Phase relationships and transformations in the ternary aluminum—titanium—tantalum system

Acta metall, mater. Vol. 43, No. 7, pp. 2625 2640, 1995

Pergamon

0956-7151(94)00468-4

Elsevier ScienceLtd Copyright :O 1995 Acta Metallurgica Inc. Printed in Great Britain. All rights reserved 0956-7151/95 S9.50 + 0.00

PHASE RELATIONSHIPS A N D T R A N S F O R M A T I O N S IN THE T E R N A R Y A L U M I N U M - T I T A N I U M - T A N T A L U M SYSTEM M. L. WEAVER and M. J. KAUFMAN Department of Materials Science and Engineering, University of Florida, Gainesville, FL 326l 1, U.S.A. (Received I6 February 1994; in rezised.~)rrn 8 November 1994)

Abstract--The solid-state phase transformations and phase equilibria in a series of AI-Ti-Ta alloys in the range 38 55%A1, 19 30%Ti and 20-38%Ta have been investigated. It is shown that a variety of phases exist over this composition range and that the microstructures vary dramatically with prior thermal history. For example, the disordered fl phase exists at high temperatures over a large composition range and its decomposition depends strongly on both composition and cooling rate in a manner similar to that reported in titanium rich Ti AI Nb alloys. In addition, the ~ phase is shown to extend to higher AI contents than originally reported and its decomposition is similar to that observed in binary Ti-A1 alloys (i.e. it forms "massive" y at high cooling rates and ~ + 7 or a2 + "/at slower rates). The ~-(Ta~Ti~ +)2A1 phase becomes stable at lower temperatures in the tantalum-rich alloys studied and its formation is also shown to be a strong function of cooling rate and composition. The various results are used to suggest revisions to the ternary AI Ti-Ta isothermal sections at both 1723 and 1623 K.

I. INTRODUCTION Currently, gamma titanium aluminide alloys containing additions of refractory elements such as V, Cr, Mn, Mo, Nb, or Ta are receiving attention as potential high temperature structural materials. Although many ternary and quaternary alloys have been investigated, their development into viable engineering materials has been hindered by a lack of detailed information concerning the phase transformations and phase equilibria in these systems and in the related binaries. Undoubtedly, a detailed understanding of these issues is necessary before alloy development can proceed effectively. A review of the available literature reveals similar isothermal sections at 1723 and 1373K for the AI T i - T a system [1,2]. Although these studies concur on the existence of the c~-Ti, ~2-Ti3AI, fl-Ti, 7-TiAI, r/-A13(Ti,Ta ), and a-Ta2Al phases, recent revisions of the binary A I T i phase diagram and discrepancies concerning the crystal structures and stoichiometries of the phases present in the A I - T a system have rendered these findings questionable [3 6]. Das et al. [7-10] and McCullough et al. [11, 12] recently re-evaluated the ternary phase equilibria at 1373 K. Both agreed on the presence of the ~, ce2, fi, 7, r/, and a phases; however, they indicated significant differences in the locations of the various phase boundaries. They also reported the extension of the 6-(AIzTa) and E-(AITa) phase fields into the ternary system and indicated the occurrence of a series of complex solid-state phase transformations during coiling. Similar observations were reported by Jewett a~

4~

I

and coworkers [8, 13] and Weaver et al. [14] at even higher temperatures (1713 and 1723 K respectively). Discussions of the solid-state phase transformations in this system are limited. Kim et al. [15] showed that alloys containing 3 2 - 4 0 % A l and approx. 2 5 % T a ordered to the B2 (CsCI) structure during splat quenching. Re-annealing of the splat quenched specimens below 1473 K resulted in the decomposition of this B2 phase into a fine two-phase mixture of ?, and a. More recently, McCullough et al. [11, 12] using combinations of in situ high temperature X-ray diffraction ( X R D ) and transmission electron microscopy (TEM), found that small additions of Ta (~< 10%) had little effect on the solidification behavior of TiA1 alloys containing 4 5 M 8 % A I and that these alloys solidified much like binary Ti-AI alloys containing 48-56%A1 with a forming as the primary phase and 7 forming as an interdendritic segregate. Alloys containing ~ 5 0 A I and > 10%Ta solidified with fl as the primary phase. This fi was surrounded by peritectic layers of c~ and 7 segregate and was significantly cored with dendrite cores reportedly containing as much as 15% more Ta than the dendrite peripheries; this coring resulted in the formation of different microstructural constituents during solid-state cooling. For example, the dendrite cores containing ~ 2 5 - 3 0 % T a reportedly transformed to ~ and then to polycrystalline ?, upon cooling. The dendrite cores containing 3 5 - 4 0 % T a in the higher Ta alloys precipitated a and a upon cooling, eventually transforming completely to c~and ¢. Further cooling caused the a to transform to either 100%?,, ? + o , or ~2+?' plates depending on its

