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JANUARY 2016 / WELDING JOURNAL 27-s

Introduction P91 steel, known as modified 9Cr-1Mo-V-Nb steel (ASTM A335 P91 forpipe), has been widely used for pres-sure vessels and superheaters in fossilfuel power plants due to its high corro-sion resistance and creep resistance(Refs. 1–3). The as-received P91 steelis usually normalized at 1038°–1149°Cand tempered at a minimum tempera-ture of 732°C (Ref. 4). The microstruc-ture of the P91 base metal is reported

to consist of tempered martensite anddispersed carbides/carbonitrides atthe prior austenite grain boundaries(PAGBs) and martensite lath/packetboundaries (Refs. 4–6). The 10,000-hour lifetime rupture strength of P91steel can be 200 MPa at 500°C (Ref. 7).However, creep strength of P91 weldsdramatically decreases due to the non-equilibrium structure of the heat-affected zone (HAZ) caused by expo-sure to various peak temperatures andhigh heating and cooling rates in the

welding thermal cycles. The HAZ of aP91 weld is divided into coarse-grained HAZ (CGHAZ, Tp>> AC3), fine-grained HAZ (FGHAZ, Tp > AC3), andintercritical HAZ (ICHAZ, AC3 > Tp >AC1) based on the prior austenite grainsize in various peak temperatures(Refs. 8, 9). Figure 1 shows typical heat-affectedzone regions and their dimensions in a25-mm-thick P91 pipe joint, fusionwelded using a matching flux-coredwire. The microstructure of the ICHAZ(Fig. 1D) shows a mixture of finegrains formed during welding thermalcycle and coarser grains retained fromthe base metal. Several factors contribute to thelong-term creep performance of P91welds, including chemical compositionof base metal and filling materials,joint design, welding procedure, andheat treatment. In the lifetime of aP91 weld, it experiences three thermalstages, including the as-welded (AW),postweld heat treated (PWHT), andhigh-temperature service exposed(crept). Many researchers have investi-gated the microstructure of welds atthe individual thermal stage. To theknowledge of the authors, no re-searcher has studied a given set ofwelds and tracked their microstructureevolution for all three stages. The relatively low peak temperature(just above Ac3) and short dwell time inthe FGHAZ during the welding thermalcycle are believed to limit the graingrowth after diffusive / transforma-tion and dissolution of coarse precipi-tate particles. This low peak tempera-ture produces the fine prior austenite

Microstructure Evolution of FineGrainedHeatAffected Zone in Type IV Failure of P91 Welds

This work studies the precipitation behavior of Crrich M23C6 and MX carbonitrides andmartensitic/ferritic grain evolution in FGHAZ of P91 welds over three thermal stages

BY Y. WANG AND L. LI

ABSTRACT A microstructure analysis has been conducted on the finegrained heataffectedzone (FGHAZ) of P91 welds that have undergone sequential thermomechanical histories: the aswelded (AW), after a postweld heat treatment, and after a creep rupturetest. The Type IV failure is observed to have initiated by wedgetype cavities that coalesce to form an intergranular separation with a high local deformation at the outeredge of the FGHAZ. The ratio of the highangle grain boundary (HAGB) frequency tolowangle grain boundary (LAGB) frequency in the FGHAZ increases from the aswelded 0.48 to the crept 1.06, resulting in an increased grain size from 1.2 m (AW)to 3.1 m (crept). The area fraction of the fully recrystallized ferrite grains increasesfrom 14% (AW) to 45% (crept). Coarse, undissolved, Crrich M23C6type carbidesremain on prior austenite grain boundaries and packet/block boundaries after welding. V (C, N) and Nb (C, N) carbonitrides distribute at HAGBs. The MoC/Mo2C andM7C3 precipitates form a grain boundary network that is believed to help the nucleation of cavities during creep. The kernel average misorientation (KAM) maps showthe higher local strain level in the particlecontaining grains. Straining variation of theparticlecontaining grains and particlefree grains in the crept FGHAZ can be tracedback to the degree of austenitization in the welding thermal cycle. Creep deformationmismatch between fully recrystallized grains that are free of precipitates and the unrecrystallized grains that contain carbide particles seems to have contributed to nucleation of the initial creep cavities.

KEYWORDS• P91 Steel • Electron BackScatter Diffraction (EBSD) • FineGrained HeatAffected Zone (FGHAZ) • Microstructure • Type IV Creep Rupture

Y. WANG is a PhD candidate and L. LI ([email protected]) is a professor with the Department of Chemical and Materials Engineering, University of Alberta, Edmonton, AB, Canada.

