Introduction

Joining or repair of superalloy gasturbine hot section components canbe achieved using welding or brazing.For superalloys containing substantialgamma prime (γ ′) or gamma doubleprime (γ ″), welding of these alloys cancause excessive cracking in the heat-affected zone and fusion zone in addi-tion to distortion. Brazing, on theother hand, has advantages over weld-ing in terms of its capability in joininghard-to-weld superalloys and reducing

cost due to batch processing. Duringthe brazing cycle, the brazing alloymelts, joins the superalloys, and solidi-fies during cooling or via anisothermal diffusion process. The bulksuperalloys being joined stay in solidstate during brazing such that brazingalloys must have a lower meltingrange/temperature than that of super-alloys. Traditional nickel-basedbrazing alloys contain relatively largeamounts of a melting point depressant(MPD) such as boron (B), silicon (Si),and phosphorus (P) to reduce themelting temperature of nickel,

nickel-chromium, or cobalt matrix tosuitable brazing temperature rangesbetween 1000° and 1250°C (Ref. 1).Furthermore, these MPD elements arealso responsible for the wetting andflowing behavior of brazing alloys onthe superalloy substrate during thebrazing cycle. Boron, due to its highdiffusivity, is preferred where homoge-neous joint compositions are required. However, the use of traditional B,Si, and/or P-containing brazing alloyscan lead to the formation of brittle,hard phases (often containing Cr)within the joint and reductions of ten-sile, fatigue, and creep properties, inaddition to compromised corrosion re-sistance due to chromium depletion.Brittle borides (or carboborides), in ei-ther discrete or eutectic form, formedduring the brazing cycle have been ob-served to affect the mechanical prop-erties of the joint, noticeably the duc-tility and fatigue life (Refs. 2–4). Asshown in Fig. 1, the cycles to failureduring low-cycle fatigue (LCF) tests ofboth wide-gap braze (WGB) andnarrow-gap braze (NGB) joints aresubstantially reduced due to the pres-ence of brittle, B-containingintermetallic compound(s). The ductil-ity of NGB joint was also significantlylower than that of the base metal;crack initiation in NGB joint wasfound to be associated withintermetallic phases (Ref. 5). A processto improve the mechanical integrity ofthe braze joint/repaired area generallyrequires the use of a costly diffusion

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SUPPLEMENT TO THE WELDING JOURNAL, JULY 2014Sponsored by the American Welding Society and the Welding Research Council

Brazing of CMSX­4 with a Boron­ and Silicon­FreeNi­Co­Zr­Hf­Cr­Ti­Al Brazing Alloy

X. HUANG (Xiao.huang@carleton.ca) is with Dept. of Mechanical and Aerospace Engineering, Carleton University, Ottawa, Ont., Canada.

ABSTRACT In this study, a new boron- and silicon-free brazing alloy containing Ni-Co-Zr-Hf-Cr-Ti-Al was used to join CMSX-4 under three conditions. The resultantjoint interface microstructure and hardness of the constituents were evaluatedand are reported here. It was found that by using a combination of Hf and Zr asa primary melting point depressant, the amount of each element could beeffectively reduced while still achieving a relatively low liquidus for the brazingalloy. The new brazing alloy is able to successfully join CMSX-4 with completemetallic bonding. Although two types of interfacial Zr-rich intermetallic com-pounds were observed after a 40-min brazing cycle, their hardness values aresimilar to that of a superalloy substrate. With the application of pressure dur-ing brazing or postbrazing homogenization heat treatment, the intermetalliccompounds are substantially reduced. Furthermore, the formation of cellularγ/γ ′in the interfacial region is discussed in this paper.

