Comparison of the Microstructure, Tensile, and Creep Behaviorfor Ti-22Al-26Nb (At. Pct) and Ti-22Al-26Nb-5B (At. Pct)

C.J. COWEN and C.J. BOEHLERT

The effect of the addition of 5 at. pct boron on the microstructure and creep behavior of anominally Ti-22Al-26Nb (at. pct) alloy was investigated. The boron-modified alloy containedboride needles enriched in titanium and niobium, and because to these borides, this material wasconsidered to be a discontinuously reinforced metal matrix composite. These needle-shaped bo-rides made up to 2 pct of the volume and were up to 158-lm long and 22-lm wide. The effect ofboron on the mechanical properties was evaluated through in-situ creep testing and tensile testingat room temperature (RT) and 650 �C. Overall, the addition of 5 at. pct boron proved to bedetrimental to the tensile and creep behavior. The composite exhibited a brittle failure and lowerelongations-to-failure than the monolithic material. The in-situ tensile and creep experimentsrevealed that the deformation process initiated in the boride needles, which cracked extensively,and significantly greater primary creep strains and creep rates were exhibited by the composite.

DOI: 10.1007/s11661-006-9004-6� The Minerals, Metals & Materials Society and ASM International 2007

I. INTRODUCTION

THE addition of trace amounts of boron (B) toconventional titanium (Ti) alloys, such as Ti-6Al-4V (wtpct), has been shown to decrease the as-cast grain size byapproximately an order of magnitude.[1] This drasticreduction in the as-cast grain size leads to significantbenefits including increasing yield strength while reduc-ing or avoiding time spent on expensive and energy-intensive thermomechanical processing. The addition ofB produces titanium boride (TiB) whiskers within theconventional a + b microstructure.[1,2] The presence ofthese TiB whiskers within these microstructures has ledto substantial increases in room-temperature (RT)strength and Young’s modulus (E), while maintainingadequate elongation-to-failure (ef) values.

[3] Such bene-fits of the TiB phase are exhibited through either ingotmetallurgy or powder metallurgy processing routes.

Titanium-aluminum-niobium (Ti-Al-Nb) alloys differfrom conventional Ti alloys in that their constituentphases may include the ordered intermetallic ortho-rhombic (O) phase, based on Ti2AlNb; the orderedhexagonally-close-packed (hcp) intermetallic a2 phase,based on Ti3Al; and the body-centered-cubic (bcc) phase,whose ordering is dependent upon composition.[4–8] Ti-Al-Nb alloys have shown potential to exceed theelevated-temperature capabilities of conventional

a + b Ti alloys.[5,9–15] In particular, Ti-Al-Nb alloys ofnominal compositions close to Ti-22Al-26Nb (at. pct)have exhibited an attractive balance of RT and elevated-temperature properties.[16–22] However, relatively fewstudies have focused on investigating the effects of B

additions on the microstructure and mechanicalproperties of Ti-Al-Nb alloys.[23–28] The purpose of thiswork was to determine the effect of a nominally 5 at. pctB addition on the microstructural features of a nominallyTi-22Al-26Nb (at. pct) alloy (henceforth, all composi-tions will be given in atomic percent) and the impact thiscauses on the tensile and creep deformation behavior.The alloy composition was selected based upon studiesby Rowe,[19] Rowe and Larson,[20] and Smith etal.,[16,17,18] which suggested that this alloy exhibited anexceptional balance of mechanical properties comparedto other ternary O-based Ti-Al-Nb alloys. In order toexamine a fabrication route alternative to ingot metal-lurgy, powder metallurgy processing was performed onthe Ti-22Al-26Nb and Ti-22Al-26Nb-5B materials. Thislatter approach has the potential to produce significantcost reduction compared to some alloy and metal matrixcomposite (MMC) processing methods.[16] A limitedamount of work has been done to understand the effectsof powder microstructures on the mechanical perfor-mance of O-based alloys and MMCs.[16,17]

