Part I. The Microstructural Evolution in Ti-Al-Nb O …boehlert/GROUP/publications/...Part I. The...

Part I. The Microstructural Evolution in Ti-Al-Nb O 1 BccOrthorhombic Alloys

C.J. BOEHLERT, B.S. MAJUMDAR, V. SEETHARAMAN, and D.B. MIRACLE

Phase transformations and the resulting microstructural evolution of near-Ti2AlNb and Ti-12Al-38Nb O 1 bcc orthorhombic alloys were investigated. For the near-Ti2AlNb alloys, the processingtemperatures were below the bcc transus, while, for Ti-12Al-38Nb, the processing temperature wassupertransus. Phase evolution studies showed that these alloys contain several constituent phases,namely, bcc, O, and a2; when present, the latter was in small quantities compared to the otherphases. The transmission electron microscopy (TEM), scanning electron microscopy (SEM), and X-ray investigations of samples that were solutionized and water quenched were used to estimate thephase fields, and a pseudobinary diagram based on Ti 5 50 at. pct was modified. The aging-transformation behavior was studied in detail. For solutionizing temperatures between 875 8C andthe bcc transus, the phase composition and volume fraction of the near-Ti2AlNb alloys adjustedthrough relative size changes of the equiaxed B2, O, and a2 grains. The aging behavior followed threedistinct transformation modes, dependent on the solutionizing and aging temperatures. Widmanstattenformation was observed when a new phase evolved from a parent phase. Thus, Widmanstatten Ophase precipitated within the B2 phase for supertransus fully B2 microstructures, as well as forsubtransus a2 1 B2 microstructures. Similarly, Widmanstatten B2 phase can form from a fully Omicrostructure, a transformation that has not been observed before. In the case of equiaxed O 1 B2solutionized and water-quenched microstructures, Widmanstatten O-phase formation occurred onlybelow 875 8C. For the subtransus-solutionized and water-quenched microstructures, a second agingtransformation mode, cellular precipitation, was dominant below 750 8C. This involved formation ofcoarse and lenticular O phase that grew into the prior B2 grains from the grain boundaries. A thirdtransformation mode involved composition-invariant transformation, where the fully B2 supertransus-solutionized and water-quenched microstructure transformed to a fully O microstructure at 650 8C.This microstructure reprecipitated B2 phase out of the O phase with continued aging time. For Ti-12Al-38Nb, Widmanstatten O precipitation remained the only transformation mode. It is shownthat subtransus processing offers flexibility in controlling microstructures through postprocessingheat treatments.

I. INTRODUCTION (designated as b). Henceforth, such alloys will be designatedas “O 1 bcc” alloys. In order to optimize properties forTHE orthorhombic (O) phase (cmcm symmetry based specific applications, it is important to understand the phase

on Ti2AlNb) has similarities with the hexagonal close- composition, morphology, and structure-property relationspacked (hcp) a2 phase (Ti3Al, DO19 structure), yet differs of such alloys. This need formed the rationale for the workby the lattice arrangement of Nb with respect to Ti.[1,2] Alloys described herein.with a significant volume fraction of O phase have shown Several works have discussed the phase equilibria of Oattractive mechanical properties,[3–10] owing to the excellent alloys.[9–19] Several of these studies have focused on thecreep resistance of the O phase combined with some room- lower-temperature stability and aging transformation behav-temperature (RT) ductility of the O phase arising from the ior.[9,10,13,15–19] Past investigations revealed that alloy compo-existence of multiple slip systems. Compositions that have sition and the temperatures of processing, solutionizing, andreceived recent attention are Ti-25Al-17Nb, Ti-22Al-23Nb, aging all have significant influence on the mechanisms gov-Ti-25Al-25Nb, and Ti-23Al-27Nb.* Those containing Nb erning the transformation behavior. The metastable B2 phase

decomposes into a2 and/or O through three distinct transfor-* All alloy compositions are given in atomic percent.mation modes. The first transformation mechanism is Wid-

concentrations of 25 pct and higher show very little a2 phase manstatten precipitation of a2 and/or O that involves bothand a much larger O 1 bcc phase field, where the bcc a shear transformation as well as diffusional transport. Aphase may either be ordered (designated as B2) or disordered second decomposition mode is through a composition-

invariant mechanism, which rapidly transforms the super-transus parent B2 to either a metastable a2 phase at low Nbcontents or metastable O at higher Nb levels.[13,18] A thirdC.J. BOEHLERT, Postdoctoral Fellow, is with the Department of

Mechanical Engineering, Johns Hopkins University, Baltimore, MD 21218. decomposition mode is through discontinuous precipitationB.S. MAJUMDAR and V. SEETHARAMAN, Senior Scientists, are with of a2 1 bcc or O 1 bcc; the resultant cellular microstructureUES, Inc., Dayton, OH 45432-1894. D.B. MIRACLE, Research Group replaces fine intragranular matrix precipitation.[6–9,19,20]Leader, is with the Air Force Research Laboratory, Wright-Patterson AFB,

Independent of the aging transformation, the B2 phase tendsOH 45433-7817.Manuscript submitted March 26, 1998. to undergo a composition-induced disordering resulting in

METALLURGICAL AND MATERIALS TRANSACTIONS A U.S. GOVERNMENT WORK VOLUME 30A, SEPTEMBER 1999—2305NOT PROTECTED BY U.S. COPYRIGHT

Table I. Chemical Analysis of the Near Ti2AlNb Alloysthe disordered b structure at low temperatures.[17–19,21] Thisis attributed to the decreasing-Al and increasing-Nb content

At. Pct Wt. (ppm)of the B2 phase in the ternary system as the temperatureMaterial Ti Al Nb N Fe Odecreases. The temperature and B2 phase compositional

dependence of these transformation modes has yet to be Ti-23Al-27Nb bal 23.2 27.2 200 1100 1160Ti-25Al-24Nb bal 25.4 24.2 110 350 280determined, and the low-temperature stability of these alloysTi-25Al-23Nb bal 24.7 23.3 150 290 930remains a concern.Ti-12Al-38Nb bal 13.2 39.2 70 255 575In this work, Ti-Al-Nb alloys containing approximately

Ti-50 at. pct were studied. The Nb/Al ratios were unity andhigher: nominally, Ti-25Al-25Nb, Ti-23Al-27Nb, and Ti-12Al-38Nb. At the start of the investigation, available iso-

temperature and then unidirectionally forged to an approxi-therms suggested that such a range of compositions mightmately 3:1 ratio. After the initial forging, the workpiecesallow a systematic investigation of the relative effects of thewere EDM cut, recanned, and forged in a direction perpen-O and bcc phases on the mechanical properties withoutdicular to the original forging direction under identical condi-significant complexity arising from the presence of a thirdtions. The canned material was then isothermally soaked atphase a2. Also, these O 1 bcc alloys are structurally attrac-815 8C for 1 hour, followed by a soak at the rolling tempera-tive, since alloys containing any significant amount of theture for 15 minutes prior to rolling. Unidirectional multipassa2 phase suffer from ductility problems, arising primarilyrolling was carried out with interpass reheating at the rollingfrom the single ^a& slip of the a2 phase. Details of thetemperature for 5 minutes. The reduction per pass rangedmechanical properties of these O 1 bcc alloys are availablebetween 5 and 10 pct, and the total reduction per pancakein Reference 9.was approximately 60 pct. The total effective true shearThe objectives of this work were twofold: (1) to identifystrain was calculated to be on the order of 3. After the finalthe temperature ranges of the separate phase regimes forpass, the assemblies were reheated at the rolling temperatureO 1 bcc alloys and (2) to understand the compositional andfor 3 minutes and then cooled slowly in vermiculite. Thetemperature dependence of the separate aging transformationestimated cooling rate was 3 8C/min. The can was thenmodes and associated phase morphologies. Whereas mostremoved. The final thickness of the sheets was approxi-past studies have concentrated on processing O 1 bcc alloysmately 10 mm. Additional details on processing, experi-in the supertransus B2 or b phase field, the approach in thisments, and microstructural and mechanical results arework was to process the alloys in the lower-temperatureavailable in Reference 9.subtransus domain. As will be shown in this article, this

Samples for the phase evolution and aging study wereapproach provided significant flexibility in controlling thediamond-cut from the rolled sheet, then cleaned and wrappedresultant microstructures through postprocessing heat treat-in tantalum foils prior to being encapsulated in quartz tubesments. Special emphasis was focused on the lower-tempera-that were backfilled at a low pressure of high-purity argonture aging transformation behavior and identifying the order/gas. The capsules were subjected to solution treatments atdisorder transition of the B2 phase. In addition to the descrip-different temperatures, followed by controlled cooling ortion of the microstructural evolution behavior, the opportuni-water quenching. Selected samples were then reheated at aties and limitations of using processing and heat-treatmentlower temperature, followed by water quenching, in orderschedules for obtaining stable microstructures that containto study the aging transformation behavior. The chemicaldifferent grain sizes, phase volume fractions, and morpholo-composition distribution between the different phases wasgies are discussed.measured using a JEOL* 733 electron microprobe analyzer.