KEYWORDS• Boron­ and Silicon­Free Braze Alloy • CMSX­4 • Narrow Gap Brazing

Combining Hf and Zr as a melting point depressant improved ductility and reduced thebrazing temperature to a range more compatible when joining superalloys

BY X. HUANG

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cycle. This approach may reduce theamount of brittle boride phases to acertain extent, but the properties ofthe superalloys may be affectedadversely. Alternatively, the use of Ni-Cr-Hf or Ni-Cr-Zr brazing alloy hasshown to alleviate some of the disad-vantages associated with B- or Si-containing brazing alloys and provideimproved high-temperature tensileproperties (reaching 76% of the basemetal’s yield strength and 93% of theductility [Refs. 6, 7]), as well as LCFand creep resistance. Close to 80% ofthe base metal’s fatigue life and 90%of the rupture strength have been re-ported with the use of Ni-Cr-Zr braz-ing alloy (Refs. 6, 7). The use of hafnium as a meltingpoint depressant for a nickel-basedbrazing alloy provides a differentapproach to produce a ductile brazingalloy with moderate brazing tempera-tures at or below 1240°C. Ductile braze alloys containing Ni-Hf-Cr, Ni-Hf-Co, and Ni-Hf-Mo wereinitially developed by Buschke and

Lugscheider (Ref. 8) as alternatives toboron- and silicon-containing brazingalloys. While these alloys showed goodductility, they require relatively highbrazing temperatures (1235°C) inorder to produce good wetting. Somesuperalloys may suffer incipient melt-ing at this temperature. Also, thejoints with Ni-Hf-Cr showed galvaniccorrosion when tested in an aqueoussalt solution (Ref. 9). An alloy with thecomposition of Ni-18.6Co-4.5Cr-4.7W-25.6Hf was used to join superal-loys at a brazing temperature of1240°C (Ref. 10). The W in the brazingalloy was found to increase themelting temperature of the alloy. Nickel-based boron-free brazing al-loys containing either Hf or Zr,reported in the literature, have compo-sitions of about 26–34 wt-% Hf or11–19 or 40–60 wt-% Zr (Ref. 11). Theintermetallic phases formed in the Ni-Zr eutectic alloys were found to besofter than borides, a good indicationof reduced brittleness of the joint.However, the addition of Hf or Zr in

some alloys was quite high, requiringlong diffusion heat treatment (up to 36 hours) at high brazing tempera-tures (up to 1320°C) in order for thebrazed joint to reach a compositionsimilar to a superalloy substrate. Withthe addition of hafnium or zirconiumalone in the nickel substrate, the low-est achievable melting temperaturesare 1190° and 1170°C, respectively. Another nonboron-containingnickel-based brazing alloy with man-ganese was developed for epitaxialbrazing of single crystal superalloys(Refs. 12, 13). The alloy containsmanganese in the range of 20–58 wt-%; in particular, Ni-(20–23)Ge alloyswere developed for brazing singlecrystal superalloys. Ge is a meltingpoint depressant that has also shownto assist the formation of γ /γ ′ in thebrazement (Ref. 14). A hardness testshowed the microhardness in thejoint to be equivalent to that of thebase metal (Ref. 15). However, beinga substitutional element in the Ni thematrix, the diffusion time for thepurpose of homogenizing the joint isquite lengthy (Ref. 16). To develop new brazing alloysystems that are devoid of B or Si, thecompositions of eutectic constituentsin precipitation-hardened superalloys

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Fig. 1 — LCF property of NGB and WGB joints with B­containingbrazing alloy in comparison to base metal X­40 (Ref. 4).

Fig. 2 — Heating (top) and cooling (bottom) heat flux curvesfor the braze alloy. A heating/cooling rate of 25°C/min wasused.

Table 1 — Braze Alloy Composition and Melting Range

Ni Co Zr Hf Cr Ti Al Solidus °C Liquidus °C

Development Alloy 1172.3 (heating) 1201 (heating)Composition (BF­1) Bal. 15 10 7.5 7.5 6.4 3.6 1137.6 (cooling) 1175.2 (cooling)Actual Alloy Ingot (BF­1) Bal. 13.52 9.07 7.2 6.94 6.9 3

Table 2 — Composition of CMSX­4 (wt­%)

Cr Co Al Ti W Mo Ta Re Hf Zr Ni

CMSC­4 6.4 9.6 5.6 1.0 6.4 0.6 6.5 2.9 0.1 – Bal.

