The effects of forging and rolling on microstructure in O …boehlert/GROUP/publications... · 2003-05-21 · The effects of forging and rolling on microstructure in O˙BCC Ti

Materials Science and Engineering A279 (2000) 118–129

The effects of forging and rolling on microstructure in O+BCCTi�Al�Nb alloys

C.J. Boehlert *Department of Mechanical Engineering, Johns Hopkins Uni6ersity, 3400 North Charles Street, Baltimore, MD 21218, USA

Received 1 July 1999; received in revised form 8 September 1999

Abstract

The effects of hot upset forging and hot pack rolling on microstructure of orthorhombic (O)+body-centered cubic (BCC)Ti�Al�Nb alloys was investigated. The starting materials were melted ingots of nominal compositions: Ti�25Al�25Nb(at.%),Ti�23Al�27Nb(at.%), and Ti�12Al�38Nb(at.%). Smaller cigar-shaped Ti�25Al�25Nb ingots were examined to understand theeffect of rolling preheat treatment on microstructure. It was found that super-transus preheat treatment results in large prior BCCgrains and surface edge cracking. For larger castings, forging and rolling procedures were carried out after heating the materialsbetween 932–1000°C. These temperatures were below the BCC-transus temperature for Ti�23Al�27Nb and Ti�25Al�25Nb andabove the transus for Ti�12Al�38Nb. This resulted in a significantly larger grain size for the as-processed Ti�12Al�38Nbcompared with the other two alloys. The Ti�25Al�25Nb alloy required the greatest forging and rolling loads, while the fully-BCCTi�12Al�38Nb alloy exhibited the best workability and required the lowest forging and rolling loads. This was related to thealloys’ aluminum contents and O-phase volume fractions. Sub-transus processing of the near Ti2AlNb alloys proved to be a viabletechnique for obtaining homogeneous microstructures containing fine O and BCC phases and lacking large prior BCC grains,which can be detrimental to the mechanical performance. © 2000 Elsevier Science S.A. All rights reserved.

Keywords: Titanium alloys; BCC phases; Orthorhombic phase

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1. Introduction

Since the discovery of the orthorhombic (O) phase ina Ti�25Al�12.5Al(at.%)1 alloy by Banerjee et al. [1],titanium aluminides containing the O phase (based onTi2AlNb) have been of interest for high-temperaturestructural applications, primarily because of their highspecific strength and stiffness as well as their creep andoxidation resistance. Recent results have shown that Oalloys offer major performance improvements overcommercial titanium alloys [2–10]. Like commercialtitanium alloys, the properties of O alloys dependstrongly on the microstructure and therefore the pro-cessing. To date, the relationship between the process-ing and microstructure of O alloys has beeninvestigated to a limited extent [1,5,6,8–16]. Much of

the thermomechanical processing has been based onprevious methodologies developed for conventional a–b titanium alloys and the intermetallic a2 titaniumaluminides. Early alloy development efforts focussed onextrusion, forging, or rolling operations on arc-meltedingots, with the primary intent being to characterize theequilibrium phases, phase transformations, and me-chanical behavior [5–8,10,11]. Later, Smith et al. [2–4]used foil processing to examine the development ofmicrostructure in Ti�22Al�23Nb and Rhodes et al. [12]studied microstructural evolution and crystallographictexture in Ti�22Al�23Nb and Ti�22Al�27Nb sheet andfoil products during hot rolling, cold rolling, and subse-quent heat treatment. More recently detailed studies ofthe development and control of microstructure duringforging of a Ti�22Al�27Nb alloy [9] and hot packrolling of a Ti�22Al�23Nb alloy have been performed[13]. These studies have focussed on a relatively narrowrange of O alloy compositions, containing the O, a2

(ordered hexagonal close packed), and body-centeredcubic (BCC) phases, which are being considered for

* Tel.: +1-410-5162876; fax: +1-410-5167316.E-mail address: [email protected] (C.J. Boehlert)

1 All alloy compositions are given in atomic percent.

0921-5093/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved.

PII: S0921 -5093 (99 )00624 -3

C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 119

metal matrix composite applications using foil-fiber-foilprocessing.

