Introduction

Allvac 718Plus® is a newly developed γ ′ — strengthened-nickel-based superal-loy. It was developed by ATI Allvac (Ref.1) as a derivative of the widely used In-conel® 718 superalloy. The alloy was de-veloped to overcome service temperaturelimitations that are usually encounteredwith components made with 718. Alloy718 loses its strength above 650°C due toinstability of its main strengthening phase,γ ″. Above this temperature, the D022 or-dered body -centered tetragonal phase ofγ ″ becomes unstable and transforms to anequilibrium delta (δ) phase with or-thorhombic (D0a) crystal structure with aplate-like morphology, which impairs itshigh-temperature mechanical properties.In developing 718Plus, major composi-tional changes, which are detailed in thework of Cao et al. (Ref. 2), were made to

its principal alloy, 718. These changes re-sulted in a change in the main strengthen-ing phase of 718Plus from γ ″ to γ ′ phasethat is known to be more stable at hightemperatures and thus improves thestrength of the superalloy. Thus, it isclaimed (Ref. 2) that hot-section compo-nents of both aero-engine and land-basedpower-generation turbines made of718Plus superalloy can operate at temper-atures up to 55°C higher than those madeof 718.

Although 718Plus can withstand highservice temperatures, recent studies on itsweldability (Refs. 3, 4) have shown that itis susceptible to heat-affected zone (HAZ)cracking during high-power beam welding,a problem that plagues many other nickel-based superalloys. It is known that HAZcracking during welding of superalloys canbe significantly reduced or eliminated bythe use of appropriate preweld heat treat-ments (Ref. 5). Therefore, the aim of thepresent work was to investigate the influ-ence of preweld heat treatments on thecracking behavior of 718Plus. Ultimately,a preweld heat treatment schedule thatwould be capable of preventing the occur-rence of HAZ cracking during weldingwas sought.

Experimental Procedures

Materials and Preweld Heat Treatments

The 718Plus superalloy used in thisstudy was supplied by ATI Allvac in theform of 305 × 127 × 16 mm hot rolledplate. The alloy contained about 30 ppmof boron, and the rest of its elementalcomposition is given in Table 1. Rectangu-lar sections measuring 12 × 12 × 100 mmand 3 × 3 × 2 mm were cut normal to therolling direction of the plate, and then so-lution heat treated for 1 h in an argon at-mosphere in a furnace at 950°, 1050°, and1150°C. The heat-treated coupons werecooled from the solution heat-treatmenttemperatures at various rates ranging fromiced-water quenching (≈500°C/s), air cool-ing (≈25°C/s), to furnace cooling(≈0.25°C/s), as shown in Table 2.

SIMS Analysis

The SIMS technique was used to ana-lyze grain boundary distribution of boronin this study. In the SIMS (Refs. 5, 6), asolid sample is bombarded with primaryions of few keV energy. This results in thesputtering of atomic species from the sur-face of the sample. Fractions of the speciesare emitted as negative, neutral, or posi-tive secondary ions and their mass are an-alyzed in a mass spectrometer todetermine the elemental, molecular, orisotopic composition of the surface. Dueto the extremely high signal/backgroundratio of the mass spectrometer, SIMS candetect trace amounts of elements on a sur-face. It provides ion images with spatialresolution of a few μm and elemental res-olution to a few ppm. The primary ionbeam species used in SIMS include Cs+,O2+, O–, Ar+, Ga+, and Xe+. Very elec-tronegative elements such as H, O, S, andP are most sensitive to Cs+ beam, whileO2+ is used for most other elements.

In preparation for SIMS analysis in this

SUPPLEMENT TO THE WELDING JOURNAL, SEPTEMBER 2009Sponsored by the American Welding Society and the Welding Research Council

Crack-Free Electron Beam Welding of Allvac 718Plus® Superalloy

Careful control of preweld heat treatment and cooling rates was found tosiginificantly reduce HAZ cracking

BY O. A. IDOWU, O. A. OJO, AND M. C. CHATURVEDI

KEYWORDS

Allvac 718Plus®Inconel® 718SuperalloyBoronGrain Boundary SegregationGrain SizeSIMS AnalysisMicrostructural AnalysisHeat TreatmentElectron Beam WeldingWeld HAZ Cracking

O. A. IDOWU, O. A. OJO, and M. C.CHATURVEDI are with the Department of Me-chanical and Manufacturing Engineering, Uni-versity of Manitoba, Winnipeg, Manitoba,Canada.

ABSTRACT

Allvac 718Plus® was welded in dif-ferent preweld heat-treated conditionsin order to investigate the effects ofpreweld heat treatments on its crack-ing behavior during welding. The oc-currence of HAZ cracking in the alloywas found to be closely related tograin boundary boron segregation andvariation in grain size, both of whichwere controlled by preweld heat treat-ments. It was demonstrated that crack-free welds of the alloy can be made bya proper selection of preweld heattreatment.

