dis-tances from the weld interface of SMA and laser weldments wasjust above the Ac1 temperature but below the Ac3 temperature.Thus, very fine austenite grains (1–2 μm) nucleate along the prioraustenite grain boundaries, lath, block and packet boundariesduring heating. This region is known as ICHAZ [3,8]. Austeniteformed in this region is lean in alloying elements, precisely carbon,is well reported in literature [3]. Consequently the martensiteformed does not have a typical lath morphology due to low c/aratio [2,3,5]. Presence of very fine grains is indicated by red arrowsin Fig. 4(c) and (f). From these figures, it is clearly evident thatsubstructure is completely developed in ICHAZ of SMA weldmentas compared to laser weldment. It is worth noting that the dwelltime at a particular peak temperature in the HAZ is smaller for the

laser weldment than that of SMA weldment. Accordingly thewidth of the HAZ is smaller in the laser weldment as observed,Fig. 2(b). As TEM specimen preparation is difficult from actualweldment, an EBSD study was carried out for SMA weldment asdifferent regions in this HAZ of the weldment were much biggerthan that of the laser weldment.

The inverse pole figure (IPF) maps obtained from CGAHZ andICHAZ of the weldment are shown in Fig. 5(a) and (b). The prioraustenite grain boundaries are prominent in Fig. 5(a) unlike inFig. 5(b) where refined/recovered grains are observed. The bigdark regions were observed at the grain boundaries in IPF map(Fig. 5(b)). This could be due to higher step size (0.5 mm) used forthe data acquisition which was bigger than the microstructuralfeatures. Regular grains observed in Fig. 5(b) as compared toacicular morphology in Fig. 5(a) clearly indicate that the lathmartensite is present in the CGHAZ. The features with similar colorin the map indicate the regions having the same orientation. Theplots of grain orientation spread (GOS) for CGHAZ and ICHAZ areshown in Fig. 6, which show marginal shift of GOS towards the leftfor ICHAZ as compared to CGHAZ. Cheong and Weiland reportedthat the GOS value decreased during the recrystallization in Alalloy indicating recovery/recrystallization [24]. Similar observa-tions were made in ICHAZ of the SMA weldment. This is inagreement with lower dislocation density in this region ascompared to CGHAZ [3,5]

3.1.2. Microstructure of simulated HAZBright field TEM images of P91 specimens which were sub-

jected to heat treatment at 1175 and 875 1C are shown in Fig. 7(a)

Fig. 5. IPF maps of (a) CGAHZ and (b) FGHAZ of SMAW weldment.

Fig. 6. Grain orientation spread for CGHAZ and ICHAZ of SMAW weldment.

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and (b). The morphology of the features observed in these imagesis in good agreement with that observed in SEM images (Fig. 4) aswell as IPF maps (Fig. 5). Lath martensite with fine precipitates atthe lath boundaries stacked parallely is clearly evident in Fig. 7(a).On the other hand, the specimen heat treated at 875 1C showedpresence of large precipitates at the boundaries and substructure(Fig. 6(b)). Presence of large precipitates inside the grain/substruc-ture suggests the coalescence of boundaries leaving the precipi-tates inside the grain [25].

3.2. Microhardness distribution

The hardness distributions across the weld interface of P91weldments of SMA and laser weldments after PWHT are shown inFig. 8(a) and (b). The hardness values of SMA and laser weldmetals were found to be �400 to 450 HV0.05 and 420–450 HV0.05

respectively. This marginal difference could be due to the differ-ence in cooling rate experienced in the respective weld metalswhich also resulted in marginal difference in microstructuralfeatures. The hardness values of the CGHAZ at 0.5 and 1.0 mmdistance from the weld interface of SMAW and laser weldmentswere 495 and 510 HV0.05 respectively. The hardness valuedecreased to 240 HV0.05 at 1.8 mm distance in the ICHAZ of theweldment prepared by SMAW process whereas the hardness valuewas 250 HV0.05 at 0.75 mm distance from the weld interface forlaser weldment. Lower hardness value observed in this region ascompared to CGHAZ is in agreement with that of the microstruc-tural features observed in respective regions as discussed inSection 3.1. The hardness values in all the regions of HAZdecreased after PWHT. It is important to note that the hardnessdip was not observed in laser weldment as was observed in SMA