2625

2626

WEAVER

and KAUFMAN:

TERNARY

composition. Meanwhile, the a reportedly transformed to a + ?, via a cellular reaction sequence. Boettinger e t al. [16] used differential thermal analysis (DTA), high temperature XRD, scanning electron microscopy (SEM), and TEM to evaluate the phase equilibria and microstructural stability of an AI-25Ti 25Ta alloy over a wide temperature range. They found that this alloy consists of 100% /~ above 1798 K and :c + fl between 1723 and 1773 K. Between 1573 and 1673 K, this alloy was observed to consist of various combinations of~, fl, ?, and a, while, below 1573 K, the alloy consisted of 7 and a only. This data was used to estimate a partial isothermal section at 1603 K. Although both the studies of Boettinger e t al. [16] and McCullough e t al. [11, 12] accurately indicated the phases present at elevated temperatures in the AI- Ti Ta system, no detailed accounting of the phase equilibria at high temperatures and the solid-state transformation sequences during cooling is available. The purpose of this paper is to better define (1) the solid-state phase transformations that occur during cooling, particularly near the composition AI 25Ti-25Ta and (2) the phase equilibria in the AI Ti Ta system at 1623 and 1723 K. In order to accomplish these goals, 28 ternary alloys were selected, isothermally annealed, and either water quenched (WQ), air cooled (AC) or furnace cooled (FC). Selected samples were then characterized using combinations of light optical microscopy (LOM), SEM, TEM, XRD, electron probe microanalysis (EPMA), and DTA in order to provide information concerning the phase transformations during solidstate cooling.

Al Ti Ta SYSTEM 2.

E X P E R I M E N T A L

27 Ternary alloys (Table 1) were prepared from pure elemental constituents (99.999%A1, 99.985%Ti, 99.985%Ta) by arc melting on a water-cooled copper hearth under an atmosphere of purified argon in the form of 0.050 to 0.090 kg buttons. Each button was inverted and remelted a minimum of six times in order to insure homogeneity. Some specimens were then drop cast into copper chill molds containing rectangular (15 x 15 x 40 mm) or cylindrical (15 mm diameter x 40mm) cavities. As-cast sections were then taken from each ingot and examined using LOM and EPMA in order to confirm the chemical compositions of each alloy. In many cases, oxygen and nitrogen contents were also measured with a LECO oxygen/nitrogen gas analysis system. Oxygen and nitrogen contents prior to heat treatment ranged from 100 to 331ppm oxygen and 100 to 380ppm nitrogen by weight. Sections from the as-cast samples were cut, cleaned in ethanol, wrapped in Ta foil and placed in Ta-foillined zirconia boats. Heat treatments were conducted at 1823, 1723 and 1623 K in a continuous flow of purified argon to reduce oxidation and interstitial contamination. Heat treatment times ranged from 1 to 72 h. After heat treatment, the samples were initially WQ and analyzed to determine the equilibrium phases present at temperature. Five of the 28 alloys were annealed and AC ( ~ 450 K/min) and/or FC ( ~ 10 K/min) in order to determine how the microstructures evolved as a function of cooling rate and composition. The locations of these five alloys are shown on the "'accepted" 1373 K phase diagram

Table 1. Compositions of AI Ti T a alloys investigated in this study Nominal (at.%)

Measured (at.%)

Phases present at:

1D

AI

Ti

Ta

A1

Ti

Ta

I 2 3 4 6 7 8 16 18 25 26 30 32 33 37 38 39 40 45 46 47 49 50 51 52 53 54

65.00 55.00 50.00 45.00 35.00 30.00 25.00 45.00 75.00 55.00 36.00 65.00 40.00 50.00 32.00 25.00 72.00 65.00 55.00 45.00 72.0(I 50.00 61.66 48.70 44.00 54.00 60.00