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study, 3 × 3 × 2 mm preweld heat-treatedcoupons were mounted in an epoxy baseresin and polished to 1-μm finish using stan-dard metallographic techniques. Thereafter,they were electro-etched in 10% oxalic acidat 6 V for 5 s. This revealed the grainboundaries, which were the areas of inter-est for SIMS analyses. A Vickers micro-hardness indenter was used to mark thegrain boundaries, and the coupons were re-polished to remove the etched surface layerwhile leaving the indentations. Finally, thecoupons were removed from the epoxy baseresin in which they were mounted. Thecoupons were then examined for grainboundary segregation of boron using aCameca IMS-7f magnetic sector secondaryion mass spectrometer. The equipment wasoperated in a scanning microprobe mode.A primary ion beam of O2+ with impact en-

ergy of 10 kV and 180pA beam current wasrastered over an ap-proximate surface areaof 150 × 150 μm of thecoupon. Mass resolvedimages of boron wereobtained by imagingpositive secondary ionsof 11B+.

Welding and Micro-structural Analyses

In order to investigatethe effects of variouspreweld heat treat-ments on HAZ crack-ing of 718Plus duringwelding, the 12 × 12 ×100 mm heat-treatedspecimens were welded

autogenously along the 100 mm lengthusing a focused electron beam at 44 kVvoltage, 79 mA beam current, and 152cm/min welding speed. Transverse sectionsof the welds were cut by electrodischargemachining and prepared by standard met-allographic polishing technique for mi-crostructural study. Thereafter, thespecimens were etched electrolytically at6 V for 5 s in a solution of 10% oxalic acid.Microstructures of preweld heat-treatedand welded specimens were examinedusing an optical microscope and a JEOL5900 scanning electron microscope (SEM)in both secondary and backscattered elec-tron (BSE) imaging modes. The suscepti-bility of 718Plus to weld HAZ crackingwas evaluated by measuring total length ofcracks (TCL) observed in 8 sections ofeach weld.

Results and Discussion

Microstructure of Preweld Solution HeatTreated Alloy

Figures 1A–C and 2A–C show SEMmicrostructure of 718Plus after solutionheat treatment for 1 h at 950° and 1050°C,respectively. Samples that were heattreated at 1150°C had microstructuressimilar to those treated at 1050°C. Gener-ally, after the solution heat treatment at950°C, and at all the cooling rates em-ployed, the alloy contained dispersedblocky-shaped precipitates in intergranu-lar and intragranular regions of its γ solidsolution matrix (Fig. 1A–C). The precipi-tates were identified to be Nb, Ti-rich car-boborides and Ti, Nb-rich carbonitrideswhich are the primary solidification con-stituents of the alloy (Ref. 3). Also, grainboundary regions of the alloy were out-lined by secondary precipitates of δ phasehaving platelet or blocky morphology. Inaddition to the carbides and δ phase, pre-cipitates of γ ′ particles were observedwithin the γ matrix of the furnace-cooledsamples. A high-magnification SEMimage showing spherical-shaped γ ′ parti-cles is shown in the top-right inset of Fig.1C. The γ ′particles precipitated from γmatrix during continuous cooling of718Plus in the furnace at a slow rate (≈0.25°C/s). Specimens heat treated at 1050°and 1150°C contained microconstituentssimilar to those observed in the samplesthat were heat treated at 950°C, with theexemption of δ phase. δ-phase precipitateswere not observed after the solution treat-ments at 1050° and 1150°C because thesetemperatures are above the solvus of δphase in 718Plus, which has been reported

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Fig. 1 — SEM microstructure of 718Plus solution heat treated at 950°C. A —Iced-water quenched; B — air cooled; C — furnace cooled. The top-rightinset on 1C shows γ ′ particles in γ matrix of the furnace-cooledsamples.

Table 1 — Bulk Chemical Composition (wt-%) of 718Plus

Ni Cr Fe Co Nb Mo Al W Ti Mn C Si V P B Mg SBal 17.92 9.33 9.00 5.51 2.68 1.50 1.04 0.74 0.03 0.022 0.02 0.02 0.006 0.003 0.0008 0.0003

A

C

B

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to be in the range of 1002°–1018°C (Ref.7). A notable difference was also observedin the grain size of the samples after heattreatments. The optical micrographs inFig. 3 A–C show the grain sizes of speci-mens solution heat treated at 950°C (≈58μm grain size), 1050°C (≈120 μm grainsize), and 1150°C (≈360 μm grain size), re-spectively, and air cooled. The micro-graphs depict the variation in grain size ofthe alloy after these heat treatments. Theapproximate mean diameter of the grainsafter all the preweld solution heat treat-ments are presented graphically in Fig. 4.It is seen that no significant variation ingrain size occurred with variation in cool-ing rate. However, the grain size increasedapproximately by a factor of two and sixwhen heat treatment temperature was in-creased from 950° to 1050°C and 1150°C,respectively. The dramatic grain growththat occurred at 1050° and 1150°C can beattributed to the dissolution of δ-phase pre-cipitates during heat treatments at thesetemperatures, as shown in Fig. 2. It isknown that δ phase pins grain boundaries,and thereby prevents/limits grain growthduring high-temperature exposure (Ref. 8).Thus, no significant grain growth occurredduring the treatment at 950°C where thegrain boundaries were outlined and suc-cessfully pinned by δ phase — Fig. 1.