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WELDING JOURNAL / JANUARY 2016, VOL. 9528-s

grains and undissolved particles in theFGHAZ (Ref. 10). Sawada et al. (Ref. 11)have reported a martensitic structurewith very fine laths that contain a highdislocation density in the FGHAZ ofmodified 9Cr-1Mo steel weld. Undis-solved M23C6 and V-Nb MX particles aredetected at the grain boundaries of theFGHAZ. Paddea et al. (Ref. 12) reported

FGHAZ has a hightensile residual stressbecause of thestrength mismatchbetween the base met-al and fusion zone.The highest tensileresidual stress of 600MPa is located at theboundary of HAZ andadjacent base metals,and a decreased ten-sile residual stress of120 MPa still remainsin the vicinity of theHAZ after PWHT.

Postweld heattreatment at a tem-perature below theAc1 temperature is re-

quired to reduce the residual stressesand to temper the fresh martensite.During PWHT, recovery and dynamicrecrystallization of the martensitelaths along with coarsening and redis-tribution of precipitates are observedin HAZ. It is reported that the hard-ness in FGHAZ close to the base metaldecreases faster due to a higher anni-

hilation rate of excess dislocationsduring PWHT (Ref. 13). Liu et al. (Ref.14) found that the equiaxed grains inthe FGHAZ after PWHT are free ofdislocations, resulting in the distribu-tion of carbides/carbonitrides on thePAGBs and block boundaries. The microstructure evolution ofmartensite laths and precipitates inthe FGHAZ is reported to determinethe creep behavior of P91 welds (Refs.15–18). The complex microstructureevolution of P91 welds during creep isassociated with the recovery and re-crystallization of martensite laths andthe coarsening and redistribution ofprecipitates. As a premature failuremode of welded P91, the typical TypeIV creep rupture occurs in the transi-tion zone between the FGHAZ and theICHAZ (Refs. 19, 20). The Type IVrupture is considered brittle becauseof a small total creep deformation atfailure. The nucleation and growth ofcavities in the FGHAZ/ICHAZ are re-ported as the main mechanism of theintergranular fracture of P91 weldsduring creep (Ref. 21). Coarse M23C6

particles on the PAGBs, especially onthe intersection of grain boundaries,are suggested as the preferential nu-cleation sites for cavities (Refs. 10,11). The grain growth of the newly re-crystallized grains and coarsening ofM23C6 carbides are reported to be accel-erated by the applied stress duringcreep (Ref. 22). Parker et al. (Refs. 23, 24) reportthat a local band, free of martensitictransformation, within the FGHAZ/ICHAZ, leads to a local high residualstress, which is responsible for the TypeIV failure of P91 welds. AbdEl-Azim etal. (Ref. 25) have investigated the mi-crostructural instability of FGHAZ ofP91 welds after PWHT and long-termcreep at 650°C and concluded that theprior austenite grain size (PAGS) is themain factor that affects the creep be-havior of the welds. They further pointout that the finer the PAGS, as observedin the FGHAZ, the higher the recoveryrate of excess dislocations, which leadsto higher rates of subgrain growth andprecipitate coarsening. On the other hand, Liu et al. (Ref.26) have reported that the fine-grainedmicrostructure in the FGHAZ is not di-rectly responsible for the Type IV fail-ure. They believe that the precipitateformation behavior at PAGBs and

Fig. 1 — Microstructure of the heataffected zone (HAZ) in the aswelded P91 specimen (10%Nital etched). A — Overview of the HAZ from the fusion zone (FZ) to the base metal (BM); B —coarsegrained HAZ (CGHAZ); C — finegrained HAZ (FGHAZ); D — intercritical HAZ (ICHAZ).

Fig. 3 — The logic flow for deciding whether a grain is recrystallized, contains subgrains, or is deformed is found by comparing the grain average misorientation (GAM) with athreshold of 1.0 deg.

Fig. 2 — A — Schematic of crossweld creep test specimen extracted from the weldedpipe after PWHT; B — the dimensions of the creep rupture specimen (in mm).

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lath/block boundaries is the critical fac-tor. Boundary strengthening by precipi-tates plays the most important role inpreventing Type IV failure of the welds.They proposed a boundary-strengthen-ing method by addition of boron alloy-ing elements. Boron promotes precipi-tate formation not only at PAGBs, butalso along the lath/block boundaries,which helps to stabilize the microstruc-ture of HAZ and prolong the creep life-time of the welds.

The objective of this work is tostudy the precipitation behavior of Cr-rich M23C6 carbide and MX carboni-trides, and martensitic/ferritic grainevolution in FGHAZ of P91 weld inthree sequential thermal stages, con-sisting of the as-welded, after PWHT,and after creep test. This work sum-marizes the evolution of FGHAZ mi-crostructure in Type IV creep ruptureof P91 welds, including matrix grainsize, grain boundary types, precipitatetype, size, and distribution.