II. EXPERIMENTAL PROCEDURES

The Ti-22Al-26Nb and Ti-22Al-26Nb-5B materialswere produced in spherical powder form using inert gas(argon) atomization at Crucible Research (Pittsburgh,PA). The powders were size fractioned to 35 mesh, whichcorresponds to an average diameter of 500 lm. Eachpowder composition was placed into commercially pureTi cans and degassed at 300 �C for 24 hours in prepara-tion for hot isostatic pressing (HIP) at the Air ForceResearch Laboratory. The inner diameter of the cans was22.9 mmand the outer diameter of the cans was 25.4 mm.Two cans were made for each composition. The Ti-22Al-26Nb cans were filled with 0.135 and 0.065 kg of powderand the Ti-22Al-26Nb-5B cans were filled with 0.149 and

C.J. COWEN, Graduate Assistant, and C.J. BOEHLERT, Assis-tant Professor, are with the Department of Chemical Engineering andMaterials Science, Michigan State University, East Lansing, Michigan48824-1226, USA. Contact e-mail: boehlert@egr.msu.eduManuscript submitted August 8, 2006.

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0.76 kg of powder. The cans were subsequently hotisostatic pressed at 207 MPa and 1027 �C for 4 hours.The initial ramp up to temperature in the HIP vessel wasperformed at a heating rate of 18 �C/min from RT to 700�C and 28MPa. Thereafter, theHIP vessel was heated at arate of 17 �C/min to 1027 �Cand207 MPa.The cans thendwelled at a pressure of 207 MPa and a temperature of1027 �C for 4 hours. The HIP vessel was then cooled at arate of 24 �C/min to 300 �C and 0 MPa. The temperaturewas then decreased to RT rapidly. Both the temperatureand pressure profiles during these ramps were approxi-mately linear. A processing diagram, which outlines thetime-temperature-pressure profile used during the HIPconsolidation process, is presented in Figure 1. It is notedthat the materials remained below the bcc-transus tem-perature throughout the HIP process. No subsequentworking or heat treatment of these materials was per-formed after the HIP process. The material was removedfrom the cans using electron dischargemachining (EDM).

Bulk chemical analysis was performed using induc-tively coupled plasma optical emission spectroscopy andinert gas fluorescence. Samples for microstructuralevaluation were sectioned from the processed alloysand metallographically polished, using colloidal silica(60 nm) for the finishing polish. The microstructuralevolution from powder through the HIP process wasevaluated as both the powders and HIP materials wereimaged in a CAMSCAN 44FE scanning electronmicroscope (SEM). Backscattered electron (BSE) SEMimages were used to examine the phase distributions andtheir morphologies. The average chemical compositionof the boride phase in Ti-22Al-26Nb-5B was determinedusing a JEOL* JXA-8200 electron microprobe (EMP)

with high-purity standards, where over 20 boride needleswere analyzed. Phase volume fractions were determinedwith an accuracy of ±5 pct from approximately 20 BSEimages acquired at different magnifications for each

microstructure. ImageJ image analysis software wasused to calculate the phase volume fractions. Grain sizewas determined using the mean line intercept method.[29]

Attempts to validate the bcc-transus temperature weremade using differential thermal analysis (DTA). Endo-thermic and exothermic inflection points were takenfrom the DTA heating and cooling curves (both heatingand cooling were performed at 10 �C /min) to access thebcc-transus temperature.Specimens were machined into dogbone coupons via

EDM. Recast layers, which formed during EDM, wereremoved from the coupon edges using diamond polish-ing. Coupon dimensions measured 38 mm in length witha 1.9- by 3.6-mm reduced gage section. Tensile and creeptests were conducted over the temperature interval of23 �C to 650 �C, and selected samples were polished to ametallographic finish in order to reveal surface defor-mation features.The RT tensile tests were performed using an MTS