In this study, the processing-microstructure relation-ship for O alloys containing a wide range of composi-tions; namely Ti�25Al�25Nb, Ti�23Al�27Nb, andTi�12Al�38Nb alloys, was examined. Due primarily tothe high Nb content in such alloys, the two-phaseO+BCC regime has a wider temperature range thanthe a2+B2, a2+B2+O, and O-phase fields. Hencesuch alloys are termed ‘O+BCC’ alloys. The evolutionof microstructure, from melted ingot to forged pancaketo rolled sheet on the order of millimeters thick, wasexamined. In attempt to avoid large prior-BCC grains,which have shown to be detrimental to mechanicalbehavior [17–19], conservative thermomechanical pro-cessing techniques comprising non-isothermal forgingand hot pack rolling at relatively low processing tem-peratures were chosen. In addition to further develop-ing the understanding of microstructural evolutionduring processing of O+BCC alloys, this work de-scribes how processing temperature affects the ability tocontrol microstructural features, especially grain size,which strongly influence the mechanical behavior.

2. Experimental procedures

2.1. Materials and microstructural characterization

The studied alloys were grouped into two categories:near Ti2AlNb and Ti�12Al�38Nb. The target composi-

tions of the near Ti2AlNb alloys were Ti�25Al�25Nband Ti�23Al�27Nb. The initial portion of this studyinvolved examining the effect of processing parameters,and in particular the rolling preheat temperature, onmicrostructure using smaller ingots prior to attemptinglarge-scale deformation on larger castings. The termused to describe the smaller ingots is ‘cigar melts’because the dimensions, 150 mm in length and 30 mmdiameter, are similar to that of a cigar. Three 300-gcigar-melts were triple-melted using a vacuum inductionmelter at the Air Force Research Laboratory MaterialsDirectorate of Wright–Patterson Air Force Base, OH.The heats were formulated using elemental Ti, Al, andNb according to the stoichiometric mixture Ti2AlNband their measured compositions are listed in Table 1.The larger castings consisted of 175–500 mm longcylinders of 75 mm diameter. The large, near Ti2AlNbingots, whose compositions are listed in Table 2, were‘induction-skull’ melted at Flowserve (formerly DurironCorp.), Dayton, OH. Note that for Ti�23Al�27Nb, theTi, Al, and Nb contents adhered well to the targetcomposition, while for Ti�25Al�25Nb, the measuredcomposition was close to Ti�25Al�23Nb. The largeTi�12Al�38Nb ingot, whose measured composition wasclose to Ti�13Al�39Nb (see Table 2), was vacuum arcmelted at Pittsburgh Materials Technology Inc., Large,PA.

Several samples, diamond cut from each material,were analyzed for their constituent elements. The Ti,Al, Nb, and Fe contents were analyzed by means ofsolution X-ray fluorescence spectrometry and the datawere obtained using a Kerex Corporation Model 770Delta Analyst. The amounts of nitrogen and oxygenwere quantified using a Leco Corporation Model TC-136 oxygen/nitrogen analyzer. Chemical compositiondistribution between the different phases was measuredusing a Japan Electron Optics Ltd electron microprobeanalyzer (JEOL 733). Grain size (d) and phase volumefractions were determined quantitatively using NIHimage analysis software of digitized, high-contrast,back-scattered-detector (BSD) images taken using aLeica 360 field-emission scanning electron microscope(SEM). Transmission electron microscopy (TEM), per-formed using a JEOL JEM-2000FX electron micro-scope, and X-ray diffraction (XRD) were used toconfirm the presence of the different phases.

2.2. Procedures

2.2.1. Forging and rolling procedures for thecigar-melted ingots

Following melting, the cigar melts were cut by a wireelectron discharge machine (EDM) to a rectangulargeometry, measuring 125×25×20 mm, and coatedwith high-temperature glass for lubrication and protec-tion from the environment. They were then sealed in 6

Table 1Chemical analysis of the Ti2AlNb cigar-melted ingots and the corre-sponding as-processed sheetsa

Weight (ppm)Atomic percentMaterial

NbAlTi N Fe H O

Bal 24.8 24.5Ingot A 140 na 250na11024.225.4BalSheet A na 280350

Bal 26.6Ingot B 23.5 135110 na 650Bal na 230Sheet B 26.2 25.1 110 460

na na 290100Ingot C 25.124.6Bal25.0 110 530 na 890Sheet C Bal 24.6

a na, not available.