Materials andExperimental ProcedureMaterials and Welding Process

The P91 steel pipe used as the basematerial had a 219-mm (8.625-in.) out-er diameter (OD) and 29-mm (1.143-in.) thickness. It had been normalizedfor 8 min at 1060°C and tempered for45 min at 786°C. Flux cored arc welding(FCAW) and gas tungsten arc welding(GTAW) were used to weld two 124-mm-long pipes with a 60-deg double-Vweld groove and 1.5-mm root face, re-spectively. Chemical compositions ofP91 base material, GTAW filler metalER90S-B9 (1.2-mm diameter) andFCAW filler metal ER91T1-B9 (1.2-mmdiameter) are presented in Table 1.

Gas tungsten arc welding parametersused for the root pass were 300 A DCand 1.27 m/min wire feed speed. FCAWparameters for the filling passes were26.1- and 27-arc voltage and 6.35 and7.62 m/min wire feed speed. A steppermotor-controlled fixture was used to

maintain 0.14m/min linear travelspeed. A preheattemperature of150°C, an interpasstemperature of300°C, and a wirefeed speed of 7620mm/min at 27 Vwere selected.Shielding gases usedfor GTAW andFCAW were pure ar-gon and mixed75/25 argon/CO2,respectively. For hy-drogen bake-out, apostweld isothermalhold was conductedat 250°C for 4 h.Postweld heat treat-ment of the weldwas performed in aCress electric fur-nace ModelC162012/SD. ThePWHT parameter isthe code recom-mended followingASME B31.1 —760°C for 2 h.

MechanicalProperty andMicrostructureCharacterization

Creep tests wereperformed on theP91 weld joint afterPWHT and the as-received base met-al. A cross-weld specimen for creeptest was extracted from the center ofthe weld after PWHT. This cross-weldspecimen includes FCAW weld metal,and HAZs on both sides of the weldmetal and the base metal. The rectan-gular cross-section creep specimenhad a cross section of 8.8 × 12.7 mmand an effective gauge length of 84mm, shown in Fig. 2. The creep testwas conducted at applied test temper-

ature of 650°C and engineering stressof 70 MPa (10 ksi). The creep test pro-cedure was conducted following thestandard test procedures guided byASTM E139-06 (Ref. 27). Vickers hardness across the cross-section of polished welds was meas-ured with 1 kgf and a dwell time of 10s by using a Tukon 2500 automatedhardness tester. For microstructurecharacterization, six specimens from

Fig. 4 — Creep curves of the P91 crossweld specimen after PWHT(760°C for 2 h) and the asreceived base metal. Creep tests wereperformed at 650°C with a stress level of 70 MPa. The test of theasreceived base metal was terminated in the secondary creepregime. Creep rupture occurred in one HAZ of the crossweld specimen. The other HAZ (not ruptured but contained cavitation damages) was used for the posttest evaluations.

Fig. 5 — Microhardness across the P91 weld joint in three thermomechanical conditions. The criteria for the CG/FG and FG/ICboundaries are determined by grain size measurements.

Table 1 — Chemical Composition of Grade 91 Base Material and Filler Metals (wt%)

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

ER90SB9 GTAW 0.097 8.830 0.928 0.560 0.250 0.307 0.002 0.197 0.064 0.004 0.006 0.030 E91T1B9 FCAW 0.100 8.830 0.880 0.790 0.280 0.550 0.001 0.200 0.030 0.008 0.020 0.050 Base Metal 0.110 8.470 0.940 0.370 0.370 0.080 0.002 0.190 0.071 0.002 0.016 0.048


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HAZ were cut from the three thermalconditions, as-welded (AW), afterpostweld heat treatment (PWHT), andafter creep test (crept). These speci-mens were mounted, ground, and pol-ished by using a conventional mechan-ical polishing method. Grit #360,#600, #1200 SiC sandpapers were usedfor grinding. Three-m and 1-m dia-mond suspension, 0.5-m and 0.05-m alumina suspension, and 0.02-mcolloidal silica were used for polishing.Electron backscatter diffraction(EBSD) analysis on FGHAZ was con-ducted on two specimens for each testcondition by using a Zeiss Sigma field-emission SEM (FESEM) equipped withan Oxford AZtec system under 20-kVaccelerate voltage, and 60-m objec-tive aperture. A spot size of 10 nm and

step size of 0.15, 0.10, and 0.13 mwere used in EBSD for AW, PWHT, andcrept specimens, respectively. The EBSD-analyzed areas for eachtest condition were 390 m2 (as-weld-ed), 390 m2 (PWHT), and 1049 m2

(crept). The increase in the area for thecrept samples was to make sure a com-parable number of grains was ana-lyzed, as there was significant graingrowth in the crept samples. Speci-mens for EBSD characterization werein as-polished condition without etch-ing. The polished specimens for opticaland FESEM microscope analysis wereetched with 10% Nital for 2 to 10 min. Phase identification and orienta-tion relationship of the matrix grainsand precipitates were analyzed byEBSD with Oxford AZtec software.