HAZ Microstructure of Welded Alloy

The SEM backscattered electron (BSE)image in Fig. 5 shows weld HAZ mi-crostructure of the specimen that was

preweld heat treated at950°C and quenched iniced water. Liquation ofMC-type carbide particleswas occasionally observedin the HAZ region closerto the fusion zone of theweld. Also, grain boundaryliquation was observed inHAZ region up to about250 μm from the fusionzone, but HAZ cracking was not observed.The samples that were preweld heat-treated at 950°C and air or furnace cooledwere also crack free. Their HAZ mi-crostructures were similar to those of thesamples that were quenched in iced water— Fig. 5. However, SEM microstructuresof the preweld and weld HAZ of the fur-nace-cooled sample (Fig. 6A, B) suggeststhat liquation of γ ′ particles occurred in theheat-affected zone during welding. The un-derlying mechanism of the constitutional liquation of γ ′ precipitates in γ ′ strength-ened nickel-based superalloys has been dis-cussed in detail in the work of Ojo et al.(Refs. 9, 10). An explicit consideration ofthe potential effects of this on the weld-ability of 718Plus (in aged condition) is asubject of future work. However, it hasbeen reported (Refs. 9, 10) that liquid filmfrom constitutionally liquated γ ′ precipi-tate can contribute to weld HAZ cracking ifit penetrates and wets HAZ grain boundaries.

The weld HAZ microstructures of sam-ples that were preweld heat treated at1050° and 1150°C, respectively, were sim-

ilar to those that were heat treated at950°C. However, HAZ cracks were ob-served in the samples that were preweldheat treated at 1050° and 1150°C. Repre-sentative optical and SEM images of acracked region is shown in Fig. 7A and B,respectively. As shown in the figures, thecracks were mostly present within the neckregion of the nail-head-shaped welds. Thedegree of cracking in these samples wasquantified by the measurement of totalcrack lengths (TCL) in them. The averagevalue of TCL observed in variously heattreated material is shown in Fig. 8. It isseen that cracking was not observed in thewelds in samples that were preweld heattreated at 950°C and cooled at the threecooling rates. However, as shown in Fig. 8,HAZ cracking occurred in the welds whenthe preweld heat treatment temperaturewas increased to 1050° and 1150°C. Figure8 also shows that, at 1050° and 1150°C, thehighest degree of cracking occurred in air-cooled samples. This was followed by fur-nace-cooled samples, while the lowestextent of cracking was found in samplesthat were quenched in iced water. At allthe three cooling rates, the degree of

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Fig. 2 — SEM microstructure of 718Plus solution heat treated at1050°C. A — Iced-water quenched; B — air cooled; C — furnacecooled. The top-right inset on 1C shows γ ′ particles in γ matrix of fur-nace-cooled samples.

Table 2 — Preweld Heat-Treatment Temperatures and Cooling Method

Temperature (°C) 950 1050 1150Cooling method ID-WQ(a) ID-WQ ID-WQ

AC(b) AC ACFC(c) FC FC

(a) ID-WQ: Iced-water quench at ≈ 500°C/s(b) AC: Air cooled at ≈ 25°C/s(c) FC: Furnace cooled ≈ 0.25°C/s

A

C

B

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cracking was always higher in samplespreweld heat treated at 1150°C comparedto those that were heat treated at 1050°C.Although constitutional liquation of MC-type carbides was observed in this work,and has been reported to contribute toHAZ cracking in 718Plus (Ref. 11), a care-ful study of the variation of cracking withpreweld heat-treatment temperature andcooling rate, which did not affect the vol-ume fractions of constitutionally liquatingMC-type carbides, suggested that crackingwas more influenced by two main factors,viz., segregation of B and grain size of thematerial, which are discussed next.

Grain Boundary Segregation of Boron

The cause of intergranular crackingthat often occurs in the HAZ of austeniticalloys during welding is generally attrib-uted to a combination of mechanical driv-ing force for cracking, threshold tensilewelding stresses, and a crack susceptiblemicrostructure such as with liquated grainboundaries. The liquation of a HAZ grainboundary during welding causes its solid-solid interfacial bond to be replaced by a

weaker solid-liquid bond.Consequently, this re-duces the threshold ten-sile welding stress that isrequired to cause the li-quated grain boundary’sdecohesion. Thus, the in-herent resistance of analloy to HAZ cracking isreduced by liquation re-action. Grain boundary li-quation can occur byeither equilibrium super-solidus or nonequilibriumsubsolidus melting. How-ever, the mere occur-rence of liquationreaction is not sufficientto produce a crack sus-ceptible microstructure.It is essential that the liq-uid film penetrates andwets the grain boundaries

and is stable over a wide range of temper-atures to allow enough stresses to build upduring weld cooling. The presence of pos-itively partitioning elements (i.e., k < 1),particularly boron, enhances the wettabil-ity and increases the stability of grainboundary liquid film, and boron has beendescribed as the most detrimental elementcausing grain boundary liquation cracking(Refs. 12–14), although it improves creep-rupture life of the alloy. Boron is particu-larly more detrimental as it has a tendencyto preferentially segregate to grain bound-aries during thermo mechanical process-ing of an alloy, or thermal treatments priorto welding.