Table 2Chemical analysis of the large ingots

Atomic percent Weight (ppm)Material

OFeTi NNbAl

Bal 23.3 150 290 930Ti�25Al�25Nb 24.7Bal 1160Ti�23Al�27Nb 110020027.223.2

5752557039.213.2BalTi�12Al�38Nb

C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129120

Table 3Rolling procedures and parameters for the Ti2AlNb cigar-melted ingots

Ingot A B C

Preform Forged (50%) Forged (50%)Forged (50%)Can dima (mm3) 150×66×12140×60×12 122×66×12Pre-rolling – 1200°C/24 h—Heat treatmentRolling-preheat 815C/1 h/1000°C/0.25 h 1040°C/1 h/1000°C/0.25 h1040C/1 h/1000°C/0.25 hInterpass reheating 1000°C/2–3 min1000°C/2–3 min 1000°C/2–3 minIntermediate anneal – 1060°C/.5 h/1000°C/0.25 h1060°C/0.5 h/1000°C/0.25 hReduction per pass 10% 10%10%Roll speed (m2 min−1) 2.3 2.32.3Can dim (mm3) 412×64×3 419×74×3391×66×3Sheet dimb(mm3) 381×46×2.7 389×50×2.8353×48×2.8

a After forging, the cans were weld repaired, reevacuated, and sealed for use in rolling.b All rolled sheets were reheated at 1000°C for 3 min and then cooled slowly in vermiculite (�3°C min−1).

mm thick stainless steel cans and unidirectionallyforged according to a 2:1 ratio from 25 to 12.5 mm inair at a rate of 150 mm min−1. Preheat treatment of theingots included an isothermal soak at 1050°C for 15min followed by a 1000°C soak for 2 min prior toforging. After forging, the pancakes were slowly cooledin vermiculite2. The cans were weld repaired, re-evacu-ated and sealed at room temperature (RT), then sub-jected to different preheat treatments prior tosubsequent rolling. The rolling operations, performedusing a two-high laboratory mill, were chosen in orderto input more work into the forged pancakes andfurther homogenize the microstructures. Pack rollingwas used to minimize the temperature transients thatoccur during processing of thin sheet [13,20]. Therolling procedures for each pancake (labeled A, B, andC) are given in Table 3. Three different heat-treatmentand rolling schemes were chosen to produce a range ofmicrostructures. The heat treatments consisted of low-temperature (A), intermediate-temperature (B), andhigh-temperature (C) exposures in air followed by asub-transus preheat stage in the O+BCC regime(1000°C for 15 min) prior to rolling. Preheat treatmentof pancake A comprised an initial sub-transus preheat-ing stage (815°C for 1 h) to dissolve most of the BCCphase. The objective of this heat treatment and rollingcycle was to produce a fine-grained O microstructurewith a small volume fraction of BCC-phase particles.Preheat treatment of pancake B consisted of an initialpreheating stage near the BCC transus (1040°C for 1 h)to stabilize a larger BCC-phase volume fraction whilelimiting grain growth. Pancake C received a 24 hpre-rolling heat treatment at 1200°C, which is wellabove the BCC-transus temperature, in order to ho-mogenize a large-grained, fully-BCC microstructure.This was followed by the same preheat treatment sched-ule as that of pancake B prior to rolling. The rolling

was performed on a cold die and therefore after eachpass the pancakes were reheated at 1000°C for 2–3min. After finishing half of the rolling passes, each ofwhich resulted in a 10% reduction in thickness, pan-cakes B and C were annealed at 1060°C for 30 min andthen reheated at 1000°C for 15 min before continuingthe remaining rolling passes. As a result of the rollingoperation, significantly longer workpieces were ob-tained and the resulting sheets were approx. 3 mmthick. Fig. 1 compares the size of a forged and forged-and-rolled cigar-melted ingot.