Postprocessing EBSD data was con-ducted by using a Channel 5 softwarepackage. Phase distribution maps werecreated based on the indexed Kikuchipatterns. The area fraction of the pre-cipitates was calculated according tothe indexed pixels. Matrix grain orien-tation was presented by the inversepole figure (IPF) in Z [001] direction(perpendicular to the screen). Low-angle grain boundaries (LAGB, misori-entation between 2 and 15 deg) andhigh-angle grain boundaries (HAGB,misorientation > 15 deg) were dis-tinguished by the scalar misorienta-tion between adjacent pixels. Localmisorientation maps were processedby calculating kernel average misorien-tation (KAM) between all pixels of akernel (Refs. 28–30). The KAM valuewill be relatively larger if the calculat-ed kernel contains a grain boundary. Ahigher local misorientation correlateswith a higher local strain and deforma-tion. Grain average misorientation(GAM) with a threshold (c) was usedto quantitatively distinguish the un-recrystallized grain and recrystallizedgrains. The GAM is the average misori-entation between each neighboringpair of the measured pixels withineach grain, which evaluates the localstrain of individual grain (Refs. 27,28). The GAM is calculated as follows (Ref. 31):

where N is the number of pixels in agrain, n is the number of neighboringpixels, and ji is the misorientation be-tween the measured pixel j and neigh-boring pixel i. It is proved that GAMprovides a reliable measurement of re-crystallized grains (Refs. 32, 33). Inthis study, the recrystallized fractionof grains in the FGHAZ was evaluatedby using GAM with a commonly usedthreshold (c) of 1.0 deg. The logic flow for classification ofthe grains is shown in Fig. 3. The inter-nal average misorientation angle withinthe grain and misorientation betweensubgrains was measured to categorizethe grains. The grain with an internalgrain average misorientation (GAM)higher than the threshold (c) is definedas an “unrecrystallized grain.” When theGAM of the grain is lower than thethreshold (c), but the misorientationbetween the subgrains (SB) is higher

GAM = 1N

1n

Nj = 1� n

i = 1� � ji j � i( ) (1)

Fig. 6 — Typical Type IV failure of the crosssectional weld after creep test. A — Macrostructure of the ruptured weld; B — backscattered electron (BSE) image of the failed transition zone between FGHAZ and ICHAZ indicates the threestage lifetime of the cavity:initial cavity, growing cavity, and connected cavity.

Fig. 7 — Micrographs of FGHAZ in three thermal conditions. A — Aswelded; B — afterPWHT; C — after creep test.

Fig. 8 — A — Backscattered electron (BSE) images of the aswelded FGHAZ; B — deepetched with 10% Nital to show the prior austenitic grain boundaries (PAGBs).

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than threshold (c), this grain is classi-fied as a “substructured grain” or “par-tially recrystallized grain.” The grainwith GAM lower than the threshold (c)or subgrain misorientation lower thanthe threshold (c) or without subgrainsis defined as a “recrystallized grain.”

Experimental Results

Creep Test and Microhardness

Figure 4 shows creep curves of theP91 cross-weld specimen after PWHTand the as-received base metal tested at650°C with a stress level of 70 MPa. Thecreep curve of the P91 cross-weld showsa typical three-stage curve, includingprimary creep regime, secondary creepregime, and tertiary creep regime. Theprimary creep regime of base metal ismuch longer than that of the weld. Thecurve of the weld shows it quickly en-ters the secondary creep regime with aconstant strain rate. In the tertiarycreep regime, the cavities promote therapid deformation of the weld with anincreasing strain rate. The weld rup-tured at testing time 649 h with an av-erage tensile strain of 1.2% over the en-tire gauge length. The test of the as-received base metal was terminated inthe secondary creep regime at 2402 hwith an average tensile strain of 1.0%over the entire gauge section. It is obvi-ous that the creep rupture strength ofthe weld is much lower than that of thebase metal. Microhardness traverse distribu-tion as affected by the thermal histo-ries is shown in Fig. 5. Hardness of allregions decreased after PWHT and thecreep test. The fusion zone shows thehighest hardness, 463 HV1, in the AWcondition due to untempered marten-site formed following weld cooling.Hardness in AW-CGHAZ gradually de-creases along with the gradient peaktemperature (above AC3) and coolingrate during the welding process. TheFGHAZ is the softest region, with ahardness value of 220 HV1 for AW,206 HV1 for PWHT, and 115 HV1 forthe crept condition,respectively. Figure 6 shows a typical Type IVcreep rupture of the cross-weld P91joint tested at 650°C with a stress levelof 70 MPa. Failure of the weld joint oc-curred in the outer edge of the FGHAZand ICHAZ close to base metal. The