810 tensile testing machine equipped with a 100-kNload cell. Strain was measured during the tensile testswith an MTS model 632.31F-24 12-mm gage-lengthextensometer attached directly to the gage section of thesample. The RT tensile tests, performed at a strain rateof 10–3s–1, were performed in duplicate or triplicate, andthe average tensile properties were reported. For the650 �C tensile and tensile-creep experiments, flat dog-bone-shaped samples, with gage portion dimensions of3-mm wide by 2.5-mm thick by 10-mm long, were EDMcut. A slight reduction of 0.6 mm in the width of thesample near the middle of the gage section wasintroduced to ensure that failure occurred at thislocation. The specimens were glued to a metallicplaten and polished using an automatic polishingmachine. The 650 �C tensile and constant-load tensile-creep experiments were performed using a screw-driventensile stage (built by Ernest F. Fullam, Incorporated,Lantham, NY) placed inside the SEM chamber. Theload-displacement relationship was obtained live duringthe experiments using MTESTW version F 8.8e dataacquisition and control software (Admet, Inc.,Norwood, MA), which allowed for either load- ordisplacement-control testing. It is noted that the dis-placement data acquired during the experiments com-prised that of the sample as well as the gripping fixtures.Thus the displacement values reported in this studydo not represent the sole displacement of the reducedgage section of the sample. The 650 �C tensile exper-iments were conducted at a strain rate of approximately10–3s–1. The 650 �C tensile-creep specimens were loadedat rate of 3.7 N/s until reaching the desired creep stressafter which the load was held constant. Temperaturewas controlled using a constant-voltage power supply toa 1 cm diameter tungsten-based heater located justbelow the gage section of the sample. An open-bath,closed-loop chiller was used to circulate distilled waterat RT through copper tubes to prevent the tensile stagefrom overheating. A fine-gage type K thermocouple wasspot-welded to the gage section of each sample tomonitor the temperature. All samples were held attemperature for a minimum of 30 minutes prior to theaddition of load.

Fig. 1—Processing diagram outlining the time-temperature-pressureprofile used during the HIP consolidation process.

*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 38A, JUNE 2007—27

III. RESULTS

A. Microstructure

Table I lists the measured bulk chemical compositionsand Table II lists the measured phase volume fractionsand grain sizes for the materials investigated. Theaverage composition of the borides in the Ti-22Al-26Nb-5B composite was 46.3B-31.5Nb-22.2Ti. Back-scattered electron photomicrographs of the Ti-22Al-26Nb and Ti-22Al-26Nb-5B gas-atomized powders andhot isostatic pressed microstructures are presented inFigure 2. Large boride needles, up to 158-lm long and

up to 22-lm wide, were present in both the atomizedpowders and the HIP consolidated compact. Based onthe BSE images, the borides appeared to exhibit acompositional gradient and/or appears to be coring orfaulting from the boride-matrix interface into the centerof the boride (Figure 3). It is noted that although noporosity was evident in any of the HIP materials, themicrostructure of the Ti-22Al-26Nb-5B material wasnot homogeneous because some of the gas-atomized

Table II. Measured Phase Volume Fractions and Grain Size

Composition a2 Vp O VpBccVp

BorideVp

GrainSize* (lm)

Ti-22Al-26Nb 6 72 22 0 59±3Ti-22Al-26Nb-5B 20 66 12 2 23

*Grain size represents prior bcc grain size, except for Ti-22Al-26Nb-5B, where it represents all the equaixed phases present in themicrostructure.

Table I. Measured Chemical Compositions

NominalComposition

Ti(At. Pct)

Al(At. Pct)

Nb(At. Pct)

B(At. Pct)

O(ppm)

Ti-22Al-26Nb 50.7 21.4 26.8 <0.15 1127Ti-22Al-26Nb-5B 47.4 21.4 25.6 4.5 1653

Fig. 2—Backscattered electron SEM images of Ti-22Al-26Nb-xB powder and hot isostatically pressed alloys: (a) Ti-22Al-26Nb gas-atomizedpowder, (b) Ti-22Al-26Nb-5B gas atomized powder, (c) Ti-22Al-26Nb low magnification, (d) Ti-22Al-26Nb higher magnification, (e) Ti-22Al-26Nb-5B low magnification, and (f) Ti-22Al-26Nb-5B higher magnification. The a2 phase is the darkest phase, the O phase is the gray lathphase, the bcc phase is the light matrix phase, and the borides in (e) and (f) are the bright needles.