2.2.2. Forging and rolling procedures for the largeringots

Based on the low-oxygen pick-up, the relatively lowdegree of edge cracking (see Fig. 2), the fine-grainedmicrostructure (see Fig. 3a), and the balance of RTtensile strength (1237 MPa) and elongation (5%) exhib-ited by sheet A [21], a similar rolling and forgingschedule was devised for the larger ingots. Forgingpreforms, 60 mm in diameter and 150 mm tall, wereEDM cut from each of the larger castings and coatedwith high-temperature glass then sealed in 6 mm thickstainless steel cans for protection from the environment.The can assemblies, made up of both the workpiece andthe outer can, were unidirectionally forged (3:1 ratio) toa final height of 50 mm in air at a rate of 150 mm

Fig. 1. Size comparison of (a) forged and (b) forged and rolledcigar-melted ingots.2 The estimated cooling rate was 3°C min−1.


Fig. 2. Low magnification photograph of the three as-rolled Ti2AlNbsheets.

not undergo uniform deformation, no surface crackswere observed on these workpieces after the initialforging. This is discussed in the results and discussionsection, which describes the added procedures used toprocess the Ti�25Al�25Nb workpiece. TheTi�23Al�27Nb and Ti�12Al�38Nb workpieces werethen re-canned and forged to 25 mm in a directionperpendicular to that of the original under identicalforging conditions. Similar to the first forging run, thecan assemblies were uniformly deformed. Again thework pieces were removed and the sides of the pan-cakes were cut parallel to a width of 75 mm.

After forging, the EDMed Ti�23Al�27Nb andTi�25Al�25Nb pancakes were re-canned and isother-mally soaked at 815°C for 1 h followed by a 982°Csoak for 15 min prior to rolling. The unidirectionalrolling steps consisted of several passes on a cold dieeach after a soak at 982°C for 5 min. For theTi�12Al�38Nb pancake, the soak temperature (932°C)was also identical to the forging temperature used forthis alloy. The reduction per pass for each pancake wasbetween 5–10%. The total reduction was approx. 60%and the total shear strain, calculated for both theforging and rolling procedures, was on the order ofthree. The rolling loads were recorded for each pass ona stripchart and measured between 50–70 tons depend-ing on the reduction and the alloy. After the final pass,the sheets were reheated at the respective soaking tem-peratures for 3 min and then removed from the furnaceand hung vertically for 1 min to promote creep straight-ening. They were then cooled in vermiculite. The finalthickness of the sheets was approx. 12 mm and nofurther effort to produce thinner sheets or foils wasmade.

3. Results and discussion

3.1. Microstructural e6olution

3.1.1. Cigar-melted ingotsA qualitative assessment of the effect of the rolling

preheat treatment and the intermediate annealing stepsis shown in Fig. 2, which depicts the three as-rolledsheets fabricated from the cigar-melted ingots. Notethat the greatest amount of edge cracking was exhibitedby the sample which underwent the most severe pre-rolling heat treatment (1200°C/24 h). Fig. 3a–c depictsthe corresponding as-rolled sheet microstructures. Foreach of the cigar-melted ingots the forging and rollingprocedures successfully broke down the O+BCCplatelet morphology, however some chemical banding,indicated by the alternating light and dark layers in theBSD SEM images of Fig. 3a–c, remained. Table 1 liststhe measured compositions of each of the ingots afterprocessing. The largest oxygen increase due to process-

Fig. 3. As-rolled microstructures of sheets (a) A (b) B, and (c) C.These BSD SEM images were taken from the thickness section andthe rolling direction is horizontal.

min−1. For the near Ti2AlNb alloys, prior to forgingthe can assembly was heat treated at sub-transus tem-peratures, 1000°C for 15 min followed by a 982°C soakfor 2 min. For the Ti�12Al�38Nb alloy, the can assem-bly was heat treated at 950°C for 15 min followed by a932°C soak for 2 min prior to forging. Note thatalthough the forging temperature for Ti�12Al�38Nbwas lower than that for the near Ti2AlNb alloys, it wasabove the BCC-transus due to the decrease in BCC-transus temperature with decreasing Al content [19,22].After forging, each pancake was cooled in vermiculite.The Ti�23Al�27Nb and Ti�12Al�38Nb workpieces,which underwent uniform deformation, were removedfrom the can assemblies and EDM cut to a height of 75mm. Unlike the Ti�25Al�25Nb workpieces, which did