15.00 20.00 25.00 30.00 40.00 45.00 50.00 50.00 12.50 40.00 28.00 10.00 45.00 5.00 55.00 25.00 18.00 15.00 20.00 20.00 23.00 30.00 6.78 23.44 5.00 3.00 15.00

20.00 25.00 25.00 25.00 25.00 25.00 25.00 5.00 12.50 5.00 36.00 25.00 15.00 45.00 13.00 50.00 10.00 20.00 25,00 35.00 5.00 20.00 31.56 27.86 51.00 43.00 25.00

63.99 56.60 49.34 44.61 35.34 31.56 25.96 43.73 75.07 53.73 34.62 66.03 41.49 42.85 33.35 25.05 70.72 63.08 57.84 42.54 71 9 5 47.29 60.06 49.19 43.79 53.44 62.66

14.51 19.03 25.16 3(I.35 40.30 44,17 49.59 50,67 12.31 40,00 22,31 10,72 43.76 4.36 54.07 22.90 17.25 16,87 18.82 19.68 22.09 30.39 7.49 23.82 5.29 3.54 12.38

21.50 24.37 25.50 25.05 24.36 24.27 24.45 5.60 12.62 6.27 43.07 23.26 14.75 52.78 12.58 52.05 12.03 20.05 23.35 37.78 5.96 22.32 32.45 26.99 50.92 43.02 24.96

1623 K

1723 K

?÷c~ ~+~r

¢e ~+/~

~+/~ m

/~+~

*/ ./

/~+~

q+L ~ + ' / + t/ o-÷7

-.

/~+o

-o:+o"

q-L ,x÷o-

q+a ~ +[~ ¢$ --

:~+3+q

W E A V E R and K A U F M A N : Ta

TERNARY

,;),. B ~

2

6

~

~

Ti ~

AI 80

60

40

Fig. I. Schematic representation of the 1373 K ternary isotherm o f D a s et al. [8] showing the composition of the five alloys used to study microstructural evolution in this work.

in Fig. 1. Where possible, alloys were solution treated in a single phase region to eliminate casting segregation prior to the treatments at the selected temperatures. The alloys were then sectioned and analyzed via LOM, room temperature XRD, EPMA, SEM, and TEM. Sections for LOM, SEM, and EPMA were cut using a low speed diamond saw, mounted in bakelite, and polished to 0.05 #m with colloidal silica. LOM and SEM specimens were examined in both the etched and unetched conditions. SEM was conducted on either a JEOL 35C, JEOL 840, JEOL 6100 or a Cambridge 100 operated at accelerating voltages of 15-25 kV. A modified Kroll's etchant (10 vol.% HF, 14vo1.% HNO~, and 76vo1.% H2) was used for etching prior to LOM analysis. Quantitative chemical information was acquired via EPMA on unetched specimens using a JEOL 733 Superprobe operated at 20kV equipped with a Tracor wavelength dispersive spectroscopy (WDS) Table 2. EPMA compositions of phases in multiphase alloys after heat treatment at 1623 K History time/cooling

l

41 h/WQ

2

43 h/WQ

3

48 h,'WQ

4

ST + 43 h/WQ

26

ST + 21 h;WQ

49

48 h/WQ

45

24 h/WQ

50

48 h/WQ

Phase

%AI

%Ti

%To

a r/ ~r ,' ~ o" c~

45.83 71.38 46.42 57.46 49.86 40.52 45.45 39.39 37.61 34.29 49.29 40.49 44.85 58.36 46.62 73.57

8.06 14.07 1 I. I 1 20.81 28.83 15.90 33.46 20.24 34.61 17.97 30.55 20.50 9.61 20.31 4.03 9.15

46.11 14.56 42.47 21.73 21.31 43.58 21.09 40.38 27.78 47.74 20.16 39.01 45.55 21.33 49.34 17.28

tr :t a o" ?, ~r r/

History time/cooling

WQ = water quenched. ST=solution treated at 1823K/5h followed by WQ prior to annealing at 1623 K.