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

The grain size and phase volume fractions were determinedThe starting materials of this work comprised induction quantitatively using NIH image analysis software of digi-

skull-melted ingots of Ti-25Al-25Nb and Ti-23Al-27Nb tized, high-contrast, back-scattered detector (BSD) imagesalloys acquired from Duriron Metals Inc. (Dayton, OH) and taken using a Leica 360 field-emission scanning electronvacuum arc-melted ingots of the Ti-12Al-38Nb alloy sup- microscope (SEM). Transmission electron microscopyplied by Pittsburgh Materials Technology Inc. (Large, PA). (TEM) was performed using a JEOL JEM-2000FX electronHenceforth, the Ti-25Al-25Nb and Ti-23Al-27Nb alloys will microscope. The TEM foils were prepared by ion millingbe referred to as near-Ti2AlNb alloys. The as-cast micro- mechanically dimpled samples. X-ray diffraction (XRD) wasstructure of the near-Ti2AlNb alloys contained large B2 used to confirm the presence of the different phases, andgrains, on the order of 500 mm, containing primarily O and differential thermal analysis (DTA) was performed to esti-some a2 platelets. The as-cast Ti-12Al-38Nb microstructure mate the bcc transus temperature.contained similarly large b grains. Hot-working proceduresconsisted of pancake forging and pack rolling performed at

III. RESULTS1000 8C, 982 8C, and 932 8C for the Ti-25Al-25Nb, Ti-23Al-27Nb, and Ti-12Al-38Nb alloys, respectively. Forging A. The Near-Ti2AlNb Alloyspreforms were cut from the ingots by an electrodischargemachine (EDM), and they were coated with a protective The average chemical compositions of the near-Ti2AlNb

alloys are listed in Table I. For the two alloys that werelayer of glass before being sealed in 6-mm-thick stainlesssteel cans. The cans were used to prevent oxidation during targeted for the stoichiometric Ti2AlNb composition, the Al

content was close to 25 at. pct, while the Nb compositionsprocessing. The assemblies were heat treated at the forging

2306—VOLUME 30A, SEPTEMBER 1999 METALLURGICAL AND MATERIALS TRANSACTIONS A

Table II. The Measured Average Grain Sizes, Phase Volume Percents, and Compositions for the Ti-23Al-27Nb Heat-Treated Microstructures

a2 bcc O dHeat Treatment Al Nb Vp Al Nb Vp Al Nb Vp (mm)

875 8C/100 h/WQ — — 0 pct 15.1 36.5 27 pct 24.8 24.2 73 pct 3.8875 8C/40 h/WQ/650 8C/336 h/WQ — — 0 pct na na 14 pct na na 86 pct 3.8900 8C/45 h/WQ — — 0 pct 17.0 33.8 63 pct 24.4 24.7 37 pct 5.8950 8C/45 h/WQ — — 0 pct 20.1 30.7 52 pct 25.2 23.6 48 pct 6.6950 8C/45 h/WQ/650 8C/373 h/WQ — — 0 pct na na 16 pct na na 84 pct 6.61000 8C/45 h/WQ 25.6 16.1 3 pct 18.8 31.7 43 pct 25.0 24.5 54 pct 5.01000 8C/45 h/WQ/650 8C/304 h/WQ 25.3 15.8 3 pct na na 16 pct na na 81 pct 5.01025 8C/200 h/WQ 25.6 15.5 12 pct 22.1 28.5 88 pct — — 0 pct 8.71025 8C/45 h/WQ/650 8C/450 h/WQ 25.6 15.4 12 pct na na 14 pct na na 74 pct 8.71090 8C/0.5 h/WQ — — 0 pct 22.0 27.3 100 pct — — 0 pct 177.01090 8C/0.5 h/WQ 650 8C/112 h/WQ — — 0 pct na na 20 pct na na 80 pct 177.0

WQ: water quenched; bold data represents the ordered bcc, B2, structure; na: not available.

varied between 23.3 and 24.2 at. pct (Table I). Henceforth, all the solution-treated microstructures represented in TableII contained only two phases, and none of the equiaxedthese alloys will be referred to by their measured composi-

tions: Ti-25Al-24Nb and Ti-25Al-23Nb. For Ti-23Al-27Nb, phases present were as fine as the a2 phase within the three-phase microstructure.the Ti, Al, and Nb contents indicated good adherence to the

target composition. The interstitial contents and, in particu- The phase volume percent (Vp) and chemical compositionsare also reported for each microstructure in Table II. Measure-lar, the oxygen content of the Ti-25Al-23Nb and Ti-23Al-

27Nb alloys were significantly higher than those of Ti-25Al- ment errors are estimated to be 64 pct for Vp and 61 pct forthe phase composition. Table II illustrates a wide variation in24Nb (Table I). Also, note the relatively high Ti-23Al-27Nb

alloy iron (Fe) content, which has proven to be significant Vp with heat treatment. With regard to the chemistry of eachphase, the O phase corresponded closely to the stoichiometricfor Ti-Al a2 alloys in terms of mechanical performance.[22]

The DTA results indicated the B2 transus to be 1070 8C Ti2AlNb composition. The chemistry of the a2 phase was alsorelatively constant at Ti-25Al-16Nb and in good agreementfor Ti-23Al-27Nb, 1030 8C for Ti-25Al-24Nb, and between

1051 8C and 1064 8C for Ti-25Al-23Nb. Confirmation of with the composition that has been obtained by Rhodes[17]

for a nominally Ti-22Al-27Nb alloy following 1000 hours ofthese transus temperatures was obtained using phase-disappearance studies, involving solutionizing at different exposure at 1038 8C, which is within the a2 1 B2 phase field.

The a2 phase was in a lath-type morphology in that study, duetemperatures, followed by quenching, and SEM and XRDconfirmed the disappearance of the O phase and any a2 to supertransus processing, rather than of the equiaxed type

present in the current investigation. The largest variation inphase. The DTA results were in relatively good agreementwith the cooling dilatometric transformations reported by elemental composition occurred for the bcc phase, where the

Al and Nb content varied according to the phase equilibriaRowe et al.[10]

The as-rolled sheet microstructures contained elongated requirements. Solution treatment at temperatures higher than875 8C favored the formation of the B2 phase (data printed inB2 and O and/or a2 grains, with fine O-phase precipitates

within the B2 grains. This precipitation occurred during the bold type); however, TEM studies revealed that the B2 phasedisordered below 750 8C, in agreement with previousgradual cool to RT. Table II describes the microstructures

obtained after different heat treatments for Ti-23A1-27Nb. works.[6,18]

Table III lists the measured microstructural parametersThe average spatial grain size (d ) was calculated using afactor of 1.74 times the planar grain size, determined by the for the Ti-25Al-24Nb and Ti-25Al-23Nb alloys. The phase

compositions and grain sizes followed similar trends to thatmean line-intercept method.[23] It is important to note thatthe d values listed in Table II do not represent only one of Ti-23Al-27Nb. For a given temperature in the O 1 B2

phase field, those alloys containing a greater Al concentra-phase, but these values represent the average equiaxed grainsize of the entire microstructure and include all the equiaxed tion exhibited lower B2 volume fractions. Concentrating

attention on the 875 8C heat treatment, note that the Ti-25Al-grain–shaped phases present. The sizes of the extremelyfine lath-type phases, which precipitated within the equiaxed 24Nb alloy exhibited a fully-O microstructure, whereas the

Ti-25Al-23Nb and Ti-23Al-27Nb alloys contained approxi-prior B2 grains after aging, were not included. The structuresof these phases will be described in more detail in Section mately 10 and 27 pct bcc phase, respectively (compare rows

3 and 9 of Table III and row 1 of Table II). The compositionIII–A–2. In general, the grain size increased with increasingsolution-treatment temperature (Table II). The exception to of the B2 phase, for the same solutionizing temperature, was

also different for the different alloys; compare the 950 8Cthis trend was the 1000 8C solution-treated microstructure,which contained a smaller average grain size than that of the treatment for Ti-23Al-27Nb and Ti-25Al-23Nb. Thus, both

the phase composition and volume percent of the bcc phase900 8C and 950 8C solution-treated microstructures (comparerows 3, 4, and 6 of Table II). The reason for this result is are alloy and temperature dependent, unlike the O and a2

phases that tend to maintain a fairly constant composition.believed to be the existence of extremely fine a2 phase at1000 8C, which restricted the B2-phase grain growth. Except Figures 1(a) through (d) depict portions of the 1025 8C,

1000 8C, 950 8C, and 875 8C ternary isotherms exhibitingfor the 1025 8C and 1000 8C solution-treated microstructures,

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 30A, SEPTEMBER 1999—2307

Table III. The Measured Average Grain Sizes, Phase Volume Percents, and Compositions for the Ti-25Al-24Nb,Ti-25Al-23Nb, and Ti-25Al-27Nb Microstructures

a2 bcc OAlloy and dHeat Treatment Al Nb Vp Al Nb Vp Al Nb Vp (mm)

Ti-25Al-24Nbas-rolled (subtransus) — — * na na 63 pct na na 37 pct 4.5875 8C/100 h/WQ — — 0 pct — — 0 pct 25.3 25.5 100 pct 17.1875 8C/100 h/WQ/900 8C/8 h/WQ — — 0 pct 19.8 29.5 18 pct 24.9 24.7 82 pct 17.1930 8C/100 h/WQ — — 0 pct 17.2 32.6 5 pct 24.9 24.8 95 pct 10.1975 8C/100 h/WQ — — 0 pct 21.9 28.4 21 pct 25.4 24.5 79 pct 10.4Ti-25Al-23Nb650 8C/166 h/WQ — — 0 pct na na 5 pct 24.8 24.9 95 pct 3.8875 8C/100 h/CC — — 0 pct 15.9 31.8 10 pct 24.5 22.9 90 pct na910 8C/8 h/WQ — — 0 pct na na 6 pct na na 94 pct 8.4910 8C/8 h/WQ/650 8C/218 h/WQ — — 0 pct na na 4 pct na na 96 pct 8.4950 8C/45 h/WQ — — 0 pct 20.8 27.5 25 pct 25.2 22.7 75 pct na950 8C/45 h/WQ/650 8C/223 h/WQ — — 0 pct na na 5 pct 25.2 22.6 95 pct na1010 8C/8 h/WQ 25.7 15.2 6 pct 23.8 24.8 94 pct — — 0 pct na

WQ: water quenched; CC: control cooled at 28 8C/min to RT; na: not available; *: a small volume of a2 phase may have been present;bold data represent the ordered bcc, B2, structure.