Fig. 4. BSD SEM image of the sheet A microstructure after asub-transus heat treatment of 975°C for 100 h followed by waterquenching.

second phase was found to be a combination of globu-lar/elongated particles within the grains as well as lay-ers decorating the prior-BCC grain boundaries, see Fig.3c. The prior-BCC grain sizes of sheets A and B, whichdid not undergo the 1200°C heat treatment, were morethan an order of magnitude finer than that for sheet C.For sheet A, fine O- and BCC-phase particles werepresent throughout the microstructure. The as-pro-cessed sheet B microstructure was severely segregatedand contained particles which were highly concentratedwith Al and Nb (Ti�32Al�34Nb), see white particles ofFig. 3b. This segregation was determined to have beena result of the inhomogeneous starting ingot material[19,21], and it was also considered to be the reason whythe large difference in Nb content was measured for thepre- and post-processed materials (see Table 1). Due tothe fine, equiaxed microstructure, similar processingconditions, and in particular the rolling operations, asthose used for sheet A, which was devoid of largeprior-BCC grain boundaries and severe segregation,were chosen to process the larger-scale ingots. It isnoted that the chemical banding present in the as-pro-cessed sheet A was removed through sub-transus heattreatment, see Fig. 4, and a balance of RT tensilestrength and elongation and elevated-temperature creepresistance resulted [21].

3.1.2. Larger ingotsThe lower Al-containing ingots, Ti�23Al�27Nb and

Ti�12Al�38Nb, were uniformly deformed around theircenters during forging. The post-forged can assemblyfor Ti�12Al�38Nb is depicted in Fig. 5a and b. The

ing was measured for sheet C, which is expected to haveoccurred during the 1200°C exposure for 24 h, and as aresult the a2-phase precipitated at near-surface loca-tions, see the bottom of Fig. 3c. Also due to this severepreheat treatment, the as-rolled microstructure of sheetC retained large elongated grains, which were on theorder of 500 mm in length. The elongated grains in Fig.3c suggest that a substantial amount of deformation(approx. 4:1 thickness reduction) had been imposedduring rolling and that BCC-phase recrystallization didnot take place during deformation as large BCC grainswere present after the 1200°C/24 h heat treatment. The

Fig. 5. The (a) top and (b) side views of the can assembly after the first forging run for Ti�12Al�38Nb.


Fig. 6. The (a) top and (b) side views of the extracted workpiece after the second forging for Ti�12Al�38Nb.

circular cross-section of the workpiece can be clearlyseen in Fig. 5a, while a bulge in the stainless steel can ispresent at its midpoint, see Fig. 5b. Fig. 6a and b depictthe extracted Ti�12Al�38Nb workpiece after the secondforging run. The as-cast and forged pancake mi-crostructures for Ti�12Al�38Nb and Ti�23Al�27Nb aredepicted in Fig. 7a and b and Fig. 8a and b, respec-tively, and the O and BCC phases were identified usingboth TEM and XRD. It is evident in Fig. 7a and b thatthe BCC grain size was significantly reduced forTi�12Al�38Nb. TEM investigations revealed that theBCC phase was disordered (designated as b), a result ofthe low Al content [19,22]. For Ti�23Al�27Nb, most ofthe prior-BCC grain boundaries were serrated duringforging and although elongated O+BCC platelets re-mained, they were shorter and more ‘blockier’ thanthose in the as-cast material. Thus, the ingot mi-crostructure had been broken down and the orientationof the platelets was more random than that of theas-cast microstructure. It appeared that neither the Onor BCC phase had been recrystallized and that thestructure was merely a more highly wrought version ofthe as-melted ingot.

Ti�25Al�25Nb did not undergo uniform deformationduring forging and sheared at an approx. 45° angle

with respect to the forging direction. Photos of thepost-forged can assemblies for Ti�25Al�25Nb are de-picted in Fig. 9a and b. The circular shape of the top ofthe severely displaced workpiece is easily recognized inFig. 9a. The shearing resulted in a displacement of thestainless steel can which bulged on opposite sides at thetop and bottom of the can assembly, see Fig. 9b. Thisbehavior was reproducible as an identically configuredworkpiece and can exhibited similar behavior under thesame forging conditions. However, in this case, the

Fig. 7. Comparison of the Ti�12Al�38Nb (a) as-cast and (b) forgedmicrostructures. The forging direction was vertical. Note the reduc-tion in grain size after forging.