The thermally induced grain boundarysegregation of solute atoms, e.g., boron, isa diffusion controlled process that canoccur by two mechanisms, viz., equilibriumsegregation and nonequilibrium segrega-tion. Equilibrium segregation occurs whena polycrystal is held at a sufficiently hightemperature to permit appreciable diffu-sion of misfit impurity atoms within itsgrains to loosely packed interfaces, e.g.,grain boundaries. The impurity atoms are

absorbed by such interfaces in order to re-duce their free energy (Ref. 15), and arethereby said to “segregate” to the inter-face. The extent of such equilibrium seg-regation is known to decrease with anincrease in heat-treatment temperature,but is independent of cooling rates. On theother hand, thermally induced nonequi-librium segregation (Refs. 16, 17), as itsname suggests, occurs under nonequilib-rium thermal conditions, particularly whencooling from elevated temperatures overa range of cooling rates. This type of seg-regation has been widely described by asolute drag mechanism (Refs. 18, 19),which involves the following. During high-temperature annealing, an equilibriumconcentration of vacancies is generatedand distributed within the grain interior ofa polycrystal. As the temperature de-creases during cooling, the grains becomesupersaturated with vacancies. Since grainboundaries act as perfect sinks for pointdefects, a vacancy concentration gradientis set up between the supersaturatedgrains and the boundaries (Ref. 20). Con-sequently, vacancies diffuse down the con-centration gradient to grain boundarieswhere they are annihilated. However, ifsome of the vacancies are attracted to im-purity atoms, vacancy-solute complexeswill be formed, and a proportion of the im-purity atoms (those with positive vacancy-solute binding energies [Ref. 13]) will bedragged with vacancies toward grainboundaries. Thus, the grain boundarieswill become gradually enriched with im-purity solute atoms while the vacancies arebeing annihilated.

In contrast to equilibrium segregation,the degree of nonequilibrium segregationdepends on cooling rates and vacancy-solute binding energy. The degree of non-equilibrium segregation increases with anincrease in heat treatment temperature.The cooling rate is, however, very critical.If cooling rate is very rapid, sufficient timemay not be available for a significant dif-fusion of complexes to grain boundaries tooccur (Ref. 14). On the other hand, if the

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Fig. 3 — Optical micrographs (100×, 10% oxalic etch) of 718Plusshowing the variation in grain size after solution heat treatment. A —At 950°C; B — at 1050°C; C — at 1150°C. All were followed by air cooling.

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cooling rate is very slow, desegregation ofsolute atoms may occur. That is, segre-gated solute atoms on grain boundariesmay diffuse away from their grain bound-ary sites back to grain interior when asolute concentration gradient occurs be-tween the grain boundaries and the graininterior (Refs. 21, 22). Therefore, the de-gree of nonequilibrium segregation is usu-ally highest at intermediate cooling rateswhen sufficient time is available to let va-cancy-solute complexes diffuse to grainboundaries but is not sufficient to let thedeposited solute atoms diffuse away fromthe grain boundary region.

Figures 9A–C to 11A–C show a set ofSIMS images illustrating grain boundarydistribution of boron in 718Plus after heattreatments at 950°, 1050°, and 1150°C atcooling rates corresponding to iced-waterquench, air cool, and furnace cool. Theimages were recorded at the same instru-mental settings of SIMS. Within the reso-lution limit of SIMS, boron segregationwas significantly minimized in samples

heat treated at 950°C, and cooled by anyof the three cooling rates employed (Fig. 9A–C). However, brighter boron signals,the degree of which varied with coolingrates and appeared to increase with heat-treatment temperature, were observed insamples heat treated at 1050° and 1150°C,respectively (Figs. 10A–C and 11A–C).The brightest boron intensity was ob-served in air-cooled samples while iced-water quenching and furnace coolingresulted in a significant reduction in theintensity. The intensities of boron signalscould not be quantified by SIMS, thus, thiswork limits itself only to a qualitative com-parison of the signals.