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can serve as a reference point. A braz-ing alloy, Ni-10Co-8Cr-4W-13Zr,developed based on a eutectic compo-sition, was used to successfully brazesingle-crystal superalloys at a brazingtemperature of 1270°C (Ref. 17). Itwas reported that a eutectic Ni-basedbrazing alloy with a composition of3.1–8.2% Co, 6.8–38.5%Cr,0–12.6%Al, 0–11.5%Ti, 0–1.3%Mo,0–23.1Ta, 0–2.4%W, 0-5.1%Nb,0–1%Re, 0–0.4%Hf, and 0–0.6%Y hadsufficient fluidity during brazing oncethe eutectic temperature TE wasreached (Ref. 18). This research was initiated with theobjective to develop a boron-freemulti-element brazing alloy systemthat could be tailored to different sub-strate alloys and brazing temperatureranges while providing a joint compo-sition similar to that of thesuperalloys. The selection of alloyingelements and their percentages wasbased primarily on the followingstrategies: 1) eutectic composition(s)of binary Ni-M system (Ni-Hf, Ni-Zr,Ni-Al, and Ni-Ti); 2) elements that arepresent in current superalloy systems;and 3) compositions of eutectic phasesfound in superalloys. In the system developed in thisstudy, Hf and Zr were included, in asmaller percentage, in addition to

other elements (Co, Cr,Ti, and Al) found in theeutectic phases. In a con-trolled-compositionrange, Zr can improve theductility and stress rup-ture life of the brazed

joint. Hf, on the other hand, has theability to strengthen the γ ′ phase inprecipitation-hardened superalloysand improve the oxidation resistanceand coatability of the brazed joints. Aseries of alloys containing Ni, Co, Zr,Hf, Cr, W, Ti, and Al was cast and theliquidus and solidus of the alloys wereevaluated. Among these alloys, one ofthem (Table 1) exhibited the followingmelting and solidification characteris-tics (Fig. 2) and was used in this workto join CMSX-4, a second-generation,single-crystal superalloy. CMSX-4 is a nickel-based single-crystal superalloy known for its highstrength and creep resistance at veryhigh temperature. It is commonly usedin the gas turbine industry for turbineblades and vanes. The γ ′solvus isgreater than earlier generations of su-peralloys due to the removal of grainboundary strengtheners, and it is esti-mated to be as high as 1286°C (Ref.19). As the alloy is exposed to severe

operating conditions, there is always aneed to repair gas turbine componentsin order to reduce the overall life cyclecost. When selecting repair and refur-bishment technologies, it is to be real-ized that DS and SX high-strength su-peralloys are among the most difficultalloys for joining and repair. Duringrepair, several operations, includingmachining, grinding, and heattreatment, are known to cause recrys-tallization (Ref. 20) and formation ofcellular γ/γ′ (Ref. 21). The formationof cellular γ/γ′microstructure in thealloy has been observed to reduce thefatigue strength of CMSX-4 (Ref. 22).Reported in this reference, the fatiguecracks were found to be predominantlyinitiated from the cellular area and ar-rested at the interface between the cel-lular and unaffected regions. One ofthe reasons for such reduction infatigue strength is believed to be thelack of grain boundary strengtheningelements (B, Zr, and Hf) in SX alloys(Ref. 21). In one particular study, aboron-containing coating (called a“damage cure” coating) was used tosupply grain boundary strengthener toa heavily worked and heat treatedCMSX-10 containing localized cellularγ/γ ′and the result was a substantial

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Fig. 3 — Vacuum furnace brazing setup (samples were in­side graphite holder and weights were applied in one trial).

Fig. 4 — Ni­Co­Zr­Hf­Cr­Ti­Al braze alloy.

Table 3 — Summary of Sample Processing Conditions

Sample No. Condition

NGB1 1240˚C for 40 minNGB2 1240˚C for 40 min with pressure 50 lb/in.2

NGB3 1240˚C for 40 min + 1200˚C for 10 h for homogenization

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increase in fatigue life. The addition ofgrain boundary strengthener relocatedthe fatigue initiation site from cellularsite to internal casting defects. Thisfurther supports the beneficial effectsof using a brazing alloy containinggrain boundary strengtheners such asZr and Hf as included in the brazingalloy developed in this study.