Fig. 8. Comparison of the Ti�23Al�27Nb (a) as-cast and (b) forged microstructures. The forging direction was vertical.

Fig. 9. The (a) top and (b) side views of the can assembly after the first forging run for Ti�25Al�25Nb. Note the non-uniform deformation.

forging press was halted immediately after the initialshear deformation was observed, see Fig. 10. The onsetof such instability and flow localization is typicallyobserved for materials which exhibit high degrees offlow softening and/or low strain rate sensitivities [23].High degrees of flow softening in conventional a�btitanium alloys have several main sources including (i)forging at low temperatures and high strain rates atwhich the flow stress and deformation heating-inducedsoftening are high or (ii) microstructural based soften-ing such as occurs during the break-down of coarse-grain lamellar microstructure [23]. In this case theformer source is more likely as both near Ti2AlNbingots contained similar Widmanstatten microstruc-tures, yet Ti�25Al�25Nb exhibited a greater compres-sive flow stress than Ti�23Al�27Nb, consistent withprevious results on these materials [19]. Thus the dis-placement rate of 150 mm min−1 or an approx. 0.017

Fig. 10. The Ti�25Al�25Nb can assembly after an interrupted forgingrun. The forging direction was horizontal. Note the non-uniformdeformation.


Fig. 11. The (a) top (b) bottom and (c) side views of the shearedTi�25Al�25Nb workpiece after the initial forging run. Note that theside depicted in (c) was EDM sliced after forging.

depicted in Fig. 11a–c. Along with the shearing defor-mation, severe cracking occurred along the top surfaceof the workpiece, see Fig. 11a. Such shearing andcracking have been observed in isothermal forging ofconventional a–b titanium alloys which also has beenrelated to high rates of flow softening and low values ofthe strain rate sensitivity index, m [23]. To preventnon-uniform deformation behavior, both the can as-sembly height and the reduction ratio were reducedprior to the second forging attempt. The slightly de-formed workpiece, taken from the interrupted forgingrun (see Fig. 10), was EDM cut into two pieces, whichwere subsequently re-canned to forge in the same direc-tion as that of the first. The can assembly prior toforging was approximately 62 mm tall and the intendedreduction ratio was 2:1. All other parameters, includingforging speed, remained constant. The forging deforma-tion was uniform for this run as depicted in Fig. 12aand b. Thus by simply reducing the can assemblyheight, uniform forging deformation was possible at982°C. To acquire a similar amount of deformation forTi�25Al�25Al as that for the other alloys, theTi�25Al�25Nb workpieces were again removed, EDMcut, and re-canned in preparation for a third forgingtrial. The third forging step was performed in a direc-

Fig. 12. The can assemblies of the two uniformly deformedTi�25Al�25Nb pancakes after the second forging runs.

s−1 compressive strain rate, which did not induce insta-bility and flow localization for Ti�23Al�27Nb, was toorapid for the stronger Ti�25Al�25Nb ingot at 982°C.This indicates the important influence of deformationheat-induced softening during forging. The extractedTi�25Al�25Nb workpiece after the initial forging run is


Fig. 13. Comparison of the Ti�25Al�25Nb (a) as-cast and (b) forged microstructures. The forging direction was vertical.