weldment. This is in agreement with microstructural observationshown in Fig. 4(f). These hardness variations clearly demonstratethat the width of the HAZ is lower for laser weldment as observedin macrographs (Fig. 2).

3.3. Creep deformation

The comparison of creep curves of P91 base metal and heattreated specimens (875 1C/5 min) at a stress level of 120 MPa and600 1C is shown in Fig. 9. Commonly, creep deformation of thissteel is characterized by a negligible instantaneous strain due toinitial loading, a small transient creep strain or primary creep,where the creep rate decreases as a function of time followed by aminimum creep rate region and finally tertiary or acceleratedcreep region, where the creep rate increases. The accumulation ofcreep strain is high in tertiary creep as compared to that in

Fig. 7. Bright field TEM image of base metal heat treated at (a) 1175 1C and (b) 875 1C for 5 min.

Fig. 8. Distribution of hardness values across the weld interface of weldment prepared using (a) SMAW and (b) laser welding processes.

Fig. 9. Creep curves of base metal and simulated ICHAZ (875 1C/5 min) tested at600 1C and stress level of 120 MPa.

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transient and secondary creep regimes. It was observed that thesudden accumulation of creep strain occurred in all the specimensas a result of necking just prior to fracture (Fig. 9). Fig. 9 alsoreveals that the creep life of heat treated specimen (875 1C/5 min)is lower than that of the base metal under the same stress level.Fig. 10(a) and (b) shows the bright field TEM images, obtainedfrom the shoulder and the gauge region of the base metal speci-men, creep tested at 120 MPa and 600 1C. From these secondaryelectron images it is apparent that the recovery was less in theshoulder region than in the gauge region. This also suggests thatthe coarsening of precipitates was more in the gauge region thanthe shoulder region. This observation is in line with that reportedby Watanabe et al. [5].

Abe et al. correlated the decrease in creep rate with theformation of fine Laves phase in 9Cr steel [25,26]. Other research-ers reported that the precipitation of Laves phase results in thedecrease of solid solution strengthening in the vicinity therebydecreasing creep strength [26–29]. At the initial stage of Lavesphase precipitation, it impedes the movement of dislocationsresulting in lower creep rate. The recovery of dislocations withinthe lath boundaries and recombination of two boundaries canresult in the disappearance of lath boundaries, leaving a row ofM23C6 carbides inside the matrix. These two simultaneous metal-lurgical processes result in lath coarsening [25]. Reduction innumber of precipitates in the gauge section of the specimen ascompared to the shoulder region is clearly evident from the TEMimages shown in Fig. 10(a) and (b). EDS spectra taken from one ofthe precipitates at the substructure boundary, indicated by anarrow in Fig. 10(b), is shown in Fig. 10(c). Composition of this

precipitate in wt% is found to be Cr: 19.6, Fe: 47, Mo: 28.8, Ni:1.4 and Si: 1.9 which is rich in solid solution alloying elements andthis phase is known as the Laves phase.

The plots of variation in rupture life as a function of appliedstress are shown in Fig. 11. It is clearly evident that the rupture lifeof the base metal is higher than the SMA as well as the laserweldment. The rupture life of the heat treated ICHAZ specimen(875 1C/5 min) is found to be lowest among all specimens undersimilar test conditions. From Fig. 11 it is clear that the rupture lifeof laser weldment tested at 120 MPa is higher than the SMAweldment. But, with decreasing applied stress the differences inrupture lives diminished. It is believed that the triaxial stress statein the weldment, prepared by different welding processes, would

Fig. 10. Bright field TEM images of specimen creep tested at 120 MPa and 600 1C (a) shoulder and (b) gauge, and (c) EDS obtained from the precipitate marked with an arrowin (b).