Materials andExperimental Procedures

CMSX-4 bars were purchased fromAlcoa Howmet Research Center

(Michigan). The 0.625-in.- (1.6-cm-)diameter bars, with <001> directionalong the longitudinal direction of thebar, were supplied in the solution-treated condition (usually takes placeat temperatures between 1277° and1318°C (Ref. 19)). One of the boron-and silicon-free brazing alloys was castby Sophisticated Alloys, Inc., (Pa.). Theactual brazing alloy composition isprovided in Table 2. Foils were fabricated from the castbrazing alloy bar (0.75 cm diameter) toa thickness of 250 μm. Base materialsample buttons (1.6 cm diameter) weresectioned from the CMSX-4 bar afterthe secondary orientation was marked.

After gentle polishing of one side of theCMSX button surfaces to 600-gritfinishing, the samples were cleaned andassembled into a graphite sample holderwith braze alloy foil sandwiched inbetween. The two CMSX-4 buttonswere aligned in the sample holder tomaintain the secondary orientation.The samples inside the graphite jig wereplaced into a vacuum furnace as shownin Fig. 3. Three different vacuum braz-ing cycles were conducted under theconditions detailed in Table 3. The peaktemperature was determined to bebelow the solution heat treatment tem-perature for CMSX (1277°C) and to re-duce the likelihood of recrystalization,which occurs at temperatures above1240°C (Refs. 23, 24). For the initialbrazing, the following cycle parameterswere used: • Evacuate the furnace for 2 h andpurge with Ar. Evacuate again to 10–5

torr before turning heating elementson. • Heat from room temperature to1080°C at 10°C/min. • Hold at 1080°C for 1⁄2 h. • Heat to 1240°C at 25°C/min. • Hold at 1240°C for 40 min. • Furnace cool to roomtemperature.

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Table 4 — EDS Results of Phases in NGB1

Element Concentration, at.­%White Phase (A) in Fig. 5B Gray Matrix (B) Cubic­Shaped Phase (C) in Fig 5C

in Fig. 5B

Al 8.88 12.44 5.53Ti 2.42 4.19 4.63Cr 5.54 10.76 6.50Co 11.03 11.62 7.85Ni 62.75 58.95 43.13Zr 7.17 1.44 17.09Hf 2.21 0.00 8.32W 0.00 0.60 2.58Re 0.00 0.00 6.95

Fig. 5 — SEM images before etching (NGB1)with different phases labeled A, B and C..

Fig. 6 — SEM images of NGB1 interface in the etched condition.

A B

C

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After brazing, samples weresectioned, ground, and polished. Themicrostructure of the brazed sampleswas first observed in the as-polishedcondition using backscatter electron(BSE) microscopy. The percentages ofkey alloying elements in the brazedregion of three samples weremeasured by energy-dispersive spec-trometry (EDS). After EDS analysis,the samples were electrolyticallyetched in a solution of 35.6 mLH2SO4, 37.5 mL HNO3, and 9.4 mLH3PO4 at 7 V for 120 s. Themicrostructural analysis wasperformed using a TESCAN scanningelectron microscope (SEM). Lastly,the microhardness values of variousphases were measured using a Vick-ers hardness tester (ClemexTechnologies Inc., Longueuil, QC,Canada).

Results Visual inspection was conductedafter the completion of three brazingcycles. Indication of the brazing alloywas observed around the joint periph-ery of all samples, suggestingsufficient flow of the alloy duringbrazing (the brazing alloy foil accountsfor only one-fourth of the total mating

surface area of the single-crystal bar).In separate WGB trials utilizing the in-filtration method (brazing alloy placedon top of filler powder paste), thebrazing alloy had penetrated the entiredepth of a 7-mm joint clearance (Ref.25), further confirming the flowabilityof this alloy.

Microstructure of the Brazing Alloy

The microstructure of the as-castbrazing alloy is shown in Fig. 4. It con-tains a network of eutectic phases in aγ-Ni matrix.

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Fig. 7 — Elemental maps of the NGB1 interface. A — BSE SEM image; B — Hf; C — Zr; D — Al; E — Ti; F — Cr.

Fig. 8 — Elemental distribution across the braze joint (NGB1).