The dependence of nonequilibriumsegregation on cooling rates suggests thatthe characteristically dynamic segregationpattern observed in the present work dur-ing preweld heat treatments of 718Plus, asshown in Figs. 9–11, is predominantly non-equilibrium in nature. It has been widelyreported (Refs. 22, 23) that the segrega-tion of boron to grain boundaries in

austenitic steels occurs mainly during cool-ing from high temperatures and by non-equilibrium mechanism. Karlsson et al.(Ref. 24) used SIMS to study the nature ofgrain boundary segregation of boron inType 316L austenitic stainless steel at heattreatment temperatures and cooling rateswhich ranged from 900° to 1250°C, and530° to 0.25°C/s, respectively. They ob-served an increase in the degree of grainboundary segregation of boron when thecooling rate was decreased from 530° to27°C/s. However, a further decrease in thecooling rate to 0.25°C/s resulted in deseg-regation of the boron atoms, an occur-rence which was marked by a decrease inthe intensity of boron atoms on the grainboundaries. In the present investigation,boron segregation was significantly mini-mized in 718Plus that was heat treated at950°C and cooled at various rates (Fig.9A–C), viz., iced-water quench (≈500°C/s),air cool (≈25°C/s), and furnace cool(≈0.25°C/s). However, boron segregated tograin boundaries when the heat-treatment

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Fig. 4 — Variation in average grain size of 718Plus with preweld heat-treat-ment temperature and cooling rate.

Fig. 5 — SEM micrograph showing the absence of HAZ cracking in theweld of 718Plus that was preweld solution heat treated at 950°C and iced-water quenched.

Fig. 6 — SEM micrographs showing γ ′ particles. A — Before welding (i.e., after 950°C + furnace cool); B — after welding.

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temperature was increased to 1050° and1150°C, respectively. As shown in Figs.10A–C and 11A–C at both temperatures,the degree of segregation varied with cool-ing rates. The brightest boron intensitieswere observed in samples that were cooledat the intermediate rate (air cool), whilethe fastest cooling rate (iced-waterquench) resulted in an appreciable de-crease in boron seregation. Moreover, de-segregation of boron occurred duringfurnace cooling. A significant reduction inboron intensity was observed in furnace-cooled samples, although their intensitieswere always brighter than that observed inthe iced-water quenched samples.

Boron is considered to be the mostdetrimental element causing grain bound-ary liquation cracking in Alloy 718 (Refs.12–14), although cracking is also known toresult from the constitutional liquation ofMC-type carbides. Huang et al. (Refs. 13,25) and Chen et al. (Ref. 26) evaluated theweldability of 718, in both cast and

wrought forms,after the alloy waspreweld treatedusing various heattreatment sched-ules that varied thedegree of boronsegregation on itsgrain boundaries.A close relation-ship was observedbetween thepreweld heat treat-ments, the degreeof boron segrega-tion, and the sus-ceptibility of 718 toHAZ cracking,which was evalu-ated by measuringthe total cracklength (TCL) that

were observed in the HAZ. A heat treat-ment that increased boron segregationalso increased cracking in the alloy andvice versa. In the present investigation, nocracking was observed in weld HAZ of thealloy, which was preweld heat treated at950°C and cooled at various rates, whichranged from iced-water quench to furnacecool — Figs. 5, 8. SIMS analyses of thealloy in the preweld heat-treated condi-tions did not reveal a significant segrega-tion of boron to its grain boundary regionseither — Fig. 9. Evaluation of the weldsthat were preweld heat treated at 1050°and 1150°C, respectively, revealed thatHAZ cracking occurred during welding —Figs. 7, 8. The degree of cracking (TCL)in the HAZ of the welds increased with anincrease in the degree of boron segrega-tion, which occurred on the grain bound-aries of the alloy during the respectivepreweld heat treatments. As shown inFigs. 8, 10, and 11, iced-water quenching(fastest cooling rate) after preweld solu-

tion heat treatments at 1050° and 1150°C,respectively, resulted in the lowest degreeof boron segregation and weld HAZcracking. However, the highest degree ofboron segregation and weld HAZ crack-ing were observed in air-cooled (interme-diate cooling rate) coupons. Incomparison to air cooling, the degree ofHAZ microfissuring was reduced incoupons that were furnace cooled; a cool-ing condition where desegregation ofboron from grain boundaries had startedto occur and which reduced the amount ofsegregated boron on the grain boundaries.

To further confirm the contribution ofboron to HAZ cracking in 718Plus, a ver-sion of the alloy that has a higher boronconcentration (60 ppm of boron) waspreweld heat treated for 1 h at 950°, 1050°,and 1150°C, respectively, and air cooled.Thereafter, the coupons were weldedusing the same electron beam weldingconditions that were reported in the ex-perimental technique section of this com-munication. SIMS images of segregatedboron atoms on grain boundaries of thepreweld heat-treated coupons, as well asthe degree of HAZ cracking that occurredduring their welding are shown in Figs. 12and 13, respectively. Similar to the obser-vations in the lower boron version (≈ 30ppm) of 718Plus, boron segregation in-creased with an increase in preweld heattreatment temperature, and so did theweld HAZ cracking. However, it wasnoted that HAZ cracking was alwayshigher in the higher boron version of718Plus — Figs. 8, 13. Also, in contrast tothe lower boron version of the alloy, amore significant grain boundary segrega-tion of boron, accompanied by weld HAZcracking was observed in the higher boronalloy, which was preweld heat treated at950°C. In this heat treatment condition,grain boundary segregation was signifi-cantly reduced and no HAZ cracking was

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Fig. 7 — A — Optical; and B — SEM micrographs showing cracking in the weld HAZ of 718Plus that was preweld solution heat treated at 1150°C plus aircool.