(a) (b)

(d )(c)

Fig. 1—Portions of the Ti-Al-Nb ternary system at (a) 1025 8C, (b) 1000 8C, (c) 950 8C, and (d ) 875 8C.

the nature of the tie-lines for the a2 1 B2, a2 1 B2 1 O, 1. Solution-treatment study of near-Ti2AlNb alloysFigure 2(a) illustrates the as-cast Ti-23Al-27Nb micro-and O 1 B2 phase regions. Included in the 1025 8C and

1000 8C isotherms are data taken from Reference 4. Note structure, consisting of a high volume fraction of lenticularO phase inside large prior B2 grains. After the subtransusthat, for the 950 8C and 875 8C isotherms, the O-phase

border remains relatively unchanged, while the bcc border processing, the microstructure evolved into a duplex micro-structure (Figure 2(b)) consisting of slightly elongated Oshifts to a significantly lower Al composition at lower tem-

perature. This is consistent with the work of Rowe and Hall[6] and B2 grains, with precipitation of fine and lenticular Ophase within the latter. In order to identify the temperatureand Rhodes.[17]


(a)

(b) (c)

(d )

( f )(e)

Fig. 2—Microstructures of the (a) as-cast and (b) as-rolled (rolling direction was vertical) Ti-23Al-27Nb alloy along with selected solution-treated andwater-quenched samples. Solution-treatment temperatures were (c) 1090 8C, (d ) 1025 8C, (e) 1000 8C, and ( f ) 900 8C.

range defining the various phase fields, the as-rolled material Although the B2 and a2 phases had the equiaxed morphologytypical of primary grains, several neighboring grains ex-was solutionized at different temperatures and then water

quenched. Figures 2(c) through (f) illustrate some of the hibited an orientation relationship (OR) of [11–20]a2//[1–11]B2, (0001)a2//(011)B2, which is the well-knownmicrostructures obtained. In these BSD images, the light,

gray, and dark regions correspond to the B2, O, and a2 Burger’s OR.[24] A selected-area diffraction pattern (SADP)depicting this relationship is shown in Figure 3(a). Work byphases, respectively. The figures show that solution treat-

ments at temperatures between 875 8C and the B2 transus Sagar et al.[25] indicates that, at 975 8C, the a2 phase under-goes dynamic recrystallization which destroys the Burger’sproduce equiaxed microstructures containing the B2 phase

along with the O and/or a2 phases. The presence of the fully- OR. The current observations are at variance with that find-ing and suggest that the a2 phase has its genesis in the parentB2 grains in the quenched microstructures was confirmed by

TEM. B2 phase that came from the as-cast microstructure and thatsome very fine a2 phase likely also existed in the as-castSolution treatments above the B2 transus resulted in large-

grained, fully-B2 microstructures, (Figure 2(c)). The XRD material in a platelet form (similar to the O phase). Thiswas then passed on to the rolled material in a fine, equiaxedresults, which included the superlattice (210) peak, con-

firmed the presence of the ordered B2 phase. Solution treat- mode through breakup of the lenticular form, followed bylimited growth of the islands of the a2 phase into an equiaxedments at temperatures between 1000 8C and the B2 transus

resulted in a2 1 B2 microstructures. A 1025 8C solution morphology. Although we did not observe any a2 phasein the as-rolled material, our investigation of the as-rolledtreatment exhibited a majority of B2 phase, approximately

88 pct by volume, with the remainder as a2 (Figure 2(d)). material was not detailed enough to rule out that it indeed


existed in very small amounts. It is noted that constant-cooling transformation studies have also not depicted signifi-cant volumes of a2 phase.[10]

Just below the two-phase a2 1 B2 field lies a three-phasea2 1 B2 1 O field characterized by a relatively narrowtemperature range (975 8C to 1000 8C). A typical three-phase a2 1 B2 1 O microstructure is shown in Figure 2(e),where the B2 phase was continuous. Note that O phaseformed preferentially as a rim around the a2 phase (inset inFigure 2(e)), suggesting once again that some a2 phase mayhave been present in the as-rolled microstructure. The ‘rimO’ phase is similar to that in three-phase microstructures ofother O alloys,[1,12,16] except that this phase was much thickerfor the current material.

The O 1 B2 field ranged between 875 8C and 975 8C.Solution treatments in this field, followed by quenching,(a)resulted in equiaxed O 1 B2 microstructures without anya2 phase; the volume fraction of O increased with decreasingsolution-treatment temperature in this field (rows 1, 3, and4 of Table II and rows 3, 5, 6 and 9, 10, 12 of TableIII). Figure 2(f) reveals a 900 8C heat-treated microstructurewhich contained approximately 63 pct O by volume (row 3of Table II). Once again, an OR was observed between theequiaxed O phase and the neighboring B2 phase. The ORdepicted in Figure 3(b), [2111]B2//[1–10]O, (110)B2//(001)O, is in agreement with past observations of O plateletsembedded in B2 grains.[1,13] However, the main differencehere is that the OR was observed for equiaxed grains, sug-gesting that the equiaxed O grains also have their genesis inthe parent cast microstructure. Thus, the as-cast O platelets,which have an OR with the B2 phase, likely break up andgrow into an equiaxed morphology during high-temperatureprocessing, still retaining the OR of the cast microstructure.Presumably, the significant amount of hot deformation wasunable to nucleate new grains or misorient the old O plateletssufficiently. It may be noted that, for the microstructures of(b)Figures 2(d) through (f), the B2 phase was verified by {100}superlattice reflections in the [001] B2 SADP (Figure 3(c)).

In addition to providing data on the temperature bound-aries for the respective phase fields, the solution-treatmentstudy revealed three important features regarding the phaseevolution of near-Ti2AlNb alloys. The first observation wasthat samples subjected to supertransus solution treatments,which were either control-cooled at 15 8C/min or waterquenched and then aged in the O 1 B2 region, were devoidof equiaxed constituents, and they retained large prior B2grains containing fine precipitates of O, along with a2 precip-itates for the control-cooled case. Thus, relatively fine equi-axed grains are only obtainable when hot work is performedbelow the transus temperature. For samples initially solu-tionized and quenched from the a2 1 B2 field and thenresolutionized at a lower temperature in the O 1 B2 region,the microstructures maintained the equiaxed a2 phase butcontained only Widmanstatten O within the prior B2 grains;i.e., equiaxed O grains were unable to nucleate. In addition,the rim-O phase was observed around the a2 phase, similar

(c) to that of the samples solution treated within the three-phaseFig. 3—SADPs taken from neighboring (a) a2 1 B2 and (b) O 1 B2 field (Figure 2(e)). These features are illustrated in Figuresgrains for 1025 8C/40 h/WQ and 900 8C/45 h/WQ solution treatments, 4(a) and (b). These observations suggest that the O phaserespectively, of Ti-23Al-27Nb. The ORs are [11-20] a2//[1-11]B2, (0001) tends to grow from the B2 phase in a lenticular fashion fora2//(011)B2, [-111]B2//[1-10]O, and (110)B2//(001)O. (c) Verification of

samples solutionized in the higher-temperature B2 and a2 1the ordered B2 structure in a 900 8C solution-treated sample depicting{100} superlattice reflections from a [001] zone axis. B2 fields, followed by aging in the O 1 B2 regime. On


(a) (b)

Fig. 4—Comparison of the Ti-23Al-27Nb (a) a2 1 B2 solution-treated (HT: 1025 8C/200 h/WQ) microstructure with (b) that for the solution-treatedmicrostructure after subsequent heating within the O 1 B2 field (HT: 1025 8C/200 h/WQ/925 8C/2 h/WQ). Note the a2 phase remained and WidmanstattenO precipitated in (b).

the other hand, processing at temperatures near the border solution treatment at 930 8C (compare rows 4 and 5 of TableIII). Also, slowly heating the fully-O microstructure frombetween the O 1 B2 and O 1 B2 1 a2 fields, followed

by solutionizing within the O 1 B2 field, results in fully- 875 8C to 900 8C resulted in an equiaxed O 1 B2 microstruc-ture, rather than lenticular B2 in prior O grains. Thus, theequiaxed O 1 B2 microstructures (Figure 2(f)).