tion perpendicular to the first two and the reductionratio was again 2:1. The final height of the uniformlydeformed can assembly was approx. 25 mm. Similar tothat of Ti�23Al�27Nb, the ingot microstructure wassuccessfully broken down, see Fig. 13a and b. However,in this case the second-phase BCC particles were morespherical than those for Ti�23Al�27Nb. Also note thelower BCC-phase volume fraction for Ti�25Al�25Nb(Vf�0.05) compared to that for Ti�23Al�27Nb (Vf�0.30), compare Fig. 8 and Fig. 13, which is a result ofthe phase equilibrium [19,22]. Thus, the Ti�25Al�25Nbmaterial was workable at 982°C, though the ingotheight needed to be reduced from that used for theTi�23Al�27Nb and Ti�12Al�38Nb materials. It shouldbe noted that the applied loads for the initial forgingswere greatest for Ti�25Al�25Nb and smallest forTi�12Al�38Nb. The reason for the poorer workabilityof the Ti�25Al�25Nb ingot is believed to be due to thehigher Al content and the corresponding greater O-phase volume fraction. The poor workability of alloyscontaining high Al concentrations has been reportedpreviously [24], and microstructures containing greatervolume fractions of the undecomposed BCC phaseexhibit lower isothermal compressive flow stresses [15],which explains why the Ti�23Al�27Nb (which con-tained an intermediate BCC phase volume fraction)exhibited loads intermediate to those for Ti�25Al�25Nband Ti�12Al�38Nb. Ti�12Al�38Nb, which was fully-bat the forging and rolling temperature, was the mostamenable to forging and this is related to the low flowstress and excellent ductility of the b phase. On theother hand, the BCC phase of Ti�25Al�25Nb andTi�23Al�27Nb was ordered (B2) at the forging androlling temperatures [19,22] and this is expected to haveinfluenced the applied loads as well.

Similar to the forging procedures, the largest rollingloads and poorest workability during rolling were ex-hibited by Ti�25Al�25Nb. Single-pass rolling reduc-

tions of more than 5% were not achievable forTi�25Al�25Nb. The lowest rolling loads and best work-ability were exhibited by Ti�12Al�38Nb, which wascapable of single-pass rolling reductions of 10 percent.Fig. 14a and b depict the can assembly prior to rollingand the post-rolled workpiece for Ti�12Al�38Nb. Notethe substantial lengthening which resulted from rolling.Fig. 15a and b depict the can assembly prior to rollingand the post-rolled workpiece for Ti�23Al�27Nb, whileFig. 16a and b depict the post-rolled can assembly andthe extracted workpiece for Ti�25Al�25Nb. Unlike therolling procedures for the cigar-melted ingots, thesheets from the larger ingots exhibited almost no edgecracking. Thus, O+BCC alloys are quite amenable tointermediate-temperature rolling with the microstruc-

Fig. 14. The (a) can assembly prior to rolling and the (b) extractedworkpiece after rolling for Ti�12Al�38Nb.


Fig. 15. The (a) can assembly prior to rolling and the (b) extractedworkpiece after rolling for Ti�23Al�27Nb.

exception that the O and BCC phases were elongated inthe rolling direction. The rolled Ti�23Al�27Nb mi-crostructure also contained elongated O+BCC phases,Fig. 17b, yet they were less elongated, or more ovalshaped, than those for the forged material, see Fig. 8b.The development of this oval-like morphology, ratherthan the platelet morphology, may be attributed to thelarge amount of hot work during forging and rollingand lower amounts of growth due to sub-transus pro-cessing temperatures. In addition, recrystallization ofthe matrix phase was not apparent. Thus, due to thesub-transus rolling and forging steps, the platelet-likemorphology of the near Ti2AlNb ingots evolved to amore equiaxed microstructure. The largest averagegrain dimension of the near Ti2AlNb as-processed mi-crostructures was approximately 4 mm and a greaterBCC-phase volume fraction was evident forTi�23Al�27Nb compared to Ti�25Al�25Nb. Overall,the degree of homogeneity in the microstructure washigh after rolling for these sheets. Further rolling tothinner sheet or foil, through either hot rolling or coldrolling [12,14], would increase homogeneity en route toquality foils used in foil-fiber-foil processing. It is notedthat both sub-transus and super-transus heat treatmentsof the as-processed sheets resulted in homogeneousmicrostructures and a detailed description of the phaseevolution, which includes the temperature ranges forthe different phase regimes, is provided in Ref. [22]. Adetailed description of the creep and tensile behavior ofthese alloys is provided in Ref. [19] and [25].

For the super-transus processed Ti�12Al�38Nb, theas-rolled microstructure, see Fig. 17c, contained fully-bgrains elongated in the rolling direction. However, theoverall grain dimensions were over an order of magni-tude finer than those of the as-cast microstructure, seeFig. 7a, and the largest average grain dimension wasapprox. 30 mm. The highly elongated b grains suggestedan absence of recrystallization. Thus by performingforging and rolling at a relatively low temperature(932°C) within the single-phase b region (the b-transus

Fig. 16. The post-rolled (a) can assembly and (b) extracted workpiecefor Ti�25Al�25Nb.