Fig. 8 — Variation in HAZ total crack length (TCL) of 718Plus with preweldheat-treatment temperature and cooling rate.

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observed in the lower boron version of thealloy, as shown in Figs. 8 and 9. This fur-ther confirmed that boron played a signif-icant role in causing HAZ cracking duringwelding of 718Plus. Therefore, its concen-tration should be as low as possible to im-prove the weldability of the alloy.

Generally, boron has been consideredto influence the susceptibility of an alloyto weld HAZ cracking by the followingmechanisms:

1) Segregated boron can act as a melt-ing point depressant and reduce the melt-ing temperature of grain boundary regions

relative to surrounding matrix (Ref. 26).2) Boron can extend the solidification

range of liquid film on HAZ grain bound-aries due to its low partition coefficient.

3) Boron can decrease solid-liquid in-terfacial energy (γ SL) (Refs. 13, 27), whichwould enhance the wettability (and facili-tate spreading) and/or increase the stabil-ity of grain boundary liquid during welding.

The existence of any or a combinationof the above conditions, due to grainboundary boron segregation, which wasobserved in the present study, will increasethe susceptibility of HAZ grain bound-

aries to liquation and facilitate spreadingof the liquid film along the grain bound-aries during welding. The inability of theliquated grain boundaries to support ten-sile stresses that develop during cooling ofthe welds will result in their cracking (Ref.28), as observed in the present work

Preweld Grain Size

In addition to boron segregation, thevariation in grain size of an alloy that takesplace during preweld heat treatments canalso affect the susceptibility of the alloy to

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Fig. 9 — SIMS image of boron in 718Plus after heat treatment at 950°C. A — Iced-water quench; B — air cool; C — furnace cool.

Fig. 10 — SIMS image of boron in 718Plus after heat treatment at 1050°C. A — Iced-water quench; B — air cool; C — furnace cool.

Fig. 11 — SIMS image of boron in 718Plus after heat treatment at 1150°C. A — Iced-water quench; B — air cool; C — furnace cool.

C

C

C

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HAZ cracking during welding. Thompsonet al. (Ref. 29) reported that HAZ crack-ing of Alloy 718 depends linearly on itsgrain size. It is a common knowledge thatcasting components are usually moreprone to cracking than their wroughtcounterparts. This can be partly attributedto the usually larger grain sizes of the cast-ings, apart from contributions from theirsegregated microstructure. Generally, thegrain size of an alloy can also contribute toits weld HAZ cracking due to the following:

1) As predicted in the nonequilibriumgrain boundary segregation model ofFaulkner (Ref. 30), an increase in grainsize could decrease total grain boundarysurface area that is available for grainboundary segregation. Thus, providedheat treatment conditions are favorablefor segregation to occur, an increase in thedegree of boron segregation on the grainboundaries could arise, which could alsoincrease their susceptibility to crackingduring welding.

2) The grain size of an alloy can signifi-cantly influence the effectiveness of grainboundary liquid solidification by liquid filmmigration (LFM). This has been critically

discussed in the work of Nakkalilet al. (Ref. 31) and will only besummarized here. Nakkalil et al.(Ref. 31) evaluated the occur-rence of HAZ cracking in In-coloy® 903 that had a duplexgrain structure. A considerableamount of microfissures was ob-served on the grain boundariesof a large-grained alloy. How-ever, minimal cracking and anextensive occurrence of LFMwere observed in the materialwith fine recrystallized grainboundaries. It was argued that aneffective relief of the supersatu-ration of solutes in grain bound-ary liquid film by LFM canreduce the total solidificationrange of the liquid film. This caneffectively reduce the suscepti-bility of an alloy to HAZ crack-

ing. The velocity of LFM is dependent ongrain boundary curvature, which has beenreported (Ref. 32) to vary inversely withgrain size. Thus, a fine-grained alloy shouldpossess a substantial mean grain boundarycurvature that would enhance the solidifi-cation of grain boundary liquid film by LFMand potentially reduce the susceptibility ofthe alloy to HAZ cracking.

3) When HAZ begins to accumulatestrains due to welding stresses, grain bound-ary sliding is one mechanism that can oper-ate to accommodate the strains (Ref. 33).However, this would increase stress con-centrations at grain boundary triple pointsand could initiate microfissures. Once initi-ated, the microfissures can easily propagatealong the liquated grain boundary regions.A large grain size would cause a longer in-terface sliding, which would lead to largerstress concentrations at grain boundarytriple points (Ref. 29). This would increasethe potential for crack initiation at grainboundary triple points, and therefore, liquation cracking.