weld metal, much of the HAZ, and thebase metal show no large plastic defor-mation during creep. The plastic de-formation (shown as the shear-deformed steps in Fig. 6A) was ob-served at a region 0.5 mm wide, locat-ed near the outer edge of the HAZ,where the creep cavity concentrated —

Fig. 6B. Because of the creep strengthvariation between fusion zone andbase metal, a shear deformation oc-curred during creep test. Failure of theweld is caused by the formation of thecavities in the shear deformation zone.The cavity damage is observed to havestarted in the middle thickness of the

Fig. 9 — A — Backscattered electron (BSE) images of FGHAZ following PWHT; B — following creep rupture (10% Nital etch).

Fig. 10 — Highangle grain boundary (HAGB) and lowangle grain boundary (LAGB)distributions in FGHAZ. A — In the aswelded; B — postweld heattreated; C — creeprupture tested (dark lines indicate HAGB, misorientation > 15 deg, light lines indicateLAGB, misorientation = 2~15 deg, cavities are marked by arrows).

Fig. 11 — Relative frequency distribution of grain boundary misorientation in FGHAZ. A —Misorientation range from 0 to 15 deg; B — misorientation range from 15 to 62.8 deg.Aswelded FGHAZ (red circle and line), PWHT FGHAZ (blue triangle and line, and creptFGHAZ (black cubic and line). Mackenzie distribution for random misorientation is alsoshown in B.

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cross-weld specimen. Figure 6B pres-ents the three-stage lifetime of cavi-ties, which includes initial cavities,growing cavities, and connected cavi-ties. The connected cavities promotethe generation and propagation of themicrocracks, which reduces the creepstrength of the weld. Creep strain(shear strain) is mostly concentratedat the narrow Type IV damage zone.The local strain (tensile strain andshear strain) of 8.0% in the damagezone marked by the rectangle at theright bottom of the specimen is esti-mated by measuring the “boxed”shear-deformed step in Fig. 6A. Thisestimated local strain near the shear

steps, measured as

in Fig. 6A, is much higher than the av-erage tensile strain of 1.2% over theentire gauge length, measured at rup-ture of the specimen. The resolvedshear and tensile strains, as affectedby the groove angle of the weld joint,have significant effects on creep rup-ture of P91 cross-weld specimens(Refs. 23, 24).

Evolution of FGHAZ Microstructure

Figure 7 shows the microstructureevolution of FGHAZ in three thermalstages. Martensitic substructure can beobserved within the prior austenitegrains (PAGs) in the as-welded condi-tion. Dark dots along PAGBs or withinthe PAGs are carbides or carbonitrides.Compared with precipitates in AW con-dition, the carbide particles coarsenduring the PWHT, and redistribute

c � b( ) / b = a2 + b2 � b( ) / b

Fig. 12 — Inverse pole figure (IPF) in Z [001] direction (perpendicular to the screen)shows the grain orientation distribution. A — AWFGHAZ; B — PWTHFGHAZ; C —creptFGHAZ.

Fig. 14 — Recovery and recrystallization in FGHAZ. A — AW; B — PWHT; C — crept.Colors for uncrystallized grains (red), substructured/partially recrystallized grain (yellow), and recrystallized grain (blue).

Fig. 13 — Neighboring misorientation profileswithin the grains marked in Fig. 12 (the startpoint is the origin of the arrow).

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along the PAGBs and subgrain bound-aries — Fig. 7B. After creep test, theprecipitates grow even bigger and redis-tribute along PAGBs and within the sub-grains — Figs. 7C, 9B. All wedge-typecavities are observed along the PAGBs. Backscattered electron (BSE) imagesin Fig. 8 reveal more structural detailsof the as-welded FGHAZ. Fine andcoarse precipitates are distributed onPAGBs and subgrain boundaries. Deep-etched FGHAZ in Fig. 8B clearly showsthe short martensite laths within PAGs.Fine precipitates distribute within themartensite laths or on lath boundaries.After PWHT, the lath martensite sub-structure is tempered and decomposedinto a fine equiaxed ferrite — Fig. 9A.PAGBs with coarse precipitates are stillvisible. Coarse precipitates stay at theformer martensite lath boundaries. Fin-er particles precipitate within theequiaxed subgrains. Following furthertempering and accumulated strain dur-ing creep test, larger equiaxed grains areobserved — Fig. 9B. Some coarse pre-cipitates lie at the triple points of grainboundaries. PAGBs are not clearly re-vealed due the migration of the grain