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powders were not fully sintered. Thus, a portion of thismaterial’s microstructure contained spherical powdercompacts that retained their circular boundaries, as isreadily evident in Figure 4. This microstructural inho-mogeneity is expected to have played a role in the low efvalues measured from this material, as will be discussed.

The bcc-transus temperature was 1130 �C for bothTi-22Al-26Nb and Ti-22Al-26Nb-5B, as estimated fromthe DTA results. These values are slightly higher thanthose measured for powder metallurgy–processedTi-22Al-26Nb, which estimated the Bcc-transus to bebetween 1120 �C and 1125 �C.[16,17]

B. Tensile Behavior

Figure 5 illustrates representative RT stress vs straincurves for each of the materials, and the average RTtensile properties are listed in Table III. Figure 6 illus-trates the representative stress vs displacement behaviorat 650 �C for each of the materials. As stated in theexperimental procedures, the in-situ experiments did notprovide an accurate measure of the actual tensile strainexperienced across the gage section of the sample.Therefore, the values for E, 0.2 pct yield strength (YS),and ef could not be calculated from the curves presented

in Figure 6. However, the ultimate tensile strength (UTS)values were measured from the stress vs displacementcurves. From these curves, it is evident that a significantlygreater displacement was achieved for the monolithicalloy. The average 650 �C UTS values were 934 and

Fig. 3—(a) and (b) High-magnification BSE SEM images of the bo-rides in Ti-22Al-26Nb-5B. What appears to be coring or faulting ex-ists across the boride cross section in (a) with evidence ofinterdiffusion displayed in (b). The cracking within the boride in (a)was due to tensile deformation and was not an artifact of processing.

Fig. 4—Incompletely homogenized gas-atomized powders evident inthe microstructure of the Ti-22Al-26Nb-5B alloy after HIP.

Fig. 5—Representative RT stress vs strain curves.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 38A, JUNE 2007—29

625 MPa for the monolithic and composite materials,respectively.Surface slip deformation observations from samples

tensile tested at 650 �C are presented in Figure 7. Wavyslip was evident in the bcc phase of the monolithicalloy, along with slip compatibility between the bcc anda2 phases and the bcc and O phases (Figures 7(a) and(b)). Preferential cracking at a2 boundaries is evidentfor the monolithic alloy (Figures 7(a) and (b)). Boridecracking without evidence of matrix slip is shown forthe Ti-22Al-26Nb-5B composite (Figures 7(c) and (d)).

C. Creep Behavior

Figure 8 illustrates that the composite exhibitedworse primary and secondary creep resistance than themonolithic alloy. Under creep testing conditions ofr = 400 MPa/T = 650 �C, the minimum creep rate forthe Ti-22Al-26Nb alloy was 5.60 · 10–8s–1 compared to4.20 · 10–7s–1 for the Ti-22Al-26Nb-5B composite. Thecomposite exhibited a creep rate almost an order ofmagnitude higher than the monolithic alloy.Figures 9(a) through (h) depict the surface deforma-

tion evolution during an in-situ creep experiment atr = 375 MPa/T = 650 �C for the Ti-22Al-26Nb-5Bcomposite. The sequence of deformation eventsobserved included boride cracking immediately uponreaching the creep stress of 375 MPa, at a total machinedisplacement of 0.24 mm, followed by more extensiveboride cracking with increased displacement. The crack-ing nucleated and grew preferentially within andaround the boride needles followed by crack growthinto the matrix. Crack nucleation also occurred at a2

Table III. RT Tensile Properties

CompositionModulus(GPa)

0.2 pctYS (MPa)

UTS(MPa) ef (Pct)

Ti-22Al-26Nb 131 1008 1059 1.2Ti-22Al-26Nb-5B 144 NA* 643 0.5

*NA: not available.

Fig. 6—The 650 �C tensile stress vs displacement behavior for eachmaterial. These experiments were performed using the tensile stageinside the SEM chamber, and accurate global strain measurementswere not available.