A B C

FD E

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Microstructure of NGB1

The as-brazed NGB1 in the as-polished condition exhibits a layer ofinterface containing two noticeablephases (A and C as labeled in Fig. 5Band C). In the substrate CMSX-4, nodendritic structure was observed,likely due to the complete solutiontreatment carried out at thatsupplier’s facility. The interface isabout 100 μm thick, including a bandon either side with a darker contrast.The nature of the phases in the inter-face was further analyzed using EDSwith the results, in at.-%, summarizedin Table 4. The light phase (A) is moreenriched in Zr and Hf, as compared tothe surrounding matrix (B), but lacksW, Ti, and Al. The cubic-shaped inter-metallic phase (C) contains elevatedamounts of Zr, Hf, W, and Re andreduced amount of Ni. Further evalua-tion of the interfacial structure wasconducted under the etched conditionwith images illustrated in Figs. 6 and7. A complex microstructure wasformed during the NGB process. Theinterface has predominantly Zr enrich-ment as seen from X-ray maps — Fig.7. This suggests that a homogeniz-

ation heat treatment is required inorder to create more uniformmicrostructure after the initial brazingcycle. The microhardness test,however, revealed that the hardnessvalues of the two Zr-rich phases (Aand C) were similar to the surroundinggray matrix phase (B) (Table 5) andmuch lower than that measured onborides or silicides (Refs. 2, 24). Semiquantitative elemental analysiswas also conducted to determine Hf,Zr, Al, and Ti distribution across theinterface — Fig. 8. Similarly, theanalysis confirmed that Zr was theonly element that exhibitednonuniform distribution after theshort brazing cycle.

Microstructure of NGB2

Using the same brazing cycleparameters as NGB1 but with the ad-dition of 50 lb/in.2 pressure, the NGBjoint had much less intermetallicphases. Also, these intermetallicphases became more intermittentwith an occasional intermetallicphase-free zone — Fig. 9. ComparingFigs. 5 and 9, the amount ofintermetallic phases in the interface

region decreased by at least 50%. Onthe remaining intermetallic phases,EDS analysis (Table 6) was carriedout to discern the nature of thesephases. Similar to that observed insample NGB1, the irregular-shapedlight phase (D) had increased Zr con-tent while the cubic-shaped phase (F)had elevated amounts of Zr, Hf, andTa compared to the surroundingregion E. X-ray mapping was alsoconducted to identify elemental dis-tribution. As illustrated in Fig. 10,only Zr is seen to be concentrated inthe joint interface region. The EDSline scan analysis results in Fig. 11also reveal higher Zr content in thejoint region; however, the percentageof Zr is slightly less than that in sam-ple NGB1. Additionally, Hf from thebrazing alloy has diffused away fromthe interface, reaching a uniform dis-tribution. It is believed that the ele-

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Fig. 9 — SEM images of NGB2 in the as­polished condition. A — low magnification of the joint interface; B — high­magnification micro­graph showing three different phases labeled D, E, and F.

Table 5 — Microhardness of Different Phases in the NGB Joint (Refs. 1–3)

NGB1 NGB2 NGB3

A B D E G Eutectic Boride Containing Phases327 HV 375 HV 329 HV 340 HV 334 HV 720

Table 6 — EDS Results of Various Phases inSample NGB2

Element Concentration, at.­%D E F

Al 8.20 12.60 8.11Ti 2.08 4.43 3.48Cr 5.66 5.64 6.47Co 10.75 10.34 9.26Ni 59.87 63.67 52.54Zr 6.57 0.00 9.67Hf 2.53 0.00 5.84Ta 1.83 1.78 4.63W 1.36 1.55 0.00Re 1.14 0.00 0.00

A B

DE

F

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vated Zr content in the joint is due tothe formation of Zr-rich intermetal-lic phases. The dissociation of thesewill have to take place beforediffusion of Zr can occur.

The hardness values at locations D andE are given in Table 5. Similarly, evenwith the presence of intermetalliccompounds, the hardness is quite lowcompared to a conventionally brazedjoint with B or Si. Sample NGB2 was subsequentlyetched to reveal the microstructure ofthe joint and substrate after a brazingcycle at 1240°C for 40 min. The inter-face still had some remainingintermetallic phases from the use ofthe brazing alloy. However, the two

CMSX-4 buttons were completlyjoined with no interfacialdelamination or voids — Fig. 12A.Also observed was the formation offine, irregular γ ′ on the substrateside of the joint due to the furnacecooling from brazing cycle, as shownin Fig. 12B. There is also anindication of eutectic γ/γ ′ formationin the joint — Fig. 12C. To further ex-amine the secondary orientationalignment after the brazing cycle, aseparate sample set was brazed at thesame temperature and pressure butfor a longer duration (24 h) to furtherremove the intermetallic phases. Asshown in Fig. 13, the alignment ofsecondary orientation between thetwo CMSX-4 buttons can be seen byobserving the cuboidal γ ′phases — Fig. 13.