In wrought alloy components, a signif-icant increase in grain size will usually in-crease the susceptibility of the alloy to

HAZ cracking during welding. In the pres-ent investigation, the grain size of 718Plusincreased significantly with an increase inpreweld solution heat-treatment temper-atures ranging from 950° to 1150°C (Figs.3, 4), and so was the susceptibility of thealloy to HAZ cracking — Fig. 8. For ex-ample, SEM observations of the welds ofsamples that were air cooled after preweldsolution heat treatment at 950°C (this re-sulted in an average grain size ≈ 58 μm) re-vealed an absence of HAZ cracking —Fig. 5. However, significant HAZ crack-ing, the extent of which increased withtemperature, occurred when the alloy waspreweld solution heat treated at 1050° (≈ 120 μm grain size) and 1150°C (≈ 360μm grain size), respectively.

A variation in weld HAZ cracking wasalso observed with cooling rates in samplesthat were preweld heat treated at 1050°and 1150°C — Fig. 8. This variation can-not be sufficiently explained with grainsize considerations alone, since the grainsize of the alloy did not vary appreciablywith the cooling rates — Fig. 4. This sug-gests that though the size of the grains of718Plus may have an effect on its suscep-tibility to HAZ cracking during welding,the cracking is however influenced bygrain size and boron segregation. There-fore, it is concluded that both grain sizeand grain boundary segregation of boroncontributed to the observed variation inHAZ cracking of 718Plus with tempera-tures and cooling rates.

Summary and Conclusions1) Weld HAZ cracking in 718Plus was

significantly influenced by preweld solu-tion heat treatment temperature and cool-ing rate. No cracking was observed insamples that were heat treated at 950°Cand cooled either by iced water quenching,air cooling, or furnace cooling. However,HAZ cracking occurred in samples thatwere preweld heat treated at 1050° and1150°C, respectively. The degree of crack-ing in samples varied with cooling rates.

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Fig. 12 — SIMS micrographs showing grain boundary boron segregation in 718Plus (higher boron version ≈ 60 ppm) after preweld heat treatment. A — At950°C; B — at 1050°C; C — at 1150°C. All samples were air cooled.

Fig. 13 — Variation in weld HAZ total crack length (TCL) of718Plus (higher boron version ≈ 60 ppm) with preweld heat-treat-ment temperature. All samples were air cooled.

A B C

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2) In samples that were preweld heattreated at 1050° and 1150°C, respectively,the highest degree of cracking occurred inair-cooled samples. This was followed byfurnace-cooled samples, while the lowestdegree of cracking was found in samplesthat were quenched in iced-water afterpreweld heat treatment. At all the threecooling rates, the degree of cracking wasalways higher in samples preweld heattreated at 1150°C than those that were so-lution treated at 1050°C.

3) The occurrence of weld HAZ crack-ing in 718Plus is concluded to be associ-ated with nonequilibrium grain boundarysegregation of boron which depended onthe preweld solution heat-treatment tem-perature and cooling rates, and also thegrain size of the alloy prior to welding,which was also dependent on the preweldtreatment temperature.

4) It is concluded from this study thatcrack-free welds of 718Plus can be achievedby carrying out a preweld solution heattreatment at 950°C. At this temperature,both grain growth and nonequilibriumgrain boundary segregation of boron wereconsiderably minimized.

Acknowledgments

Financial support for this work wasprovided by the Natural Science and En-gineering Research Council of Canada(NSERC) and is gratefully acknowledged.The authors would also like to thank Bris-tol Aerospace Ltd. in Winnipeg for weld-ing the samples. One of the authors (O. A.Idowu) also acknowledges gratefully theawards of postgraduate scholarships byNSERC and the University of Manitoba.

Thanks are also due to Dr. Rong Liu ofthe Department of Geological Sciences,University of Manitoba, for his technicalassistance with SIMS analyses. Finally, wethank ATI Allvac for providing Allvac 718Plus alloy.

References

1. Cao, W. D. 2004. U.S. Patent No.:6,730.254 B2.

2. Cao, W. D., and Kennedy, R. L. 2004. Su-peralloys 2004. K. A. Green, T. M. Pollock, H.Harada, T. E. Howson, R. C. Reed, J. J. Schirra,S. Walston, eds. The Minerals, Metals and Ma-terials Society, pp. 91–99.

3. Idowu, O. A., Ojo, O. A., andChaturvedi, M. C. 2007. Mater. Sci. Eng. A454–455, pp. 389–397.

4. Vishwakarma, K. R., and Chaturvedi, M.C. Mater. Sci. Tech. Article in Press.

5. Huang, X. 1994. PhD thesis, The Uni-versity of Manitoba, pp. 1–295.

6. Analysis Techniques: SIMS, Cameca,www.cameca.com.

7. Cao, W. D. 2005. Superalloys 718, 625, 706and Derivatives. E.A. Loria (ed.) The Minerals,Metals and Materials Society, pp. 165–177.