Phase

%AI

%Ti

%To

e rt #

57.42 70.55 45.85 50.16 47.41 51.27 45.21 50.06 43.38 44.68 52.23 57.23 37.87 34.37 71.07 60.98 71.29 69.28 57.46 61.09 70.78 49.12 46.89 43.98 39.47 72.65 68.69 50.24 47.00 46.85 50.31 46.60 60.27 55.73 61.32 70.92

15.98 13.05 24.94 26.41 24.44 25.87 25.83 27.57 49.50 50.42 40.62 38.55 23.70 20.92 12.51 8.90 16.14 20.02 17.72 19.24 13.65 36.89 36.43 22.25 12.40 20.82 27.42 30.03 28.74 23.72 25.53 3.30 3.78 15.33 9.43 12.37

26.60 16.40 29.21 23.44 28.15 22.86 28.96 22.37 5.74 6.28 7.15 4.22 41.43 44.71 16.39 30.12 12.56 10.70 24.81 19.77 15.58 14.00 16.68 33.77 48.14 6.53 3.90 19.72 24.26 29.43 24.16 50.10 35.95 28.94 29.24 16.70

1

72 h/WQ

3

4 h/WQ

3

48 h/WQ

3

S T + 12 h/WQ

,8

16

4 h/WQ

25

48 h/WQ

26

72 h/WQ

30

72 h/WQ

39

4h/WQ

40

80 h/WQ

41

6h/WQ

46

108 h/WQ

47

4 h/WQ

49

ST + 4 h/WQ

:~ // ~ L // a ~/ 6 q L :~ 7 q c~ [/ fl a ~/ L ~

51

48/WQ

53

40 h/'WQ

54

12.5 h/WQ

20

Atomic % Ti

Alloy

2627

Table 3. EPMA compositions of phases in multiphase alloys after heat treatment at 1723 K Alloy

~k'~S 4 0 /

Al T i - T a S Y S T E M

~ 6 ~. o r/

WQ = water quenched. ST = solution treated at 1823 K/5 h followed by WQ prior to annealing,

system calibrated with elemental standards resulting in a relative accuracy of _+0.2%. Bulk alloy compositions were confirmed using inductively coupled plasma emission spectroscopy. The compositions of the various alloys used in this study are listed in Tables 1-3. XRD of bulk slices or powders ground from solid sections was performed on a Philips ADP 3720 instrument operated at 40 kV and 20 mA. Typical XRD scans were made from 20 = 5 ° to 100 using CuK~ radiation. TEM specimens were electrolytically polished in a Struers Tenupol-2 unit using a solution of 10% perchloric acid and 90% methanol at 238 K and a current of 40 mA. TEM analysis was carried out on a JEOL 200CX at an accelerating voltage of 200 keV or on a JEOL 100C at an accelerating voltage of 120 keV. DTA was conducted on a Netzsch STA 429/409 instrument in a purified He atmosphere. Small as-cast samples ( ~ 1 0 m g ) were placed in A1203 crucibles and heated next to a reference crucible containing platinum powder. Heating and cooling cycles were conducted at rates ranging from I0 to 30 K/min.

2628

WEAVER and KAUFMAN:

TERNARY AI Ti Ta SYSTEM

Fig. 2. Backscattered scanning electron micrographs of the as-cast microstructures observed in Alloys 2, 3, 4. 26 and 49. (a) Alloy 2; (b) Alloy 3; (c) Alloy 4: (d) Alloy 26; and (e) Alloy 49.

3. RESULTS 3.1. Microstructural Evolution 3.1.1. Solidification microstructures Figure 2 shows back scattered SEM micrographs taken from as-cast samples of Alloys 2, 3, 4, 26, and 49. It is clearly observed that the as-cast microstructures of all five alloys consist of coarse dendrites with varying degrees of segregation. In all cases, the dendrites appeared to exhibit orthogonal symmetries indicating that the primary solidification phase was [I (b.c.c.) [3, 11, 12, 15]. In addition, the primary den-

drites were often surrounded by two or more phases implying solidification via a series of cascading reactions (possibly the peritectic-like reactions: L +/~ --+ ~, L + :~ ---+7, and L + }, ---+r/). In addition, considerable Ta segregation towards the primary fl dendrite cores was observed [17] resulting in complex phase transformations during solid-state cooling. Current observations are consistent with those of McCullough et al. [11, 12] who have reported analogous solifidification mierostructures in alloys with similar compositions. Their high temperature X-ray diffraction experiments suggest solidification as

WEAVER and KAUFMAN: TERNARY A1 Ti Ta SYSTEM

2629

Fig. 3. TEM micrograph and respective diffraction pattern of the ordered B2 phase in Alloy 26 following an 1823 K heat treatment followed by WQ. (a) Darkfield TEM micrograph; and (b) [I 10]B2 zone axis.