Solution treatments at temperatures close to 875 8C pro- quenching treatment prior to resolutionizing appeared tofavor the lenticular nucleation and growth of the B2 phase.vide a transition regime regarding the growth of O phase in

O 1 B2 microstructures. Above 875 8C, changes in the Figures 6(a) and (b) which depict O 1 B2 microstructurescontaining similar phase-volume percents and different mor-volume fraction of the B2 and O phases, as dictated by

equilibrium requirements, are brought about by changes in phologies, show how heat treatments can be used to signifi-cantly alter the phase morphology. Because fully-Othe sizes of the equiaxed B2 and O grains. Thus, lower-

temperature solutionizing (but still above 875 8C) leads to microstructures were not obtained for the Ti-23Al-27Nb andTi-25Al-23Nb alloys, fine-gained, subtransus heat–treated,larger O grains at the expense of B2 grain size. The reverse

occurs for increasing temperatures. Below 875 8C, however, fully-lenticular O 1 B2 microstructures were also notobtained.the grain growth kinetics appear to be slow, and the volume

fraction requirements are accommodated by precipitation of In summary, through the proper selection of alloy compo-sition and thermomechanical processing steps (includingWidmanstatten O phase, as described in detail in Section

III–A–2. In addition, 875 8C solution treatments depict the heat treatments), a range of microstructures, including fully-O and fine-grained, fully-lenticular O 1 B2 microstructures,effect of nominal composition on microstructure, where

those alloys which contained higher Al contents exhibited are achievable for near-Ti2AlNb alloys.higher volume fractions of equiaxed O phase (Figures 5(a)through (c)). In fact, a fully-O microstructure, shown inFigure 5(a), was obtained for the Ti-25Al-24Nb alloy, while

2. Aging study of near-Ti2AlNb alloysthe other two alloys, which contained less Al, exhibitedsome equiaxed B2 grains. The lowest Al-containing alloy, a. Aging of subtransus-treated material

In order to investigate microstructural stability, agingTi-23Al-27Nb, exhibited the highest B2 volume fraction(Figure 5(c)). treatments were performed on selected solutionized and

water-quenched samples. Figure 7(a) depicts a Ti-23Al-The third observation regarding phase evolution is that afully-lenticular O 1 B2 microstructure was observed when 27Nb baseline equiaxed O 1 B2 microstructure solutionized

at 950 8C for 45 hours followed by water quenching. Figuresthe fully-O microstructure (Ti-25Al-24Nb solutionized at875 8C) was resolutionized at 900 8C, which is within the 7(b) through (d) illustrate the microstructures obtained after

13-hour isothermal aging treatments at temperaturesO 1 B2 region, followed by water quenching (Figure 6(a)).Figure 6(a) shows that, in this case, Widmanstatten precipita- between 650 8C and 850 8C, followed by water quenching.

All the aged microstructures were comprised of Widmanstat-tion of the B2 phase occurred within the fully-O microstruc-ture. Note that this Widmanstatten precipitation occurred on ten precipitates of the O phase within the B2 grains. The

reason this occurred is believed to be the low-energy configu-raising (rather than lowering) the temperature from the sin-gle-phase field. It is noted that this may not be an equilibrium ration for specific sets of planes in the O and B2 phases

(i.e., OR), but it is important to note that Widmanstattencondition, since the B2 phase volume fraction for this heattreatment (900 8C) was higher than that corresponding to precipitation can also occur on raising the temperature from


(a) (b) (c)

Fig. 5—Near Ti2AlNb microstructures, which have been solution treated at 875 8C: (a) Ti-25Al-24Nb, (b) Ti-25Al-23Nb, and (c) Ti-23Al-27Nb. Theseimages show the effect of alloy composition on O-phase volume fractions.

the single-phase field. The sizes of the primary O and priorB2 grains remained relatively unchanged.

The size and volume fraction of the Widmanstatten Oprecipitates depended on the aging temperature. The largestprecipitate size occurred at 850 8C and the finest size wasexhibited at 650 8C. These observations suggest that signifi-cant flexibility exists in the design of microstructures. Theprimary O grain size and volume fraction can be controlledby selecting the appropriate solution-treatment temperature,and the Widmanstatten precipitate size and volume fractioncan be controlled by selecting the appropriate aging condi-tions. Aged microstructures which were re-solution treatedat temperatures above and below 875 8C reconfirmed themorphological transition; those microstructures re-solu-tionized below 875 8C remained lenticular, while those re-solutionized above 875 8C contained only equiaxed grains.Thus, Widmanstatten precipitation starts at temperaturesbetween 850 8C and 875 8C for equiaxed, two-phase, subtran-sus solution-treated O 1 B2 microstructures in near-(a)Ti2AlNb alloys.

The microstructures obtained on aging at 650 8C and 7508C revealed discontinuous precipitation initiating at O/B2grain boundaries, indicated by the arrows in Figures 7(c)and (d). This grain boundary–related phenomenon, also des-ignated as a cellular reaction,[26,27] was investigated in detailusing long-term aging treatments. Figures 8(a) through (c)depict the microstructural changes occurring after aging a950 8C solutionized and quenched equiaxed O 1 B2 micro-structure at 650 8C for 24, 48, and 304 hours, respectively.These micrographs illustrate that aging at this temperatureresults in precipitation of coarse and platelet-like O 1 bccphases, which start at the grain boundaries and then advanceinto the prior B2 grains. The migration of the discontinuoustransformation interface resulted in the consumption of thefine, uniformly distributed Widmanstatten precipitates insidethe prior B2 grains. Since coarse precipitate islands were(b)not observed within the B2 grain interior, the coarsening

Fig. 6—Ti-25Al-24Nb heat-treated sheet microstructures: (a) 875 8C/100mechanism must be grain boundary–diffusion controlledh/WQ/900 8C/8 h/WQ and (b) 975 8C/100 h/WQ. These microstructures

depict the effect of heat treatment on morphology. rather than lattice-diffusion controlled. Thus, we designatethe coarsening process as discontinuous precipitation or a


(a) (b)

(c) (d )

Fig. 7—Effect of aging temperature on microstructure. Comparison of a Ti-23Al-27Nb solution-treated (a) (950 8C/45 h/WQ) microstructure with solution-treated and aged microstructures. The specimens were aged at (b) 850 8C, (c) 750 8C, and (d ) 650 8C for 13 h followed by water quenching. The arrowsindicated in (c) and (d) represent the initiation of discontinuous precipitation.

(c)(a) (b)

Fig. 8—Effect of aging time on microstructure. These Ti-23Al-27Nb specimens were solution treated at 950 8C and then aged at 650 8C for (a) 24 h, (b)48 h, and (c) 304 h followed by water quenching.

cellular reaction, which was first reported by Rowe.[7] After incomplete. A similar behavior was observed in the case ofan a2 1 B2 microstructure, which was solution treated at304 hours, some of the fine O precipitates remained, indica-

tive of the fact that the discontinuous precipitation was 1025 8C prior to aging at 650 8C for 610 hours. In this case,


Fig. 9—Dark-field images of the (a) O phase and (b) O 1 b phase platelets taken from the discontinuously precipitating region for a Ti-23Al-27Nb 9508C/45 h/WQ/650 8C/373 h/WQ heat-treated sample. (c) SADP taken from the same region depicting the corresponding OR: [01-1] b//[00-1] O, and (-211)b//(110)O.

the cellular products consumed a large volume fraction ofthe prior B2 grains, leaving a smaller volume fraction ofthe prior B2 grain with fine O precipitates. Thus, the transfor-mation behavior at this temperature is quite general.

The discontinuous precipitation was characterized by (1)coarse O-phase precipitation, and (2) disordering of the B2phase into the b. The latter was verified from convergent-beam electron diffraction patterns taken from the cellularregions. High-magnification dark-field images of suchregions are depicted in Figures 9(a) and (b), and the corres-ponding OR between the platelets, [01-1]b//[00-1]O,(2211)b//(110)O, is depicted in Figure 9(c). Misfit disloca-tions were observed within the discontinuously precipitatedregions and formed preferentially at O/b interfaces (Figure10). In contrast to the microstructures that contained B2grains prior to aging, the Ti-25Al-24Nb fully-O microstruc-ture (Figure 5(a)) remained equiaxed and fully-O after agingat 650 8C; i.e., there was no additional transformationinvolved. On the other hand, the O 1 B2 microstructuresdid not transform to fully-O microstructures after aging.

The b phase was not solely associated with the discontinu-ously precipitated regions. It was also found within the fine,

Fig. 10—Bright-field image of transformation misfit dislocations at O/buniformly precipitated regions. In some cases, the ordered interfaces for a Ti-23Al-27Nb 950 8C/45 h/WQ/650 8C/64 h/WQ heat-B2 phase was also retained in regions surrounding the fine- treated sample.precipitated O phase. Figure 11 depicts a transformed B2grain containing both b and B2 in the uniformly precipitatedregions. In general, discontinuous precipitation was exhib- precipitation, which initiated at grain boundaries and was

characterized by alternating platelets of the O and b phases,ited in very small volumes for those microstructures solutiontreated below 900 8C prior to aging. In those cases, the was observed in solution-treated and water-quenched sam-

ples which were aged at T , 750 8C. Although not quantified,majority of precipitation occurred uniformly throughoutthe bulk. there was clear evidence that samples which were solu-

tionized at higher temperatures and aged at lower tempera-The following summarizes the microstructural observa-tions from the aging study performed on subtransus-treated tures exhibited greater volume fractions of discontinuous

precipitation. The prior B2 regions, up to which the discon-samples. Fine Widmanstatten O phase precipitated uniformlywithin B2 grains for all the solution-treated samples which tinuous precipitation front had not advanced, contained Wid-

manstatten O precipitates with surrounding b and/or retainedwere aged at temperatures below 875 8C. Discontinuous


Fig. 11—A transformed B2 region containing the b, B2, and O phases fora Ti-23Al-27Nb 875 8C/40 h/WQ/650 8C/320 h/WQ heat-treated sample.