Fig. 7—Surface damage observations made after tensile testing to failure at 650 �C: (a) Ti-22Al-26Nb, (b) Ti-22Al-26Nb, (c) Ti-22Al-26Nb-5B,and (d) Ti-22Al-26Nb-5B.

30—VOLUME 38A, JANUARY 2007 METALLURGICAL AND MATERIALS TRANSACTIONS A

grains and grew into the bcc + O regions, but thisoccurred after cracking within the borides. Surface sliptraces were not evident in the matrix, as was exhibited intension for the Ti-22Al-26Nb monolithic samples(Figures 7(a) and 7(b)). Figures 10(a) through (c) depictselected photomicrographs obtained from a Ti-22Al-26Nb sample creep tested at r = 400 MPa/T = 650 �C. The prior-bcc grain boundaries were thelocus of damage during creep in the monolithic material.Cracks initiated preferentially within the a2 phaselocated at prior-bcc grain boundaries, and the crackpropagation path continued to follow the prior-bcc grain boundaries (Figure 10(c)) with increaseddisplacement.

IV. DISCUSSION

A. Microstructure

The addition of 5 at. pct B significantly altered themicrostructure of the Ti-22Al-26Nb alloy. Boride nee-dles were evident within the Ti-22Al-26Nb-5B compos-ite ranging in length up to 158 lm and width up to22 lm. The grain size in the Ti-22Al-26Nb-5B sampleswas significantly reduced when compared to the mono-lithic alloy (Table II). Prior-bcc grain boundaries werenot visible in the Ti-22Al-26Nb-5B composite, whilethey were evident in the Ti-22Al-26Nb alloy(Figures 2(c) through (d)). The materials were bothprocessed under the same pressure and temperatureconditions, but the Ti-22Al-26Nb-5B composite con-tained 20 pct a2 phase by volume compared to 6 pct byvolume for the Ti-22Al-26Nb alloy. The reduction ingrain size in the Ti-22Al-26Nb-5B composite whencompared to the Ti-22Al-26Nb monolithic alloy is

thought to be caused by both the increased volume ofa2 phase pinning grain growth and from the borideneedles pinning grain growth. The borides present in theTi-22Al-26Nb-5B composite occupied only 2 vol pctand were randomly distributed throughout the micro-structure.The EMP data from the boride phase in Ti-22Al-

26Nb-5B suggests a phase composition close to B2TiNb.This is significant due to the fact that B additions toTi-6Al-4V (wt pct) alloys have produced whiskers with acomposition of TiB or TiB2.

[1,2] Using the latticeparameters determined through density functional the-ory simulations by Trinkle (i.e., a = 0.6184 nm,b = 0.31262 nm, and c = 0.46833 nm),[31] selected areadiffraction patterns and EBSD Kikuchi patternsobtained from the boride phase present in a Ti-15Al-33Nb-5B alloy have been successfully indexed accordingto the space group and lattice site occupancies of theB27 orthorhombic crystal structure.[32] Because theboride composition measured in the Ti-15Al-33Nb-5Bmaterial (47.9B-29.5Nb-22.6Ti) is similar to that in theTi-22Al-26Nb-5B material (46.3B-31.5Nb-22.2Ti), thecrystal structure of the boride phase in the Ti-22Al-26Nb-5B composite is suggested to be B27 orthorhom-bic. Thus, it appears that the Nb is substituting for theTi in the borides on half of the titanium (c) Wyckoff sites(0.177, 0.25, and 0.123).