Fig. 5. Backscattered electron micrograph of Alloy 2 showing the transformed fl phase and liquid regions after 1823 K/5 h/WQ.

follows:

structure [Fig. 4(a)]. In addition, the thermal antiphase boundaries (APBs) observed within the B2 indicate that this alloy was 100% fl at 1823 K and ordered to B2 during WQ. Analysis of the martensitic transformation product in Alloys 2, 3, 4, and 49 revealed that the product martensite plates had the ordered hexagonal D019 (~2) crystal structure and were dispersed in an ordered B2 matrix. A typical microstructure is shown in Fig 5 for Alloy 4. Elemental partitioning did accompany the transformation as evidenced by the contrast observed in the backscattered image. EDS analysis in the T E M revealed the compositions of the e; and fl phases to be approx. 46AI-32Ti-22Ta and 42A1 24Ti 34Ta respectively. An additional feature observed was the presence of midribs in the larger plates. Similar features have been observed by Weykamp et al. in Ti AI Nb alloys [18]. In all four alloys, the Burgers orientation relationship was observed between the ~; plates and the B2 matrix [19] with {1 10j~2B2//"

L - - , f l + L--~ fl + ~ + L--~ fl + ~ + 7 where the fl and :¢ phases decompose further during solid-state cooling to complex mixtures of or, 7, and sometimes ~2 with the exact transformation depending upon composition.

3.1.2. Water quenched samples 3.1.2.1. Water quenched from 1823 K. Optical and S E M observations. Water quenching from 1823 K results in microstructures consisting of large equiaxed grains with grain diameters approaching 2 mm. Typical microstructures are shown in Fig. 3 where it is apparent that Alloys 2, 3, 4, and 49 transformed martensitically during cooling whereas Alloy 26 showed no apparent indications of a transformation. In addition, evidence of interdendritic boundary liquation was observed in Alloy 2 [Fig. 3(a)]. T E M observations. TEM/selected area diffraction (SAD) analyses of Alloy 26 confirmed that this alloy was indeed single phase with an ordered B2 crystal

Fig. 4.Bright field TEM micrograph of Alloy 3 illustrating the ~ and B2 phases as observed in Alloys 3, 4 and 49 after heat treating at 1823 K followed by WQ.

(0001)~(1120)~. Thus, the stable phase at 1873 K in Alloys 3, 4, 26, and 49 is fl is fl while fl and liquid phases are stable in Alloy 2.

Fig. 6. Light optical micrograph of Alloy 2 (1723K/5h/WQ). The arrow indicates a prior :~ grain boundary.

2630

WEAVER and KAUFMAN: TERNARY A1 Ti Ta SYSTEM

Fig. 7. Micrographs of Alloys 3 and 49 after WQ from 1723K showing transformed a and :~ which has transformed to }'mduring cooling. (a) and (c) Light optical micrographs; (b) and (d) backscattered electron images of Alloys 3 and 49 respectively.

3.1.2.2. Water quenching j}'om 1723K. Optical and SEM observations. Equiaxed grains were observed in Alloy 2 after WQ from 1723 K indicating the presence of a single phase at 1723 K. Clearly from Fig. 6, it is observed that these grains transformed during quenching and that the transformation product appeared somewhat different from those observed after WQ from 1823 K. XRD and TEM analysis indicated that this structure consists almost exclusively of the ordered tetragonal 7 phase, the details of which, are described below. In Alloys 3 and 49, duplex microstructures were observed with the resulting phase distribution differing depending upon the thermal history of the alloy. For example, solution treating (ST) and WQ from the fl phase field (1823 K) prior to annealing at 1723 K resulted in the formation of a second phase at the prior fl grain boundaries and in the grain interiors [Fig. 7(a and b)]. On the other hand, direct heating (DH) of as-cast material resulted in the more equiaxed microstructures shown in Fig. 7(c and d). In the ST samples, room-temperature XRD indicated the presence of ~2, fl/B2, and 7 and a small volume fraction of a following WQ. In the DH samples, however, room-temperature XRD indicated only the presence of a, ~, and ~2.