B2 phases. After aging at 650 8C for long times, fully-Omicrostructures remained equiaxed and fully-O, while O 1bcc microstructures remained O 1 bcc.

b. Aging of super-transus-treated materialAging of the supertransus-treated and water-quenched

material, which was fully B2 (as confirmed by XRD andTEM analysis), also exhibited grain boundary–related pre-cipitation. However, the precipitation behavior was differentfrom the subtransus material. Figures 12(a) through (c)depict the Ti-23Al-27Nb supertransus-solutionized micro-structure after aging times of 2, 9, and 100 hours, respec-tively, at 650 8C; refer to Figure 2(c) for the baselinesupertransus microstructure. Although SEM images of the

Fig. 12—Microstructures of a Ti-23Al-27Nb supertransus solution-treated2-hour-aged Ti-23Al-27Nb microstructure did not appear tosample (1090 8C/0.5 h/WQ), which had been aged at 650 8C for (a) 2, hbe much different than those of the unaged microstructure(b) 9 h, and (c) 100 h, followed by water quenching. The B2 phase rapidly(Figures 2(c) and 12(a)), XRD results indicated that signifi- transformed to O, as shown in the TEM image provided in the inset of

cant transformation indeed took place. In fact, after only 15 (a). O precipitates coarsened initially at (transformed-B2)/(transformed-minutes at 650 8C, the XRD peaks had shifted, as illustrated B2) boundaries, as indicated by the unlabeled arrows in (b) and then within

the bulk ((c)). The inset TEM image of (c) depicts the reprecipitation ofin Figure 13. Note that the 2u values of the peaks for theextremely fine B2 phase (white) within the O platelets (dark). In addition,aged samples are different from those of the B2 peaks fordiscontinuous precipitation is indicated by the arrow in the inset of (c).

the solutionized and unaged sample, indicating precipitationof O from B2. Several new peaks may be observed for theaged samples at 2u values expected for the O phase. High- observed, it is likely that some short-range diffusion must

also be involved at 650 8C. However, the martensitic typemagnification BSD images failed to reveal any composi-tional contrast, suggesting that the transformation of B2 to O phase was not stable, as is discussed subsequently.

Microscopy revealed that, for short aging times (t # 9O in the bulk, which occurred within a very short time at650 8C, is likely martensitic. A TEM image of the 2-hour- hours), aligned, fully-O platelets coarsened at transformed-

B2/transformed-B2 boundaries and grew into the trans-aged condition, provided as an inset in Figure 12(a), showsthe martensitic nature of the O phase. This is in agreement formed B2 grains, which were saturated with martensitic-

type O phase; note the dark regions indicated by arrows inwith observations made by Bendersky et al.[18] at 700 8Cfor a supertransus-treated Ti-24Al-25Nb alloy. The transfor- Figure 12(b). Each group of aligned platelets can be consid-

ered as a colony. With increased aging time, the colonymation is, however, complicated by the fact that if it werefully martensitic, then martensitically transformed B2 grains of platelets grew, similar to the growth of discontinuous

precipitation fronts that were observed for the subtransus-should have been observed in the as-quenched condition ofthe supertransus-solutionized material. Since this was not treated materials aged below 750 8C. However, unlike the


(a)

(a)

(b)

(b)

Fig. 14—SADPs identifying the (a) [001]B2 and (b) [100]O structureswithin precipitated regions of a 1090 8C solutionized Ti-23Al-27Nb sample,which was aged at 650 8C for 681 h.

separate cells of a colony. Diffraction patterns confirmedthe B2 and O phases in the reprecipitated regions (Figures14(a) and (b)); note that the B2 phase did not convert to bphase even after 681 hours at 650 8C. The OR between the

(c) phases was identical to that of the equiaxed, subtransus-solutionized microstructures. At some O/O cell boundaries,Fig. 13—XRD plots of intensity vs 2-theta for a Ti-23Al-27Nb (a) super-discontinuous precipitation resulted in a coarser platelet sizetransus solution-treated sample (1090 8C/0.5 h/WQ), which had been aged

at 650 8C for (b) 0.25 h and (c) 2 h, followed by water quenching. (inset of Figure 12(c)), similar to that of the subtransus-solutionized then aged microstructures. However, in thiscase, the discontinuous precipitation did not migrate far andwas not a dominant feature.subtransus microstructures, the platelets behind the migrat-

The following summarizes the microstructural observa-ing interface divided into separate cells containing distincttions from the aging study performed on the supertransus-boundaries. These cells coarsened with time and formed ansolutionized then water-quenched samples. The B2-to-Oelliptical shape, so that, for longer aging times, the micro-martensitic-type transformation occurred within 15 minutesstructure could be described as colonies containing severalat 650 8C, but the O phase was not stable. The O phaseequiaxed-shaped cells. Eventually, the separate coloniescoarsened initially at grain boundaries and then coarsenedimpinged upon each other, and, after 100 hours, the colonieswithin the prior B2 grains. The B2-phase platelets reprecipi-consumed nearly all of the prior B2 grains (Figure 12(c)).tated within the O phase.In addition to colony growth, extremely fine B2 platelets

reprecipitated within the O platelets. This reprecipitation C. Summarizing comments on the aging of subtransusand supertransus materialoccurred as early as 9 hours. The inset of Figure 12(c) is a

TEM image of a sample in an advanced state of aging The following three observations provide insight into thefactors responsible for the similarities and differences of the(681 hours at 650 8C) and illustrates the B2 platelets within


aging behavior of subtransus- and supertransus-treatedmaterials.

The first observation is with regard to the Widmanstattenstart temperature. Widmanstatten precipitation of the Ophase occurred only at temperatures below 875 8C for thesubtransus samples solutionized in the O 1 B2 regime.Above 875 8C, Widmanstatten precipitation was not evident;rather, equiaxed B2 grains coarsened at the expense of Ograins, as dictated by equilibrium requirements. Thus, aminimum DT, where DT is defined as the temperature differ-ence between the solutionizing and aging temperatures, isnecessary for Widmanstatten formation. For supertransusfully B2 and subtransus a2 1 B2 microstructures, Widman-statten precipitation initiated at a much-higher temperature(T . 900 8C).

The second observation is that lower aging temperaturesinduced discontinuous precipitation. For the subtransusmicrostructures, discontinuous precipitation of O and bplatelets occurred at O/B2 and a2/B2 boundaries for T #

Fig. 15—SEM image of the as-rolled Ti-12Al-38Nb microstructure after750 8C. For supertransus-solutionized microstructures agedaging at 650 8C for 304 h, followed by water quenching.at 650 8C, coarsening of O cells occurred preferentially at

transformed B2 boundaries. In both cases, the interface ofthe coarsening grain-boundary precipitates migrated towardthe grain interior. Combining the first two observations, foraging steps between 750 8C and 875 8C, Widmanstattenprecipitation was favored, while aging below 750 8C favoreddiscontinuous precipitation. These are in agreement withsimilar observations in other metallic alloy systems, wheresmaller DT values favor Widmanstatten precipitation andlarger DT values favor discontinuous precipitation.[27]

The third observation concerns the phase stability of near-Ti2AlNb alloys. Because the discontinuous precipitation ofthe subtransus-solutionized and aged microstructures con-sisted of O 1 b platelets, the O 1 b phase field is expectedto encompass the 650 8C to 750 8C temperature range. Theboundary between the O 1 B2 and O 1 b fields is expectedto lie between 875 8C and 750 8C for near-Ti2AlNb alloys. Fig. 16—Microdiffraction patterns of the O and b structures for a heat-

treated (650 8C/435 h/WQ) Ti-12Al-38Nb sample illustrating the (a) [00-However, the ability to obtain O 1 b microstructures1]O zone axis and the (b) [011] b zone axis.depends upon the volume fraction and composition of the

preaged B2 phase, since the supertransus-solutionized thenaged microstructures maintained the ordered B2 structure

slightly above 800 8C, since only a small volume of Oeven after 681 hours of aging at 650 8C. The diffusionplatelets was observed for the 800 8C condition, while morekinetics of the a2 phase appeared to be even more sluggishO phase was observed at 750 8C and no O phase wasthan that for the B2 phase, since only negligible volumeobserved at 850 8C.fractions of the a2 phase had transformed after 450 hours

Similar to the studies of the near-Ti2AlNb alloys, agingat 650 8C (compare rows 6 through 9 of Table II).treatments were performed at 650 8C in order to investigatethe microstructural stability of Ti-12Al-38Nb. Judging from

B. The Ti-12Al-38Nb Alloy the relatively constant 30 vol pct of O phase, the microstruc-tures appeared to be stable after 50 hours. Therefore, lessThe measured composition of the Ti-12Al-38Nb sheet

material, provided in Table I, was close to Ti-13Al-39Nb. time was necessary to equilibrate Ti-12Al-38Nb microstruc-tures at 650 8C compared to the near-Ti2AlNb microstruc-The as-rolled sheet consisted of b grains which were slightly

elongated in the rolling direction. No O phase was detected tures. Figure 15 depicts the as-rolled microstructure afteraging for 304 hours. The Widmanstatten O phase was identi-by TEM analysis of foils taken from the as-rolled sheet. A

DTA using 3 8C/min heating and cooling rates also did not fied in the [00-l] O microdiffraction pattern of Figure 16(a).The disordered b structure was retained in samples agedindicate any phase transformations occurring between 600