Fig. 11. Applied stresses vs. rupture life plots for base metal, SMAW and laserweldments.

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be different as the width of the HAZ would vary depending on theprocess employed. The effect of width of ICHAZ on the creepdeformation is discussed in the next section. Higher rupture lifefor laser weldment as compared to the SMA weldment at a higher

stress level is attributed to the higher triaxiality in the laserweldment [30]. On investigating the increased rupture lives on9Cr–3W–3Co–V–Nb–B weldments made from laser as well aselectron beam (EB) welding process, Shaju et al. reported a

Fig. 12. Cavities in the (a) SMA weldment and (b) laser weldment tested at 120 MPa and 600 1C.

Fig. 13. The distribution of triaxiality factor across the round cross weld specimens considering the ICHAZ width of (a) 1.5 mm and (b) 0.5 mm.

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triaxiality factor as high as �3 in the FGHAZ for the EB weldmentsat the midsection, parallel to the stress axis [5]. These authorsshowed that the rate of decrease in creep strength was more forthe EB weldment having smaller HAZ width than that for the GTAweldment. Commonly, the failure of cross weld creep specimenoccurs in the weld metal or base metal by tensile deformation atrelatively higher stress levels whereas; by creep cavitation atrelatively lower stresses in the ICHAZ. This is further supportedby the presence of more creep cavities in the ICHAZ of creep testedcross weld specimen tested at lower stress levels than the onetested at higher stress level. Therefore, the presence of metallur-gical notch increases creep rupture life at higher stress levelswhere metallurgical instability does not take place in a short

period of time. Goyal et al. reported that the increased creeprupture life accompanied by the decrease in creep ductilityresulted from triaxiality introduced through the mechanical notch[31]. Compared to SMAweldments, metallurgical notch (the size ofICHAZ) in the laser weldments would be sharper, and due to highprinciple stress along with high hydrostatic stress in the ICHAZ,growth of cavities would be accelerated in the later [32]. As theprincipal stress and hydrostatic stress increase with triaxiality, therupture life becomes marginally lower at lower stress levels for thelaser weldments. This difference in rupture lives of cross weldspecimens is more prominent at lower stress levels, which needsfurther investigation. It is worth noting that all creep testedspecimens had failed in the ICHAZ of the weldment, which is

Fig. 14. The distribution of triaxiality factor for flat cross weld creep specimen at (a) half depth and (b) 3/4 depth from the surface.

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known as “Type IV” cracking. Fig. 12(a) and (b) shows creepcavities observed in the ICHAZ of the creep tested specimens ofSMA and laser weldments respectively. Observation of highercreep cavities in laser weldment compared to SMA weldmentsagrees with the above discussion [30]. Failure of cross weldspecimens in the ICHAZ is further supported by lower rupture lifeof ICHAZ simulated specimen, under the same tests conditions.

The hardness profile taken across the weld interface, shown inFig. 8(a) and (b), shows further dip in hardness in the ICHAZ of theSMA weldment. It clearly suggests that microstructural damagewas more in the ICHAZ of SMA weldment compared to that inlaser weldment during welding. From the present investigation itis understood that the width of the HAZ has a direct influence onthe creep rupture life of P91 steel weldments at higher stress level,which was absent at lower stress level. Therefore, the use ofadvanced welding process alone may not improve the creeprupture life of P91 steel or similar steels where the microstructureis unstable in HAZ at longer creep exposure. Present results andthose reported in literature suggest that the softening of themicrostructure leads to the formation of metallurgical notch (softzone) in the weldment. With the progress of the creep test, themicrostructure evolves continuously accompanied by decrease indislocation density and coarsening of precipitates in the ICHAZ orFGHAZ. At lower stress levels, the creep cavitations take place dueto constrained deformation imposed by the surrounding hardmaterials. This results in loss of ductility of the weldment.