TEM observations. As mentioned above, XRD and TEM analysis indicated that the transformation structure in Alloy 2 consists almost exclusively of the ordered tetragonal 7 phase (Fig. 8). Significantly, this 7 contains thermal APBs, local orientation variants, twins, and dislocations indicating transformation through an intermediate face-centered tetragonal (f.c.t.) phase. In addition, no composition differences, as determined using EPMA, could be determined in the resultant microstructures. Similar structures have been reported in binary Ti-AI alloys and appear to result from a massive transformation (~--*)'m) [5, 20]. This massive transformation occurs when the y nucleates a prior ~ grain boundaries and then grows in a coherent or semicoherent manner into the adjacent ~ grains yielding orientation variants similar to those observed when 7 plates precipitate in more slowly cooled Ti-AI alloys [21]. The twins are believed to result from the stresses induced by the volume differences between the c~ and 7 phases and the small but finite c/a ratio of 7 (c/a = 1.0175) [5]. Also, the APBs appear to be associated with the nucleation and growth of the 7m and are not representative of a disordered intermediate phase.

WEAVER and KAUFMAN:

TERNARY A1 Ti Ta SYSTEM

Fig.

Fig. 8. Alloy 2 (1723 K/5 h/WQ); (a) BFTEM micrograph showing local orientation variants, twins, and APBs; (b) SADP showing an [001]~ zone axis: (c) DFTEM showing a close-up image of the twins in Alloy 2.

In Alloys 3 and 49 after ST, T E M analysis (Fig. 9) revealed a martensitic structure consisting of B2 and e~ phases corresponding to the dark regions in Fig. 7(a and b). This structure formed upon cooling as described above for the 1823 K treatment. The light regions (mottled) in Fig. 7(a and b) consist of 7m as described above for Alloy 2 indicating that the stable phases at 1723 K are /3 and ~. Interestingly, the 7 variants and ordered domains observed in the 7m appeared somewhat coarser than those in Alloy 2 indicating that either (1) the transformation from e to 7 occurred at a higher temperature in Alloys 3 and 49; (2) the cooling rate was slower; or (3) the ~ composition was leaner in tantalum since tantalum is expected to slow down the kinetics of APB coarsening. X R D and T E M analyses of Alloys 2 and 3 after D H indicated the presence of ~z (in small quantities), 7, and a. Typical T E M microstructures are shown in Fig. 10. The 7 phase was present in two morphologies: (1) as a massive transformation product, 7,7,; and (2) as lamellae in coarse a + 7 colonies. These colonies were consistently surrounded by a thick layer of a which separated them from the neighboring 7m' The cr phase in the colonies exhibits no orientation relationship with the neighboring 7m although an O R with [001]o(110),//(011)~ { l 1T}~. was observed between the a and 7 within the colonies. E P M A

2631

9. Brightfield TEM micrograph of Alloy (ST + 1723 K/5 h/WQ) showing ~; +/7 martensite.

3

analysis after quenching showed the overall composition of the a + 7 colonies to match those of the transformed/3 regions observed in the ST + W Q alloy which suggests that either: (1)/3 is present at 1723 K and decomposes during W Q to a mixture of a and 7; or (2) ~r is present at 1723 K rather than/3. However, considering the sluggish nature of most reactions involving ~r, option (1) is unlikely. A review of the literature also exposes contradictions concerning the stable phases at 1723K. McCullough et al. [10] reported the presence of stable ~ and cr phases at 1723 K in an alloy of approximate composition 45AI-28Ti-27Ta while Boettinger et al. [16] reported c~ and/3 in an alloy of composition 50AI 25Ti 25Ta. In order to address this discrepancy and to ascertain whether /3 or ~ was stable in Alloys 3 and 49 at 1723 K, a series of intermediate temperature anneals were conducted (both above and below 1723 K). Specimens of Alloy 3 were annealed at 1683 K for 1 h in order to eliminate some of the casting segregation. After WQ, the specimens were then re-inserted into the furnace at the desired temperatures and annealed for an additional one to five hours prior to WQ. Microstructural analyses of these specimens using X R D , E P M A and T E M indicated the presence of

Fig. 10. BETEM of Alloy 3 ( D H + 1723K/5h/WQ) showing lamellar ~ + 7 and )'m decomposition products.