boundaries during recrystallization andgrain growth. Fine precipitates concen-trate in some particular grains, but notin all grains. The wedge-type cavityforms on those grain boundaries withdispersed coarse precipitates. It must bepointed out that the current discussionis based on distributions of precipitatesdetectable by FESEM and EBSD (such asin Fig. 9), and the effects of precipitatessmaller than 0.1 m are not considered. Low-angle grain boundaries (LAGB,with a misorientation of 2 to 15 deg)and high-angle grain boundaries(HAGB, with a misorientation > 15 deg)of the matrix grains in FGHAZ wereclassified by measuring the misorienta-tion between adjacent grains — Fig. 10.Subgrain boundaries, such as lath andblock/packet boundaries within PAGs,have lower misorientation than that ofPAGBs. PAGBs are identified as HAGBsdue to their higher misorientation. Inthe as-welded condition, the temperedmartensite substructure obtains moreLAGBs — Fig. 10A. After PWHT, recov-ery and recrystallization of the marten-sitic grains result in more HAGBs —Fig. 10B. When creep test is performed,

grain growth and further recrystalliza-tion contribute to even more HAGBs —Fig. 10C. Misorientation distribution of thegrain boundaries in FGHAZ is shownin Fig. 11. Misorientation distributionpeaks of HAGBs in FGHAZ shift from45 deg in the Mackenzie distributionof randomly misoriented grains to50~60 deg due the recovery and re-crystallization effects shown in Fig.11B. Relative frequency of LAGB andHAGB is calculated and presented inTable 2. The ratio of HAGB/LAGB in-creases from 0.48 to 0.88 and 1.06.Fine lath-shaped subgrains are ob-served in AW-FGHAZ. Coarse lath-shaped grains and fine equiaxed grainsare the characteristic grains of PWHT-FGHAZ. Crept-FGHAZ shows coarseequiaxed grains. The average grainsizes, determined by the line interceptprocedure of ASTM E211, of the matrixgrains are shown in Table 2. The as-welded FGHAZ has the smallest aver-age grain size of 1.2 m. After PWHT,the FGHAZ has a slightly increased av-erage grain size of 1.8 m. After creeptest, the FGHAZ has a significantly in-creased grain size of 3.1 m. The inverse pole figure (IPF) in Z[001] direction (perpendicular to the

Fig. 15 — Overlapping of local misorientation (KAM) map and grain boundary map ofFGHAZ shows the high strain level in AW condition and neighboring grains around cavities. A — AW; B — PWHT; C — crept.

Fig. 16 — Kernel average misorientation(KAM) distributions of FGHAZ in three thermal stages.

Table 2 — Grain Size, Phase, GB, and Grain Distribution of FGHAZ in Three Thermal Histories

Ferrite Area Fraction (%) of GB Frequency (%) Area Fraction (%) of FerriteSpecimen Phases Ratio of Condition Grain Size Standard HAGB LAGB HAGB/LAGB μm Deviation Febcc/bct Precipitates (>15 deg) (2~15 deg) Recrystallized Substructured Unrecrystallized

AW 1.2 1.14 88.5 11.5 32.4 67.6 0.48 14.3 45.0 40.7 PWHT 1.8 0.94 88.9 11.1 46.7 53.3 0.88 27.0 68.1 4.9 Crept 3.1 2.67 88.7 11.3 51.5 48.5 1.06 45.1 50.8 3.1

A C

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screen) in Fig. 12 shows the grain ori-entation distribution in FGHAZ. A leg-end in Fig. 12C illustrates differentorientations in different colors. Theas-welded FGHAZ shows the largestgrain orientation variation due to finenucleated grains. Because of the recov-ery and recrystallization of the matrixgrains after PWHT, a relatively lowergrain orientation variation is observedin PWHT-FGHAZ. Crept-FGHAZ hasthe lowest grain orientation variationdue to the highly recrystallized grains. The misorientation profiles in Fig.13 are calculated based on the neigh-boring pixels across grains with similarorientation in three thermal stages. Itclearly shows the grain in the AW con-dition has the highest misorientation(> 4.0 deg), which indicates its greatercrystal distortion. In the PWHT condi-tion, the highest misorientation in theFGHAZ is 1.5 deg. In the crept condi-tion, the FGHAZ shows the lowestmisorientation (0.1 deg). Distribution of the unrecrystallizedgrains, substructured, and recrystal-lized grains, including precipitates, isshown in Fig. 14. The area fraction foreach of the three kinds of matrixgrains is calculated and tabulated in