8C and 1000 8C. However, aging treatments between 650 between 650 8C and 760 8C, as evidenced from the [011]b microdiffraction pattern (Figure 16(b)). The OR between8C and 800 8C confirmed the precipitation of fine O platelets

within the b grains. Thus, the cooling rates employed in the O and b was identical to the O/bcc OR of the near-Ti2A1Nballoys. Similar to that for the near-Ti2AlNb alloys, the agingDTA analysis were rapid enough to suppress a significant

volume of O-phase precipitates, as verified by the lack of transformation resulted in misfit dislocations which formedat O/b interfaces.O precipitates found in the DTA sample. Based on the disap-

pearing-phase method, the b transus was estimated to be The b grain size was observed to increase with an increase


in the solution-treatment temperatures. Table IV summariesthe grain size data obtained from the as-rolled and heat-treated Ti-12Al-38Nb specimens. Under isothermal heat-treatment conditions, normal grain growth for single-phasematerials is well described by the following empiricalequation:[28]

dn 2 dn0 5 kt [1]

where n is the grain growth exponent, d is the heat-treatedb grain size, d0 is the initial grain size (b grain size of theas-rolled sheet), and t is the annealing time. The temperaturedependence of the rate constant (k) follows the Arrheniusequation,

k 5 k0 exp (2Qapp /RT ) [2]

where k0 is a kinetic constant, Qapp is the apparent activationFig. 17—A plot of ln k vs (1/T), based on Eqs. [1] and [2], depicting theenergy for grain growth, R is the gas constant, and T is theb grain growth for heat-treated Ti-12Al-38Nb.

absolute temperature. For most single-phase materials, thevalue of n ranges between 2 and 10 due to the drag forceexerted by solute atoms on grain boundaries.[29] Assumingn 5 3 (as is typically observed for several metals[30]), thetemperature dependence of k was analyzed using the ln kvs (1/T ) plot depicted in Figure 17. These data suggest thatthere are two distinct temperature regimes, with differentvalues of the apparent activation energies. For grain growthbelow 950 8C, Qapp is approximately 258 kJ/mol. This valueis close to a value of Qapp 5 252 kJ/mol, which was predictedfrom the diffusion coefficients of the binary Ti-Nb system.[31]

Thus, it appears that the addition of 13 at. pct Al does notsignificantly affect the diffusion rates below 950 8C. Forgrain growth above 950 8C, the measured Qapp was approxi-mately 96 kJ/mol, which is much smaller than that expected.In particular, this is unexpected, because typically graingrowth is correlated with lattice self-diffusion. Anomalies Fig. 18—SEM image of a 900 8C solutionized Ti-12Al-38Nb sample, which

was aged at 650 8C for 5 h. The O phase precipitated initially near the bhave been observed for b grain growth, which result in lowergrain boundaries.values of Qapp close to the transus temperatures;[31] however,

a transus is not expected near 1200 8C and this observationremains unexplained.

IV. DISCUSSIONAfter the grain sizes were determined for the respectivesupertransus solutionizing temperatures, samples were again

A. Phase Equilibria for Near-Ti2AlNb Alloysaged at 650 8C. Similar to that for the as-rolled material, astable volume fraction of approximately 0.30 O platelets The compositions of the a2, B2, and O phases, which

are illustrated in Figure 1, for the near-Ti2AlNb alloys areprecipitated within 50 hours. In addition, the O phase firstprecipitated at the large b grain boundaries. This is depicted consistent with recent compositional data for supertransus-

processed O alloys that were heat treated for 1000 hoursin Figure 18 for a sample which was solutionized at 9008C, followed by water quenching then aging at 650 8C for and longer.[17] Thus, the times of exposure employed here

appear to be sufficient for reaching equilibrium composition5 hours.

Table IV. Ti-12Al-38Nb Heat Treatments and Measured Average Grain Sizes and Phase Volume Percents and Compositions

b O dHeat Treatment Al Nb Vp Al Nb Vp (mm)

650 8C/55 h/WQ na na 72 pct na na 28 pct 33.2900 8C/5 h/WQ 12.4 40.7 100 pct — — 0 pct 138.2900 8C/5 h/WQ/650 8C/60 h/WQ na na 71 pct na na 29 pct 138.2950 8C/5 h/WQ na na 100 pct — — 0 pct 197.5950 8C/5 h/WQ/650 8C/62 h/WQ na na 69 pct na na 31 pct 197.51200 8C/5 h/WQ na na 100 pct — — 0 pct 336.91200 8C/5 h/WQ/650 8C/60 h/WQ na na 68 pct na na 32 pct 336.9

WQ: water quenched; and na: not available.


that fully-O microstructures are stable. However, no experi-mental verification has been provided. Experimental verifi-cation of this regime is expected to require extremely longaging times, since the transformation kinetics at lower tem-peratures are slower. Higher nominal Al contents signifi-cantly increase the temperature of the single-phase Ofield.[15] This was verified by the fully-O microstructure,which was produced through an 875 8C heat treatment ofTi-25Al-24Nb. However, the processing and heat-treatmentschedules have a significant influence on the phases that arepresent. For example, O 1 B2 solution treatments followedby long-term aging treatments below 875 8C resulted in O 1bcc microstructures, rather than fully O, for Ti-25Al-24Nb.

Interstitial oxygen content also has a significant influenceon the phase boundaries, especially the B2 transus and a2

1 B2 field. The B2 transus was 21 8C to 34 8C lower forTi-25Al-24Nb than for Ti-25Al-23Nb, although the Ti, Al,and Nb contents were similar (Table I). Greater oxygencontent tends to stabilize the a2 phase,[4,33] and, in this case,the higher interstitial oxygen level is believed to haveincreased the temperature range for the a2 1 B2 field and,

Fig. 19—The Ti-22Al isopleth with the current data overlaid for a Nb subsequently, caused the B2 transus to rise for Ti-25Al-composition of 27 at. pct.

23Nb (930 ppm oxygen) with respect to that for Ti-25Al-24Nb (280 ppm oxygen). Dividing the transus differencefor these two alloys by the difference in the oxygen content,one obtains an approximately 10 8C increase in transus tem-levels, except perhaps at 650 8C, where discontinuous precip-perature for every 220 ppm level of oxygen. This magnitudeitation had not progressed to completion. The Nb contentis similar to that observed for Ti-22Al-23Nb[33] and Ti-25Al-of the a2 phase (Ti-25Al-16Nb) was significantly higher10Nb-3V-1Mo.[34] This is also similar to the work of Vanthan that found in a Ti-25Al-17Nb alloy (a2 ' Ti-25Al-Thyne and Kessler,[35] who found that the addition of 10Al11Nb[4]), and this is portrayed in the 1025 8C and 1000 8Cto pure Ti increased the oxygen sensitivity, such that theternary isotherms of Figures 1(a) and (b), respectively. Atransus temperature increased approximately 5 8C/100 ppmwider compositional range for the Nb/Ti ratio in a2 is possi-of oxygen. Clearly, interstitial oxygen levels must be consid-ble, because Nb and Ti atoms occupy the same site in theered whenever phase equilibria studies are conducted fora2 lattice. On the other hand, the Al concentration is fixedtitanium alloys.because of the specific lattice site occupancy. As shown in

As observed in the aging study, the lower-temperaturethe 950 8C and 875 8C isotherms (Figures 1(c) and (d)), themicrostructures do not always consist of only the O and bccbcc phase is depleted in Al at lower temperatures within thephases. Once the a2 phase is formed, through solutionizingO 1 bcc phase regime. This leads to composition-inducedwithin the a2 1 B2 or a2 1 B2 1 O fields or supertransusdisordering.solutionizing followed by controlled cooling, it is quite stableIn order to better visualize the phase regions of the near-and does not fully transform after long aging times at lowerTi2AlNb alloys, the Ti-22Al isopleth, first developed bytemperatures. This is an important point to consider for theMiracle et al.[32] and later modified by Boehlert et al.,[19]

phase evolution of near-Ti2AlNb alloys, as the a2 phase ishas been provided in Figure 19. The current data for the Ti-not an equilibrium phase below 975 8C. The diffusion kinet-23Al-27Nb alloy are overlaid on this figure. The B2 transusics for a2 are very sluggish at lower temperatures, sinceis approximately 1070 8C. The a2 1 B2 phase field rangesnegligible decreases in a2 phase volume percents werebetween 1000 8C and 1070 8C. The narrow three-phase a2 1observed (Table II). The rim-O phase, which forms at theB2 1 O field ranges between 975 8C and 1000 8C. Althoughperiphery of a2 (Figures 2(e) and 4(b)), is considered to bethis phase diagram and others[15,25] have included the three-partially responsible for this, as it appears to act as a diffusionphase field for nominal Nb contents up to 35 at. pct, tobarrier between the relatively low Nb–containing a2 phasethe authors’ knowledge, this is the first work to verify theand the high Nb–containing bcc phases. On the other hand,existence of this regime for alloys containing more thanthe a2 phase is easy to remove by supertransus solutionizing,20Nb. The O 1 B2 phase field ranges between 750 8C andwhereby single-phase B2 microstructures are quickly equili-975 8C. Below 750 8C, the B2 phase undergoes composi-brated. In such cases, the O phase readily transforms to B2tional-induced disordering leading to the occurrence of theon going into the supertransus field, and, therefore, it noO 1 b phase field. The B2-to-b disordering is assumed tolonger interferes with a2 transformation.be a second-order transformation. Although the O, B2, and

Although the transus between the O 1 B2 and O 1 bb phases were all observed in some regions of the micro-fields is indicated in Figure 19, few published works arestructure at 650 8C (Figure 11), this is expected to be dueavailable that document this transition for near-Ti2AlNbto the slow diffusion kinetics associated with lower-tempera-alloys.[17,18] In addition, the practical significance of theture transformation and not to an equilibrium phenomenon.