2632

WEAVER and K A U E M A N : 1600 --

oc

1550 -

v

1500 -

(a) × p •

o

I¢.

1300

I 35

30

,,I c~

I

I

1

I

40

45

50

55

Atomic % AI 1600 1550 --

P

1500

1400 -1350 --

E ~

m, -

x

1873

-

-- 1823

81 'X'~

-- 1723

ol

--

\

~

1673 E

--

I

I

I

I

I

10

15

20

25

30

130o

~

1773

--

1450 --

[..,

1673

-- 1623 1573

--

e-~

~

60

(b)

c

1823

-

°/•

1350 --

--

-- 1723

~×

1400 -

1873

-- 1773

x ¢-

1450 --

--

1623

1573 35

1550

-

(c)

1873

-

-- 1823

~,

-

1773

~

1450 --

-

1723

~

1400 --

-

1673

-

1623

~ E ~

1500

--

-~×

--

1350 --

13oo 10

I

I

I

I

I

15

20

25

30

35

I

i

40 45

Fig. 12. Backscattered electron micrograph of Alloy 26 after 1723 K/72 h/WQ revealing the presence of B2 (low Z) and a (high Z).

~

Atomic % Ti 1600

T E R N A R Y AI Ti-Ta SYSTEM

1573

w i t h the lattice p a r a m e t e r o f the fl-Ti solid solution. T h e p r e s e n c e o f the t h e r m a l A P B s in this p h a s e indicates t h a t d i s o r d e r e d fl is p r e s e n t in e q u i l i b r i u m with cr at 1723 K a n d t h a t this fl t r a n s f o r m s to B2 u p o n quenching.

3.1.2.3. Water quenching from 1623K. observations. A f t e r W Q , A l l o y s 2, 3, 4, a n d 49 exhibited c o a r s e d u p l e x m i c r o s t r u c t u r e s w h i c h were d e t e r m i n e d , u s i n g r o o m t e m p e r a t u r e X R D , to c o n s i s t o f ~r a n d 7. I n Alloy 2, after 1623 K / 2 4 h / W Q , T E M analysis (Fig. 14) s h o w e d t h a t b o t h p h a s e s were free o f a n y o f the a f o r e m e n t i o n e d Mierostruetural

50

Atomic % Ta 4 Fig. II. Composition of phases observed in Alloy 3 at various temperatures. (a) A1 content: (b) Ti content; and (c) Ta content.

+ cr below 1673 K and ~ +/3 above 1697 K (up to the fl transus). The EPMA results are shown in Fig. 11. Thus, the stable phases in Alloys 3 and 49 at 1723 K are ~ and /3, consistent with the observations of Boettinger et al. [16]. It is speculated that the microstructures observed by McCullough et al. [11, 12] are due to the presence of cr in the cast microstructures. This ~r was presumably unable to dissolve during their heat treatments due to a combination of sluggish kinetics and low driving force for ~r dissolution. The microstructure observed in Alloy 4 after WQ from 1723 K was identical to the one observed after quenching from 1823 K (i.e. transformed/3) indicating that this alloy is single phase fl at 1723 K. The microstructure of Alloy 26 after 1723 K/72 h/WQ is shown in Fig. 12. Compositional analysis of the two phases indicated that the likely phases present were/~ and ~. This was confirmed by subsequent room temperature XRD and TEM (Fig. 13) which revealed the presence of the ordered B2 (low Z) and ~ (high Z) phases after quenching. The lattice parameter of the B2 phase, as determined from the XRD patterns, was a = 0.3228 nm which is consistent

( lo'I')y//0"3o)o• ° °

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o=

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Fig. 13. BFTEM micrograph and SADPs for Alloy 26 after 1723 K/72 h/WQ showing the presence of B2 and ~. (a) D F T E M showing APBs which are indicative of an ordering reaction during cooling; (b) [001]B2 zone axis; (c) orientation relationship between B2 and tr phases [i.e. [11 I]B2(IOT)B2// [1 12]~[]'10).]; and (d) schematic diagram of (c).