Table 2. It is notable that the relativefraction of carbides remains a constantduring the evolution of the mi-crostructure, although their sizes anddistributions will change. In the as-welded condition, the FGHAZ has thehighest fraction of unrecrystallizedgrains (28.5%) and the lowest fractionof recrystallized grains (4.9%). On thecontrary, the FGHAZ in the crept con-dition consists of the lowest fractionof unrecrystallized grains (3.1%) andthe highest fraction of fully recrystal-lized grains (45.1%). The substruc-tured grains, which are in the interme-diate state between the unrecrystal-lized and fully recrystallized, are domi-nant in all three conditions of thewelds. Kernel average misorientation(KAM) is the arithmetic mean of thescalar misorientations between groupsof pixels, or kernels, that measures lo-cal strain levels of the grains (Ref. 29).Figure 15 illustrates the local misori-entation distribution in the FGHAZfor the three thermal stages (as-weld-ed, PWHT, and crept) by KAM maps.The blue color stands for a misorienta-tion of 0 deg (lowest deformation) andthe red color stands for a misorienta-

tion of 3 deg (highest deformation).The as-welded FGHAZ shows a rela-tively higher local misorientation inFig. 15A. Furthermore, grain bound-aries have a higher deformation andmisorientation. After PWHT, the unre-crystallized grains are tempered andrecovered, resulting in a lower degreeof deformation and misorientation inFig. 15B. Some intersections of grainboundaries still show a relatively high-er misorientation. After the creep test,the matrix is further tempered to givethe lowest misorientation, but the re-gions close to the cavities still show ahigh misorientation — Fig. 15C. Thestatistical kernel average misorienta-tion distribution is shown in Fig. 16.Relative frequency peak of AW-FG-HAZ shifts to higher misorientationvalues. Crept-FGHAZ obtains thelargest frequency of low-angle misori-entation value. Figure 17 shows the microstructureevolution of the matrix grains and pre-cipitates before and after the creeptest. Firstly, the recovered and recrys-tallized grains turn to coarse equiaxedgrains. Cr-rich M23C6 carbides distrib-ute within some, not all, grains follow-ing the PWHT. This uneven distribu-

Fig. 17 — Overlapping of EBSD phase map and grain boundary map of FGHAZ show coarsening and segregation of the precipitates, especially M23C6 carbides in AW, PWHT, and crept conditions. A — AW; B — PWHT; C — crept. Phases show as different color symbols. Martensite and ferrite are processed as the same phase background color in pink.

A B

C


tion of carbides to some grains is moreapparent following creep. In addition,a coarsening of MX carbonitride is ob-served on HAGBs following creep.Connected MoC/Mo2C and M7C3 car-bide GB networks are detected close tothe cavities.

Discussion The microstructure of creep-damaged FGHAZ has the followingfeatures. Firstly, the location of thewedge-type cavities that lead to TypeIV failure in P91 are in the lower-peaktemperature fine-grained heat-affect-ed zone, or the edge of the heat-affect-ed zone adjacent to the base metal.Secondly, these cavities tend to format the triple points of high-angle grainboundaries that are enriched withcoarsened M23C6 carbides and net-worked MoC/Mo2C-type and M7C3 car-bides. Thirdly, the cavities seem topreferentially form between grainsthat have fully recrystallized and un-recrystallized grains. The first and sec-ond features have been well recog-nized (Refs. 8, 13). The third feature,that the cavities tend to form betweengrains that are markedly different inmisorientation and degree of recrys-tallization, has not been reported be-fore. Based on the EBSD results, it isbelieved that this third feature holdsthe key for an understanding of thefirst two features. It is reasonable to assume that un-der a constant creep stress, the grainsthat contain precipitates are strongerand will generate a smaller strain,while the recrystallized grains that arefree of carbide particles are weakerand will generate a greater strain. Mi-crohardness (103 HV1) at the locationwith more particle-free grains is lowerthan that (126 HV1) of the locationwith more particle-rich grains. Anoth-er evidence is the deformation mis-match across grains that are adjacentto creep cavities (Fig. 15C). It is there-fore proposed that the mismatch increep straining in the FGHAZ maycontribute to the nucleation of cavi-ties. It has been suggested by Parkerthat local strain is responsible for cavi-ty formation (Ref. 23). Figure 15Cseems to provide the microstructureand straining evidence to support thatmechanism. The formation of a mixed mi-