For temperatures below 650 8C, the phase equilibria suggests order/disorder temperature is not well understood. As has


been conclusively shown here, the ability of obtaining equi- larger-grained (d 5 6.6 mm), 950 8C microstructure exhibitedWidmanstatten O at 850 8C (Figure 7(b)). This observationlibrium low-to-intermediate–temperature microstructures

depends largely on the processing and heat-treatment sched- is consistent with the fact that, for the larger-grained super-transus fully-B2 (d 5 177 mm) and subtransus a2 1 B2ules. In agreement with the current observations, it has been

shown that the a2 phase, formed by subtransus heat treat- (d 5 8.7 mm) microstructures, Widmanstatten precipitationinitiated at a higher temperature (T . 900 8C). Further workment, is extremely stable and resistant to transformation at

lower temperatures.[12,16] On the other hand, the decomposi- to understand this behavior is intended.tion of the B2 phase, although slow, is relatively rapid com-

2. Cellular precipitationpared to that of a2. During low-temperature aging, the B2Discontinuous precipitation is the second route in whichphase volume decreased, and its composition tended toward

the B2 phase transformed to O 1 bcc. The discontinuoushigher Nb and lower Al contents. In other words, the compo-precipitation event, observed in specimens solutionized atsition tended more toward the disordered b structure.subtransus temperatures then aged at 750 8C and below, isAlthough the compositions of the very fine bcc phases weredefined as cellular precipitation according to the reactionnot measured, the transition from B2 to b is expected to occur

between Ti-17Al-34Nb and Ti-12Al-41Nb, as the former B2 → b 1 O [3]composition was measured for the equiaxed B2 phase at

where B2 is the supersaturated ordered bcc phase, b is of900 8C (row 3 of Table II), and the latter was measured ina similar structure but is disordered with a lower thermody-the fully-b Ti-12Al-38Nb microstructure (Table IV). Thisnamic excess of solute (Al), and O is the equilibrium precipi-estimated transition composition is within the range of com-tate. Similar transformations have been observed in severalpositions measured for the B2 phase (Ti-20Al-36Nb) andmetallic alloy systems,[26,27,36] including the Ti-Cu, Ti-Ag,the b phase (Ti-12Al-33Nb) for Ti-20Al-37Nb and Ti-25Al-and Ti-Al systems, as well as in the current Ti2AlNb sys-18Nb alloys, respectively.[17] The highest temperature intem.[7,37,38] For example, cellular reaction of a 1 b at priorwhich the b phase was observed was 750 8C. This tempera-b grain boundaries consumes the tempered martensite mix-ture is within the order-disorder temperature range, basedture of a 1 a” in Ti-Cu. Cellular reaction also occurs in aon the work of Bendersky et al.[18]

1 b Ti-Ag systems, where a 1 Ti3Ag cellular product formsat massive a boundaries in Ti-17 wt pct Ag. Thus, it appears

B. The B2-to–O 1 Bcc Transformation to be easier for such systems to lower their total free energyby cellular transformation rather than by the nucleation and1. Widmanstatten precipitation

The B2 phase transformed to O 1 bcc mixtures at lower coarsening of the fine precipitates distributed uniformlywithin the bulk. The two factors which had the greatesttemperatures through three main routes. The first route was

through Widmanstatten O precipitation within the B2 grains. effect on the cellular precipitation were the solutionizingand aging temperatures. High solutionizing temperatures andStarting from a temperature high up in the equiaxed O 1

B2 field (maximum is approximately 975 8C), the transfor- low aging temperatures were more conducive to this typeof reaction. This may be explained in terms of the chemistrymation proceeds as follows. As long as the temperature

remains above 875 8C, a reduced B2 content with lower of the bcc phase and composition-induced disordering.Upon aging below 875 8C, Widmanstatten O phase precip-temperature is manifested through a reduction in the B2

grain size and a corresponding increase in the O grain size. itates uniformly within B2 grains, and the chemistry of theO phase typically adheres to its Ti2AlNb stoichiometry. ThisIn addition, the Nb content increases, while the Al content

decreases, in the B2 phase during this process; compare the type of precipitation has been discussed earlier and requiressome degree of supersaturation of the B2 lattice as well asB2 composition in rows 4, 3, and 1 in Table II. Below 875

8C, volume fraction changes of the B2 and O phases are no diffusive mass transport, which occurs primarily at the tipsof the precipitates where ledges are present and ahead oflonger grain-growth controlled. Rather, Widmanstatten O

precipitation takes over inside the B2 grains, (Figure 7). which large concentration gradients exist. When the tempera-ture drops sufficiently, as in the case of 650 8C aging follow-The transformation requires lattice reconstructions involving

dislocations and ledges, as in martensitic transformation ing quenching from a high solutionizing temperature, thesupersaturation becomes very high. This is because the pre-from a bcc to hcp phase of the same composition. Composi-

tional requirements require thermally activated diffusion of cipitation must also involve a transformation from the B2to the disordered b phase. At this point, one possible mecha-atoms across the O and B2 phase boundaries, where the

diffusion rate and, hence, the growth rate of the O phase is nism of microstructural change is a martensitic-type reaction,as was observed during the initial stages of transformationassisted by high-concentration gradients and large lattice

mismatch at the tips of the lens-shaped platelets. Thus, one of the supertransus material. However, because of theextremely low Al content of the b phase (Al less than 12can conceptualize two types of competing processes, whose

rates are a function of temperature. Above 875 8C, the driving at. pct) compared to the original B2 phase (Al typically 17to 22 at. pct), significant mass transport is necessary, whichforce for Widmanstatten formation is low, whereas grain-

boundary diffusion and grain growth kinetics are high makes martensitic reaction either unlikely, or, if it were tooccur, the product would be highly unstable. The secondenough to allow volume fraction adjustments through grain

size changes. Below that temperature, Widmanstatten pre- possibility is a continuation of Widmanstatten precipitation.However, the observations suggest that the diffusivity in thecipitation can occur much faster than grain size changes.

Initial results suggest the Widmanstatten start temperature bulk is likely too low at 650 8C to allow the large masstransport needed to separate the B2 into b and O phases ofis increased with increasing B2 grain size: the Ti-23A1-

27Nb fine-grained (d 5 3.8 mm), 875 8C solutionized micro- equilibria chemical composition. The third possible precipi-tation mechanism is a variation of this second possibility,structure remained equiaxed when aged at 850 8C, while the


in that the phase morphology is still of the Widmanstatten 15 minutes), these studies have reported the complete trans-formation of B2 to O. Because TEM did not reveal any B2type, but that a higher diffusivity associated with grain

boundaries is utilized. It is this mechanism that is believed phase after 2 hours at 650 8C and XRD data revealed notice-able shifts from the original B2 peaks after aging for 15to be responsible for the observed cellular precipitation.

The degree of supersaturation in the parent B2 phase minutes and 2 hours of the current work, it is likely that theB2 had completely transformed to O during the initial stagesdepends on the composition difference between the equilib-

rium bcc structures at the solutionizing and aging tempera- of aging. Thus, composition-invariant transformation occursat 650 8C for near-Ti2AlNb alloys. Within 9 hours, however,tures. As the composition difference between the B2 phase

and the b phase decreases, such as through lower-tempera- fine B2 platelets had reprecipitated throughout the micro-structure. In agreement with the work of Muraleedharan etture solutionizing and higher-temperature aging or through

slow-cooling treatments, the degree of supersaturation less- al.,[13] colonies of O grains coarsened initially at the trans-formed B2 boundaries and then within the bulk. Thus, similarens. In this case, only small chemistry changes are necessary

for the transition of B2 to b, which may be accommodated to the subtransus-solutionized then aged microstructures,lower aging temperatures promote discontinuous coarseningby diffusion within the bulk of the B2 region, i.e., by Wid-

manstatten precipitation. However, at higher solutionizing which initiates at prior B2 grain boundaries. Unlike thesubtransus solution-treated then aged microstructures, theand lower aging temperatures, the degree of supersaturation

of B2 with respect to b is increased. Note that large composi- precipitating phases of the supertransus solution-treated thenaged microstructures were O and B2 (not O and b), and thetion differences have been encountered between the B2 and

b phases for Ti-22Al-27Nb samples heat treated between fully-B2 microstructure completely transformed to O priorto reprecipitation of B2.650 8C and 1038 8C.[17] The B2 composition at 1038 8C for

that alloy was Ti-21Al-27Nb, while the b composition at 6508C was Ti-10Al-46Nb. Such large composition differences C. Ti-12Al-38Nb and the Pseudobinary Diagrambetween the parent and final phases provide high driving