3.4. Finite element analysis

It is important to understand the creep deformation anddamage accumulation in components after long-term service athigh temperatures. Though the replica method can be used for thispurpose, the application of this technique is limited to the surfaceof the component. Therefore this technique cannot be used toestimate the cavitation behavior which occurs in the interior of thespecimen. Difference in creep deformation characteristics of theconstituents of weldment leads to stress redistribution across it[32]. Further, the constrained creep deformation results in multi-axial state of stress across the weldment, as in the case of notchedspecimens [33], which influences the creep deformation andrupture behavior of the weldment. Hence, 3D FE analysis has beencarried out to understand the evolution of stress in the interior ofthe weldment and its effect on creep cavitation. Since the creepcavities were observed in the ICHAZ of the weldment predomi-nantly, only this region was considered for the analysis. The widthof ICHAZ was considered based on hardness profile and micro-structural variation across the respective weldments. The width ofICHAZ in SMAW was around 1.5 mm and 0.2 mm for laser welding.

In the present study, the triaxiality was calculated from thevariation of different components of principal and von-Misesstresses through the equation, Triaxiality Factor (TF)¼ðσ1þσ2þσ3Þ=σvm, where σ1 is the maximum principal stress andσ2 and σ3 are intermediate and minimum principal stressesrespectively and σvm is the von-Mises stress. Triaxiality factorcalculated using FE analysis and found to be maximum at thecentral region of HAZ for both the ICHAZ widths considered in thisanalysis (Fig. 13(a) and (b)). The maximum value of TF decreasedwith the increase in the width of ICHAZ. The maximum value ofTF was found to be around 2.45 for 1.5 mm width, whereas, 2.65for 0.5 mm width. This suggests that the width of HAZ hasa significant influence on TF and rupture life. In addition to that,the creep strain was also found to be maximum at the centralregion of the ICHAZ in both the weldments. The distribution of TFvalue for flat cross weld creep specimen prepared by using laserwelding process is shown in Fig. 14. The triaxiality factor with amaximum value of 3.73 was found to be in the ICHAZ of the

weldment. Higher triaxiality in the ICHAZ of laser specimen is inagreement with the observed high density of creep cavities in thesame location compared to SMAW cross weld creep specimen.Here it is worth noting that the geometry may also have aninfluence on distribution of TF. Near center region is expected todeform in plane strain condition because of the constraintimposed by the surrounding material. Whereas, near surfaceregion deforms under plane stress conditions because of therelatively unconstrained free surface. From this argument, it isclearly understood that the triaxiality would be higher at thecentral region of the specimen in the homogeneous material [32].In the present study, higher triaxiality was observed in the ICHAZthan that of the other regions of the weldment. This is attributedto the further constraint effect imposed in the soft ICHAZ by therelatively harder surrounding material in the weldment (Fig. 14(a) and (b)). This observation is in agreement with more cavities atthe middle of the specimen compared to the surface of thespecimen (Fig. 12).

4. Conclusions

Microstructural evolution in the HAZs of SMA and laser weld-ments is similar but microstructural damage is less in laserweldment. The width of the HAZ is significantly smaller for laserweldment compared to the SMA weldment. Hardness dipobserved in the ICHAZ of SMA weldment was absent in the laserweldments. Though the small width of HAZ improved the rupturelife of the laser weldment at higher applied stress, this beneficialeffect was absent at lower stress levels. All creep tested specimensfailed in the ICHAZ of the weldment due to lower creep strength ofthis region. Relatively higher creep cavitation in the laser weldedspecimen is in agreement with high triaxiality compared to theSMAW specimen. From this study it is clear that the application oflaser welding alone may not improve Type IV cracking resistanceof the P91 steel weldment.

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