8. Donachie, M. J., and Donachie, S. J.2002. Superalloys — A Technical Guide, 2nd ed.,ASM International, pp. 225, 226.

9. Ojo, O. A., Richards, N. L., andChaturvedi, M. C. 2004. Scripta Materialia 50:641–646.

10. Ding, R. G., Ojo, O. A., andChaturvedi, M. C. 2006. Scripta Materialia 56: 5.

11. Vishwakarma, K. R., and Chaturvedi, M.C. 2005. Superalloys 718, 625, 706 and Deriva-tives. E. A. Loria (ed.). The Minerals, Metalsand Materials Society, pp. 637–647.

12. Kelly, T. J. 1989. Welding Journal 68: 44-s to 51-s.

13. Huang, X., Chaturvedi, M. C., andRichards, N. L. 1997. Acta Mater. 45: 3095–3107.

14. Chaturvedi, M. C. 2007. Materials Sci-ence Forum 546-549: 1163–1170.

15. McLean, D. 1957. Grain Boundaries in

Metals, Oxford University Press, Oxford, UK.16. Westbrook, J. H., and Aust, K. T. 1963.

Acta Mater. 11: 1151.17. Westbrook, J. H. 1964. Metall. Rev. 9: 415.18. Faulkner, R. G. 1996. Int. Mater. Rev. 41:

198.19. Tingdong, X., and Buyuan, C. 2004.

Prog. Mater. Sci. 49: 109.20. Faulkner, R. G., Jones, R. B., Lu, Z.,

and Flewitt, P. E. J. 2005. Phil. Mag. 85(19):2065–2099.

21. He, X. L., Chu, Y. Y., and Jonas, J. J.1989. Acta. Metall. 37(11): 2905–2916.

22. Chen, W., Chaturvedi, M. C., Richards,N. L., and McMahon, G. 1998. Metall. Trans. A29A: 1947–1954.

23. Williams, T. M., Stoneham, A. M., andHarries, D. R. 1976. Metal Science 19: 14–19.

24. Karlsson, I., Norden, H., and Odelius,H. 1998. Acta Metall. 36(1): 1–12.

25. Huang, X., Chaturvedi, M. C., andRichards, N. L. 1996. Metall. Trans. A 27A:785–790.

26. Chen, W., Chaturvedi, M. C., andRichards, N. L. 2001. Metall. Trans. A 32A:931–939.

27. Morimer, M.A. 1974. Grain boundary inengineering materials, 4th Bolton Landing Conf.J. L. Walter, J. H. Westbrook, and D. A. Wood-ford, eds. Claitors Publication Division, BatonRouge, La., p. 647.

28. Guo, H., Chaturvedi, M. C., Richards,N. L., and McMahon, G. 1999. Scripta Materi-alia 40 (3): 383–388.

29. Thompson, R. G., Cassimus, J. J., Mayo,D. E., and Dobbs, J. R. 1985. Welding Journal64: 91-s to 96-s.

30. Faulkner, R. G. 1981. J. Mater. Sci.16:373.

31. Nakkalil, R., Richards, N. L., andChaturvedi, M. C. 1993. Acta Metall. Mater. 41(12): 3381–3392.

32. Patterson, B. R., and Liu, Y. 1992. Met-all. Trans. A23A: 2481.

33. Thompson, R. G. 1993. Welding metal-lurgy of nonferrous high temperature materials.ASM Handbook Vol. 6: 567.

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Call For PapersAWS Detroit Section Sheet Metal Welding Conference XIV

May 11–14, 2010, Detroit, Michigan

The Sheet Metal Welding Conference Technical Committee is actively seeking abstracts related to joining technologies for thinsheet fabrication. Typical categories include

• Resistance Welding Processes • Thin-Gauge Structures/Materials • Arc Welding Advances• Tubular Structures • Advanced High-Strength Steels • Coated Materials• Adhesive Bonding • Laser Welding • Innovative Processes• Application Studies • Light-Weight Materials • Fastener Attachment• Hybrid Joining Methods • Process Monitoring and Control

A technical abstract in a format that is compatible with MS Word, along with a completed Author Application Form must be submit-ted to the Technical Committee Chairman by November 6, 2009. Abstracts to be considered must be of sufficient detail for a fair evalu-ation of the work to be presented. The paper must be related to sheet metal alloys and/or joining processes used in manufacturing of com-mercial products. It is not a requirement that your presentation be an original effort. Case histories, reviews, and papers that have beenpreviously published or presented will be considered as long as they are pertinent to the general interests of the conference attendees.

All abstracts will be considered by the Technical Committee. It is expected that the Committee's selections will be announced byNovember 20, 2009. Authors must submit a manuscript to the Committee by March 19, 2010. The Proceedings will be available toall attendees at the beginning of the Conference.

You may also download additional information and the Author Application Form at www.awsdetroit.org or www.ewi.org. The com-pleted Author Application Form and abstract should be sent to Menachem Kimchi, SMWC Technical Chairman, EWI, 1250 ArthurE. Adams Dr., Columbus, OH 43221, (614) 688-5153, FAX (614) 688-5001, mkimchi@ewi.org.

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