Microstructure of NGB3

With the addition of ahomogenization treatment (fromNGB1 to NGB3), the interfacial inter-metallic phases reduced substantially.In fact, some of the joint region is in-distinguishable from the original sub-strate — Fig. 14A. This suggests thatthe homogenization treatment is ef-fective in removing the Zr-rich inter-metallic phases. The few remainingphases in the joint were further eval-uated. As shown in Fig. 14B,microstructure in this area contains afew light particles (H) and a large eu-tectic region (G), often termed cellu-lar eutectic γ/γ ′ in the literature, al-though strictly speaking, the forma-tion of the cellular region is through aperitectic reaction (Ref. 26). From theEDS compositional analysis (Table 7),region G is found to have essentiallyNi (Cr and Co) and Al. The light phase

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Fig. 10 — Elemental maps (NGB2). A — BSE SEM image; B — Hf; C — Zr; D — Al; E —Ti; F — Cr.

Table 7 — EDS Results for Phases in SampleNGB3

Element Concentration, at.­%I G H

Al 12.82 14.51 9.76Ti 1.25 2.09 1.13Cr 8.65 5.61 5.49Co 11.30 9.50 9.87Ni 62.39 64.10 63.44Zr 0.00 0.00 8.20Hf 0.00 0.00 0.00Ta 1.57 2.29 0.00W 2.02 1.90 0.00Re 0.00 0.00 0.00Re 2.11

Table 8 — Volume Percentage of Zr­ContainingPhase in Three NGB Joints

Volume Percentage of Zr­Rich IntermetallicPhases

NGB1 NGB2 NGB320.1 6.6 0.76

Fig. 11 — Elemental distribution acrossthe braze joint (NGB2).

A B

D

E F

C

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H (Fig. 14B) is a Zr-rich intermetalliccompound that has not been removedduring the homogenizationtreatment. Comparing Figs. 8 and 15,the composition across the joint after10 h of homogenization treatment is

essentially uniform, apart fromseveral remaining Zr-rich particles.Since these Zr-rich intermetallic com-pounds are very soft (Table 5), theyare not likely to embrittle the brazejoint.

Discussion

In this study, three different braz-ing processes were employed to joinSX CMSX-4 using boron- and silicon-free brazing alloy. The brazing

temperature of 1240°C was foundsuitable to join CMSX-4 withoutapparent interfacial defects. Althoughall three joints contained Zr-richintermetallic compounds, the totalvolume percentage was reduced bythe application of pressure or homog-enization heat treatment at 1200°Cfor 10 h. Table 8 shows the volumepercentages of intermetalliccompounds within a band of 75 μm inwidth along the interface. ComparingNGB1 and NGB2 to that produced byother researchers (Ref. 6) using simi-lar brazing temperature but a longercycle (Fig. 16), it is found theintermetallic compounds in NGB1and NGB2 are substantially lessevident. The effects of the pressure appli-cation for braze joints have been ex-amined by many researchers(Refs.16, 27–29). When beingapplied adequately (not to the extentof displacing the brazing alloy exces-sively from the joint or deform thejoints), it has the advantages todecrease the amounts of intermetal-lic compounds, increase jointstrength and ductility, improve inter-facial bonding, and the potential toreduce the brazing temperature andtime. An “ejection model” was usedto explain the improved jointmicrostructure and mechanical prop-erties when brazed with MBF(Metglas™ brazing foil) (Ref. 30). It’sessentially a mechanicalredistribution of the liquid metal(enriched in MPD, and hence, inter-metallic compound uponsolidification) first formed duringbrazing. This in turn reduces theamount of intermetalllic phasesformed in the solidified joint.Combining the application ofpressure and homogenization heat

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Fig. 12 — SEM images after etching (NGB2).