Fig. 17. As-rolled microstructures of the (a) Ti�25Al�25Nb (b)Ti�23Al�27Nb, and (c) Ti�12Al�38Nb sheets. These BSD SEM im-ages were taken from the thickness section and the rolling direction ishorizontal.

tures containing larger BCC-phase volume fractionsproviding better workability, which is consistent withthe observations of Rhodes et al. [12]. Low magnifica-tion BSD SEM images of the three as-rolled sheetmicrostructures are depicted in Fig. 17a–c. TheTi�25Al�25Nb sheet microstructure, Fig. 17a, was quitesimilar to the forged material, see Fig. 13b, with the


temperature was determined to be 800°C [22]), followedby a short time hold at the processing temperature, theas-cast grain size was significantly reduced while main-taining a homogeneous single-phase microstructure.

4. Summary and conclusions

The microstructural evolution during hot upset forg-ing and hot pack rolling was investigated for O+BCCTi�Al�Nb alloys. The alloys examined wereTi�25Al�25Nb, Ti�23Al�27Nb, and Ti�12Al�38Nbnominally. The former two were grouped as nearTi2AlNb alloys. Smaller ingots were used to understandthe effect of rolling preheat treatment on as-rolledmicrostructure for Ti�25Al�25Nb. The largest grainsize, oxygen pick-up, surface a2 precipitation, and sur-face edge cracking was exhibited for the super-transuspreheated sheet. Each of the larger castings was pro-cessed to sheet at temperatures between 932–1000°C.Hot forging procedures initiated the breakdown of thelarge prior BCC grains of the ingots, while hot packrolling further reduced the grain size. TheTi�12Al�38Nb alloy was the easiest to deform, which isexpected to be a result of two main factors: (i) the flowstress of the near Ti2AlNb O+BCC alloys is signifi-cantly higher than that of the fully-b Ti�12Al�38Nb;(ii) the fully-b Ti�12Al�38Nb microstructure exhibitsexcellent ductility [19,26] and was processed above theb transus, while the near Ti2AlNb alloys were processedat sub-transus temperatures.

The following lists the conclusions of this work.1. The achievability of desired microstructures is

strongly dependent on the processing and heat-treat-ment schedules. Homogeneous, fine-grained (d�4mm) microstructures of near Ti2AlNb alloys wereproduced through sub-transus processing, which re-sulted in the break down of the prior BCC grainboundaries and platelet-like morphology. Mi-crostructures containing large prior BCC grainswere produced when super-transus processing heat-treatment temperatures were used. It is concludedthat fine-grained microstructures are possible onlywhen work is performed below the transus.

2. Super-transus forging and rolling of Ti�12Al�38Nbproduced intermediate grain sized (d�30 mm) fully-b microstructures. Due to the excellent ductility andlower yield stress of the low-Al b phase, this mate-rial exhibited better workability than the O+BCCnear Ti2AlNb alloys.

3. The BCC phase does not recrystallize during sub-transus hot forging and hot rolling. The absence ofrecrystallization leads to microstructures with elon-gated BCC grains.

4. Greater Al contents and O-phase volume fractionslead to greater forging and rolling loads necessary

for equal deformation. Ti�25Al�25Nb required areduced forging height than Ti�23Al�27Nb for uni-form deformation, suggesting a difference in thedeformation heat-induced softening behavior withincreasing Al content and corresponding higher O-phase volume fraction.

Acknowledgements

This research was performed at the Wright-PattersonAir Force Research Laboratory Materials and Manu-facturing Directorate under Air Force contractsF33615-91-C-5663 and F33615-C-96-5258 to UES, Inc.The author is especially grateful to Dr V. Seetharamanfor technical guidance and Drs B.S. Majumdar, D.B.Miracle, and S.L. Semiatin for helpful discussions. Theassistance of J. and T. Brown, T. Jones, and T. Goff ofUES., Inc. in conducting the cigar melting, forging, androlling experiments is gratefully acknowledged. Theauthor would also like to acknowledge the supportreceived from Johns Hopkins University during thewriting of this manuscript.

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