B. Tensile Behavior

Stiffening was provided by the addition of 2 vol pctreinforcing boride needles. Microhardness testing wasperformed, which indicated that the boride phase washarder than that of the matrix phase. It was proposedthat for ductile Ti alloy matrices, the borides carry ahigher load than the other constituent phases in themicrostructure, until a critical amount of crackingwithin the borides occurs, whereupon the load is thentransferred to the matrix.[3] This has been previouslyshown to occur through surface observations madeduring in-situ tensile deformation in a Ti-6Al-4V-1B(wt pct) alloy.[3] In the current study, the matrixwas not ductile at RT because the average ef value was1.2 pct (Table III). For the composite, an extensiveamount of cracking was observed to occur within theboride needles, with no evidence of slip occurringwithin the matrix. This is contrary to that observed inthe Ti-6Al-4V-1B (wt pct) alloy,[3] in that after boridecracking occurred, slip was observed to emanate fromthe boride/matrix interface. The propagation of slipthrough the a + b matrix was then observed to occurexpansively, which was believed to account for thelarge ef values obtained in the Ti-6Al-4V-1B (wt pct)alloy. In the current study, preferential cracking at a2boundaries and within a2 grains was evident (Fig-ures 7(a) and (b)), which is a typical deformationfeature observed within Ti-Al-Nb alloys.[33–35]

At both RT and 650 �C, the Ti-22Al-26Nb-5B com-posite displayed an inferior tensile performance whencompared to the monolithic alloy (Figures 5 and 6). TheTi-22Al-26Nb-5B samples produced a UTS 39 pct lowerthan the monolithic alloy at RT and 33 pct lower than

Fig. 8—Displacement vs time behavior for creep specimens undercreep testing conditions of r = 375 to 400 MPa/T = 650 �C in avacuum atmosphere (10–6 torr) using the Ernest Fullam, Inc. tensilestage.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 38A, JUNE 2007—31

the monolithic alloy at 650 �C. The presence of theboride needles, in addition to a finer grain size, would beexpected to strengthen the microstructure in a similarfashion as found for Ti-15Al-33Nb-5B and Ti-6Al-4V-1B (wt pct).[3,28] The detriment in strength is believed tobe caused by the addition of the large boride needles,which initiated fracture and exhibited a brittle response.Hypereutectic B compositions have resulted in largeborides and low ef values for conventional Ti alloys.[2,36]

In addition, incompletely sintered spherical compacts(Figure 4) are postulated to be preferred sites for

cracking and decohesion under the application of atensile stress and would be expected to decrease the efvalues. It is also noted that the Ti-22Al-26Nb-5Bmaterial exhibited a greater volume fraction of the a2phase than the Ti-22Al-26Nb monolithic alloy, which asmentioned previously would lead to more cracking atlower strains and lower overall ef values.Ti-22Al-26Nb has proven to be a ductile alloy if

thermomechanically processed into the appropriatemicrostructure.[16–20] The average RT ef value for theTi-22Al-26Nb monolithic alloy in the current study, 1.2

Fig. 9—Backscattered electron SEM images taken during a 650 �C creep experiment at 375 MPa for a Ti-22Al-26Nb-5B sample at the followingdisplacement levels: (a) 0 mm, (b) 0.24 mm, (c) 0.28 mm, (d) 0.31 mm, (e) 0.39 mm, (f) 0.41 mm, (g) 0.43 mm, and (h) 0.48 mm. This samplefailed at a displacement of 0.58 mm after 11 h.

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pct (Table III), was almost identical to that measured bySmith et al.[16] for a powder metallurgy HIP Ti-22Al-26Nb alloy. However, they noted that their worked foilof this composition exhibited an ef value of approxi-mately 5 pct. Thus, with additional hot work of the HIPmaterial, the incompletely sintered compacts maybecome homogeneous with the remaining microstruc-ture, thereby enabling greater ef values. The Ti-22Al-26Nb alloy exhibited a ductile response at 650 �C basedon the stress vs displacement curve (Figure 6).Emura et al.[27] have observed improved tensile

strength due to B addition in a Ti-22Al-27Nb (at.pct)/6.5 mass pct TiB material. The RT tensile prop-erties of their work are shown in Table IV. The powdermetallurgy material produced in their study was alsoprocessed by HIP gas-atomized powders, but differsfrom the material in this study in that the HIP compactwas then hot rolled, which resulted in a fine and moreuniform distribution of borides. Thus, it is believedthat hot rolling the Ti-22Al-26Nb-5B compact wouldresult in improved mechanical properties over themonolithic material by eliminating the detrimentaldefects shown in Figure 4 and refining the size of theboride needles and homogenizing their distributionthroughout the microstructure.