Fig. 13 — SEM images showing the alignment of secondary orientation using pre­scribedindexing line (1240°C for 24 h, arrows point to the interface).

Fig. 14 — SEM image of sample WGB3. Aand B — In the as­polished condition; C — inthe etched condition.

A

C

B

I

GH

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treatment, the intermetallic-contain-ing layer will be reduced further,likely leaving behind a homogeneousjoint microstructure. It was also observed that the dif-fusion of Hf out of the joint occurredmore readily (than Zr) after theinitial brazing cycle. And diffusion ofZr also ensued rapidly after homoge-nization heat treatment. Anotherfactor that needs to be taken intoconsideration during repair ofturbine components is the residualstress on the serviced parts. Stressedmaterials, particularly single-crystalsuperalloys, will likely develop local-ized cellular transformation γ/γ ′structure upon being heated (such asthat during brazing cycle) (Ref. 21).If it is not completely removedthrough heat treatment, the presenceof cellular structure is detrimental tothe mechanical properties of thecomponents. During braze repair, ifgrain boundary strengtheners (B, Zr,and Hf) can be incorporated into thejoint, they can alleviate the damagingeffect. In this study, cellular γ ′/γ ′structure was observed in all threejoints. While there have been reportson the reduced fatigue life in thepresence of cellular structure, thepresence of Zr and Hf in the joint,particularly along the cellular bound-aries, will be beneficial in ensuringthat the detrimental effect of cellularstructure can be minimized.Additionally, since a furnace coolingwas employed in this study, the sizeof cellular γ ′/γ ′ structure wasrelatively large due to the slow cool-ing rate. If the cooling rate can be in-creased by using an Ar quench, forexample, the cellular γ ′/γ ′ structurewill likely be much smaller (Ref. 31).

Conclusions and Future Work

In this study, a new boron- and sili-con-free brazing alloy based on Ni-Co-Zr-Hf-Cr-Ti-Al was developed andused to join CMSX-4 single-crystal su-peralloy under three conditions. It wasfound that using a combination of Hfand Zr as a melting point depressant,in addition to other elements, both el-ements in the initial brazing alloy canbe effectively reduced while stillachieving a relatively low liquidus. Thenew alloy successfully joined CMSX-4with complete metallic bonding, underall three conditions. Although twotypes of interfacial Zr-rich intermetal-lic compounds were observed after 40min of brazing cycle, their hardnessvalues were similar to that of thesuperalloy substrate. Furthermore,with the application of pressureduring brazing or postbrazing homog-enization heat treatment, theintermetallic compounds weresubstantially reduced. Future work will focus on testing ofjoint mechanical properties (high-tem-perature tensile, LCF, and creep) in ad-dition to joint coatability/coating-joint

interaction, since nowadays most ofthe refurbished turbine componentsare coated before being returned toservice.

The author would like to thank Car-leton University for the ResearchAward to support this work. The helpfrom Dr. F. Gao for sample preparationand SEM imaging is greatlyappreciated.

1. American Welding Society. 1991. BrazingHandbook, 4th ed. 2. Huang, X., and Miglieti, W. 2011. Widegap braze repair — a review. Journal of Engineer-ing of Gas Turbine and Power 134(10): 010801. 3. Schoonbaert, S., Huang, X., Yandt, S., andAu, P. 2008. Brazing and wide gap repair of X-40with IN 738. Journal of Engineering for Gas Tur-bines and Power 130(3): 32101–32110. 4. Henhoeffer, T., Huang, X., Yandt, S., andAu, P. 2011. Fatigue properties of narrow andwide gap braze repaired joints. Journal of Engi-neering for Gas Turbines and Power 133(9):92101.

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Fig. 15 — Elemental distribution of Al,Ti, Zr, and Hf after homogenizationtreatment (NGB3).

Acknowledgments

References

Fig. 16 — Comparison of NGB joints with that from literature. A — NGB1 and B — NGB2joints produced in this study without the application of pressure; C — NGB joint with hy­pereutectic Ni­Cr­Hf alloy, 1238°C/18 h (200×) (Ref. 6); D — NGB joint with hypereutecticNi­Cr­Zr, 1238°C/18 h (200×) (Ref. 6).

A B

C D

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