The phase evolution in Ti-12Al-38Nb differed signifi-force for the transformation in Eq. [3]. Thus, the thermalcantly from that of the near-Ti2AlNb alloys. As a result,activation processes which favor cellular precipitation aresignificantly different microstructures were obtained fromenhanced by higher solution-treatment temperatures andsimilar processing and heat-treatment schedules. The disor-lower aging temperatures, which lead to larger bcc phasedered b phase was more stable for Ti-12Al-38Nb than forcompositional gradients. Although the phase instability atthe near-Ti2AlNb alloys. This is expected to be a result650 8C is similar to that observed by Rowe and Larsen,[8]

of the compositional dependence of the bcc structure. Thethe kinetics of the cellular precipitation in the current workordered B2 structure was not detected in the as-rolled (932were found to be unaffected by superimposed external8C) and 650 8C to 760 8C–aged conditions. Thus, the temper-stress.[9] It was not limited to segregated bands, nor was itature range of stable B2-containing microstructures is sig-associated with the fully-O microstructure.nificantly reduced for alloys containing lower Al and higherIn agreement with some past studies,[7,17] small volumeNb. This is depicted in the pseudobinary diagram for Ti 5fractions of cellular precipitation were also observed for50 at. pct, originally developed by Bendersky et al.,[18] whichsupertransus-solutionized samples that were subsequentlyhas been modified according to the current data (Figureaged within the O 1 B2 region. These observations are20). Note that, for the composition range between TiAl andconsistent with the rationale for the cellular precipitationTi2AlNb, three-phase equilibrium involving the a2, g, anddescribed previously, since large compositional differencesB82 phase was not included. For Ti-12Al-38Nb, the b transusbetween the B2 phases of supertransus-solutionized and O 1was close to 800 8C, as determined from the disappearing-B2–solutionized microstructures (Tables II and III) favorphase method. Thus, the Ti-12Al-38Nb bcc transus tempera-cellular decomposition. Cellular precipitation was notture is significantly lower than that of near-Ti2AlNb alloys.observed for supertransus-solutionized then slow-cooledBased on these observations and those previously mentionedsamples. Furthermore, subsequent aging of these samplesfor the near-Ti2AlNb microstructures, the B2 phase is consid-did not promote cellular reaction. Only a slight coarseningered to be stable only when it contains more than 12Al. No a2of the existing O 1 B2 laths occurred. This is also consistentphase was observed for Ti-12Al-38Nb, and the temperaturewith the previous reasoning, since the B2 phase graduallyranges for the a2 1 B2 and a2 1 B2 1 O phase fieldstransformed to O upon cooling and, therefore, decreased itsbecame narrow with increasing Nb and decreasing Al con-Al and increased its Nb concentrations along the way. Thus,tents (Figure 20). The diagram is approximated as a quasibi-the range of temperatures in which the lower-temperaturenary for simplicity. This is reasonable, because experimentalaging transformation mechanisms dominate is dependent onresults indicated that, for temperatures less than 800 8C, b-both the solutionized microstructure and the cooling rate.phase volume fractions were consistent with those computedfrom tie-line approximations. For example, the maximum3. Composition-invariant transformation

In specimens solution treated at supertransus temperatures O-phase volume predicted from Figure 20 for Ti-12Al-38Nbat 650 8C is approximately 30 pct, while the correspondingthen aged at 650 8C, martensitic-type O grains precipitated,

then coarsened at transformed B2 boundaries. The aging value for Ti-25Al-25Nb is nearly 100 pct. Both these valuesare close to those observed (Tables III and IV).transformation behavior of these samples was similar to the

composition-invariant B2-to-O transformation characterized No equiaxed O phase was present in the Ti-12Al-38Nballoy. This is believed to be a result of the processing temper-for Ti-24Al-16Nb aged between 800 8C and 940 8C[13] and

Ti-24Al-25Nb aged at 700 8C,[18] where, after long aging ature chosen for this alloy. Equiaxed primary O grains wouldbe expected if the processing temperatures were below thetimes, the B2 phase reprecipitated in platelet form within

the transformed B2 matrix. For very short aging times (t # transus (Tb ' 800 8C), as was the case for the near-Ti2AlNb


B2 1 O, and O 1 b phase regions were estimated forthe near-Ti2AlNb alloys. The observation of the three-phasea2 1 B2 1 O microstructure verified that a narrow three-phase field exists for alloys containing more than 20 at.pct Nb. Orientation relations existed between subtransusheat–treated, equiaxed a2/B2 and O/B2 grains, and theirexistence is explained in terms of the as-cast lath-type micro-structure. It is also to be noted that fine-grained, equiaxedfully-O microstructures are achievable through subtransusheat treatment of near-Ti2AlNb alloys, provided that thealloy contains at least 25 at. pct Al.

For solution treatments above 875 8C within the O 1 B2phase field, the volume fraction of the B2 and O phasesadjusted to equilibrium requirements through changes intheir grain size; thus, solutionizing at lower temperaturesled to smaller B2 grains and larger O grains. Associatedwith this change was a change in composition of the B2phase toward leaner Al and enriched Nb. In addition, thetransformation was reversible. Three distinct transformationmodes were observed within the B2 phase for the near-Fig. 20—The pseudobinary section of the Ti-Al-Nb system at 50 at. pctTi2AlNb alloys. Supertransus- and a2 1 B2–solutionizedTi taken from Bendersky et al.[18] with data from the Ti-25Al-24Nb, Ti-

23Al-27Nb, and Ti-12Al-38Nb alloys superimposed. “Or” represents the samples, which were aged within the O 1 B2 regime, exhib-orthorhombic phase. ited Widmanstatten O precipitation within the B2 grains.

The O 1 B2 solutionized samples exhibited WidmanstattenO precipitates within equiaxed B2 grains upon aging below

alloys. Phase stability is less complicated in Ti-12Al-38Nb, 875 8C. Below 750 8C, coincident with the temperatureas the O phase equilibrates more rapidly than the near- corresponding to the order-disorder transformation of theTi2AlNb alloys, and Widmanstatten precipitation is the dom- bcc phase as determined by Bendersky et al.,[18] Widmanstat-inant transformation mode. No cellular precipitation was ten O precipitation was replaced by a cellular precipitationobserved in Ti-12Al-38Nb. However, precipitation initiated reaction at the grain boundaries. The reaction products wereat b/b grain boundaries for specimens aged for short dura- O and the disordered b phase. The driving force for thistions (Figure 18). An important observation, based on the cellular precipitation is believed to be the large concentrationaging study of Ti-12Al-38Nb, is that fully-b microstructures difference between the B2 and b phases, with grain-bound-would be stable at T , 800 8C only for alloy compositions ary diffusion being important because of the relatively lowcontaining less than 12Al and more than 38Nb (Figure 20). temperatures involved. Higher solution-treatment tempera-Interpreting the pseudobinary diagram, Ti-10Al-40Nb would tures and lower aging temperatures were more conduciveborder the O 1 b/b regimes at 650 8C. This is in excellent to cellular precipitation. Supertransus solutionizing followedagreement with the b-phase composition measured for a Ti- by 650 8C aging resulted in a composition-invariant mecha-22Al-27Nb alloy at 650 8C.[17] nism where fully-B2 microstructures transformed to O

within 2 hours. For longer aging times, fine O plateletscoarsened initially at prior B2 grain boundaries, and thenV. SUMMARY AND CONCLUSIONSthey coarsened within the prior B2 grains. The B2 phase

The alloys examined had nominal compositions of Ti- then reprecipitated within the O grains.25Al-24Nb, Ti-25Al-23Nb, Ti-23Al-27Nb, and Ti-12Al- Low-temperature equilibrium was more readily obtained38Nb (at. pct). The former three were grouped as near- for the b-dominated Ti-12Al-38Nb alloy. In this case, Wid-Ti2AlNb alloys. Each alloy was processed from ingot to manstatten O precipitation occurred for all aging tempera-sheet form through hot-forging and hot-rolling procedures, tures between 650 8C and 800 8C, and no a2 phase waseach performed between 932 8C and 1000 8C. The near- present. Above 800 8C, fully-b microstructures equilibratedTi2AlNb alloys were processed at subtransus temperatures, and the size of the b grains was controlled by the solutioniz-while Ti-12Al-38Nb was processed within the supertransus ing temperature and hold time. No ordered B2 phase isb region. present in Ti-12Al-38Nb, indicating that greater than 12 at.

The study on the near-Ti2AlNb alloys shows that it is pct Al is necessary to induce the ordered B2 structure. Thepossible to obtain a wide array of equiaxed microstructures bcc order/disorder transition composition is between Ti-with a large range of grain sizes (4 to 177 mm), through 17Al-34Nb and Ti-12Al-41Nb. Using the heat-treatmentsubtransus processing followed by postprocessing heat treat- data for each of the O 1 bcc alloys containing Ti 5 50 at.ments. Conversely, neither fine-grained microstructures nor pct, the pseudobinary diagram, depicting the phase equilibriasuch a variety is possible when the material is supertransus of such alloys, was modified.processed, as has often been employed in the past. Thus,subtransus processing offers a much wider window for opti-mizing mechanical properties. In particular, the effects of ACKNOWLEDGMENTSthese microstructures on creep behavior is provided inPart II.[39] This research was performed at the Wright-Patterson Air

Force Research Laboratory Materials and ManufacturingThe temperature ranges for the a2 1 B2, a2 1 B2 1 O,


15. K. Muraleedharan, T.K. Nandy, and D. Banerjee: Intermetallics, 1995,Directorate under Air Force Contract Nos. F33615-C-96-vol. 3, pp. 187-99.5258 and F33615-C-96-5251 to UES, Inc. We thank Dr. R.

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18. L.A. Bendersky, W.J. Boettinger, and A. Roytburd: Acta Metall. Mater,sions. One of the authors (CB) acknowledges the support 1991, vol. 39, pp. 1059-69.received from Johns Hopkins University during the writing 19. C.J. Boehlert, B.S. Majumdar, V. Seetharaman, D.B. Miracle, and R.

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