C. Creep Behavior

The Ti-22Al-26Nb-5B composite exhibited a mini-mum creep rate almost an order of magnitude higherthan the monolithic Ti-22Al-26Nb alloy. The worsecreep resistance of the composite when compared to themonolithic alloy is expected to be a result of borideneedle cracking, which occurred at low creep displace-ments and initiated the fracture process (Figure 9).Hagiwara and Emura[37] observed improved creepstrength with increased TiB volumes for aTi-6Al-2Sn-4Zr-2Mo (wt pct)/5-20TiB pct material pro-duced through blended elemental powder metallurgyvia HIP. Our previous work indicated an order ofmagnitude decrease in the minimum creep rates for acast and hot isostatically pressed Ti-15Al-33Nb-5Bmaterial compared to a cast and hot-rolled Ti-15Al-33Nb-0.5B material.[28] To the authors’ knowledge,these are the only studies documenting the effects ofthe addition of B on the creep of conventional and O-based Ti alloys, and they show improvements with theaddition of B. However, in these studies, the borideswere significantly finer in scale (50-lm maximumlength by 5-lm maximum width in Ti-15Al-33Nb-5B)and more homogeneously distributed throughout themicrostructure than the boride needles within the Ti-22Al-26Nb-5B material. Thus, the size and morphol-ogy of the boride needles play a role in the creepbehavior, with benefits gained only when they arepresent on a fine scale and homogeneously distributed.Once the borides crack (which was observed to occurimmediately upon reaching the creep stress), the load isexpected to be shifted to the matrix and the cracks mayact as stress concentrators, thereby increasing the creeprates.

Fig. 10—Backscattered electron SEM images taken during a 650 �Ccreep experiment at 400 MPa for a Ti-22Al-26Nb sample: (a) at adisplacement of 0.37 mm, (b) at a displacement of 041 mm, and (c) adifferent area at a displacement of 0.41 mm. The white arrows indi-cated selected prior bcc grain boundary cracks.

Table IV. RT Tensile Properties for a Ti-22Al-27Nb Alloywith and without 6.5 Mass Pct TiB[27]

Alloy and Processing RouteYS

(MPa)UTS(MPa)

ef(Pct)

Powder metallurgyTi-22Al-27Nb

701 848 6.8

Powder metallurgyTi-22Al-27Nb/6.5 mass pct TiB

1068 1260 2.3

Ingot metallurgyTi-22Al-27Nb

818 992 4

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 38A, JUNE 2007—33

V. SUMMARY

The effect of a 5 at. pct B addition on the micro-structure, tensile, and creep behavior of a Ti-22Al-26Nballoy was investigated. Both the monolithic and com-posite materials were formed by powder metallurgy andhot isostatically pressed into a compact shape. Borideneedles up to 158 lm in length and 22 lm in width, witha phase composition close to B2TiNb, were observed inthe Ti-22Al-26Nb-5B composite in a small volumefraction. The addition of 5 at. pct B to the Ti-22Al-26Nb alloy resulted in inferior tensile and tensile-creepproperties, and this was explained to be a result ofthe poor elongation-to-failure of this material, where theborides were preferential sites for crack initiation. Thedeformation evolution, characterized through in-situ650 �C tensile and tensile-creep experiments, indicatedthat the boride cracking initiated at low displacementlevels and grew extensively thereafter.

ACKNOWLEDGMENTS

This work was partially supported by the NationalScience Foundation through Grant Nos. DMR-0533954and DMR-030992 and by an American Society ofEngineering Education (ASEE) Air Force ResearchLaboratory (AFRL) Summer Faculty Fellowship(Contract No. F49620-02-C-0015) to CJB. The authorsare grateful to Sesh Tamirisakandala, Daniel Miracle,and Fred Yolton for helpful assistance and guidanceand Gerald Wynick and Ward Votava for their technicalassistance with the EMP analysis.

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