crostructure of stronger (precipitation-containing) grains and weaker (precipi-tation-free) grains in the fine-grainedHAZ is believed to be the result of thewelding thermal cycle. The peak tem-perature for the edge of the fine-grainedHAZ is barely above the AC3. Under thispeak temperature, austenite hasformed, but the carbide particles in thebase metal, especially coarse M23C6 andMX carbides, are only partially dissolvedin FGHAZ. In the as-welded condition(Fig. 17A), the Cr-rich, fine M23C6 car-bides distribute within some, but notall, grains, and those undissolved coarseM23C6 carbides get trapped within PAGs.The MX carbonitrides are identified asV (C, N) in yellow and Nb (C, N) in bluein Fig. 17. Due to the lower diffusivityand smaller concentration of V and Nb,there are not many precipitated V (C, N)and Nb (C, N) carbonitrides. V (C, N)and Nb (C, N) particles stay withinmartensite laths. It is also notable thatsome grains are free of carbides that aredetectable (if larger than 0.1 m) by theEBSD at the magnification. Based on a pseudo binary phase di-agram produced with Thermo-Calc(Ref. 11), the Ac3 temperature isaround 850°C for the 9 mass-% Crsteel. The equilibrium phases are aferrite, austenite, M23C6, and MXphases at 900°C. The Ac3 temperatureof the P91 alloy of this study is 925°C,measured by dilatometry for slowheating and air cooling. During PWHT,martensite grains are further tem-pered and decomposed, as shown inFig. 17B. There is a significant graingrowth. Relatively finer M23C6 carbidesconcentrate within some specificgrains, but coarser M23C6 carbides stayat HAGBs. The majority of V (C, N)and Nb (C, N) carbonitrides distrib-utes at the HAGBs. MoC/Mo2C and Cr-rich M7C3 carbides are also identifiedand distribute along the HAGBs. Re-covery and recrystallization of themartensite laths lead to formation offiner equiaxed ferrite grains. Move-ment of excess dislocation acceleratesthe recovery and recrystallization ofthe matrix grains during creep defor-mation with the help of applied stressand creep strain (Refs. 34, 35). In-creasing ratio of HAGB/LAGB from0.48 to 1.06, as shown in Fig. 10, indi-cates the decreasing stability of grainstructure when FGHAZ is exposed tothe high creep temperature and stress.

Conclusions Characteristic features of Type IVcreep rupture of P91 welds have beeninvestigated by analyzing microstruc-ture evolution in FGHAZ, under contin-uous thermal histories, as-welded, afterPWHT, and after creep test. The mainfindings are concluded as follows: a) The as-welded FGHAZ has a mi-crostructure of martensite laths with-in fine prior austenite grains. Theundissolved, Cr-rich M23C6 carbidesare distributed on the PAGBs andsome lath boundaries. Fine M23C6 car-bides nucleate on the PAGBs and inter-nal martensite laths. EBSD GAM andKAM maps of the FGHAZ show thehighest fraction of unrecrystallizedgrains with a high strain level in theas-welded condition. b) PWHT tempering produces fineequiaxed ferrite grains and coarsenedprecipitates. The creep test furthertempers the martensite, and promotesthe grain growth with an increasedHAGB/LAGB ratio from 0.48 as weld-ed to 1.06 as crept. Coarsening andsegregation of Cr-rich M23C6 carbidesin specific grains causes a higher localstrain during creep. MX particlescoarsen and stabilize on the high anglegrain boundaries. ContinuousMoC/Mo2C and M7C3 carbide bound-ary networks on HAGBs weaken thecreep strength of GBs. Connectedwedge-type cavities on GBs cause theintergranular Type IV rupture with alow entire creep strain and a large localdeformation in the transition zone be-tween FGHAZ and ICHAZ. c) The welding thermal cycles in theICHAZ and outer edge of the FGHAZdo not have high enough peak temper-atures to promote the formation ofgrains with uniformly distributed pre-cipitates. Instead, the lower peak tem-perature tends to retain the variationsin precipitate distribution in the basemetal microstructure. Recrystalliza-tion during PWHT seems to have fur-ther enhanced the nonuniform distri-bution of carbides, i.e., the formationof precipitate-rich and precipitate-freegrains. It is proposed that the creepstrain mismatch between the unre-crystallized grains that seem to con-tain precipitates and fully recrystal-lized grains that seem to be precipi-tate-free promotes the nucleation ofinitial cavities during creep.

WELDING RESEARCH



This work has been financially spon-sored by the U.S. DOE Office of NuclearEnergy’s Nuclear Energy University Pro-grams and by the Discovery program ofNatural Sciences and Engineering Re-search Council (NSERC) of Canada.

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