WEAVER and KAUFMAN:

TERNARY Al-Ti Ta SYSTEM

b

Fig. 14. Backscattered electron micrograph of Alloy 2 after 1623 K/24 h/WQ showing ~ (high Z) and 7, (low Z) phases.

Fig.

15.

BFTEM

micrograph of 1623 K/24 h'rWQ.

Alloy

2

2633

after

transformation products signifying that the equilibrium phases in this alloy are ¢r and 7 at 1623 K. On the other hand, in Alloys 3, 4, and 49, 7 was observed in both lamellar and multivariant forms similar to the massive 7m described previously. This indicates that the low Z phase present in these alloys at 1623 K is (Fig. 15) which transforms during W Q to 7'- The phase compositions for Alloys 3, 4, and 49, determined using E P M A , were consistent with those of and/3. Thus, it is concluded that the stable phases in Alloys 3, 4, and 49 are ce and a at 1623 K. Two-phase microstructures were observed in Alloy 26 after heat treating at 1623 K followed by WQ. These phases were identified using X R D as ~r and /3/B2. In addition, the compositions of the phases observed in both samples were consistent with those Fig. 16. SEM backscattered TEM electron micrographs of Alloy 3 after 1823 K/5 h/AC. (a) SEM backscattered electron micrograph showing the as-cooled structure of Alloy 3; (b) TEM bright field micrograph showing the lamellar c~2+ 7 colonies; (c) TEM bright field micrograph showing the coarse a + ;, colonies and the divorced region; (b) SADP showing an orientation relationship ([001]~(ll0)~//~011),. {11 ]}~.) between er strands and the 7 phase observed in Allo~, 3 after 1823 K/5h/AC: and (e) schematic representation of (d).

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2634

WEAVER and KAUFMAN:

TERNARY A1 T i T a SYSTEM

Fig. 17. Optical micrograph showing the microstructure observed in Alloy 2 after air cooling from 1723 K. Note the presence of polycrystalline and lamellar regions similar to those observed after water quenching from the same temperature. of the fl and a phases observed in Alloy 26 after 1723 K72 h/WQ. Based on these observations, it is concluded that fl and a are stable at 1623 K. It is also likely, based on the observations of Das et al. [7 10], that this fl orders to B2 during W Q as T E M analysis was not performed on this alloy.

3.1.3. Air cooled samples 3.1.3.1. Samples air cooled from 1823K. Optical, XRD and SEM observations. The results for the W Q specimens indicated that fl, fl + c~, and + a, respectively, are the stable phases in Alloy 3 at 1823, 1723 and 1623 K. The microstructure observed in Alloy 3, after 1823 K/5 h / A C is shown in Fig. 16(a). Although this microstructure originated from a high temperature fl phase, room-temperature X R D revealed that a, 7 and a2 were present after cooling. EPMA and TEM observations. Subsequent T E M / S A D analysis showed the ), phase to be distributed either as equiaxed grains, as lamellae in ~2 + 7 colonies [Fig. 16(b)], or as lamellae in coarse + 7 colonies which were enclosed by a continuous envelope of a [Fig. 16(c)]. The continuous a layer is consistent with a divorcing type reaction in which the a ---* a + 7 transformation occurs after the ~ ---, ~ + 7 reaction and, therefore, the a near the a/a + 7 boundaries grows on the existing ? resulting in the divorced structure. The 1' within the ~ + 7 colonies bore no apparent crystallographic relationship to the lamellar ~2 + )' colonies. Some a was observed in the form of precipitates at 7 grain boundaries or as "strands" within the polycrystalline 7 and within the ~2 + 7 plate regions where cr appears to grow preferentially into the high Z ~ lamellae. These a strands and precipitates exhibited the following crystallographic relationship with the adjacent -; phase [Fig. 16(d)]

(OOl]o{l lO)~//(ol I), {lii}~. A comparison of this microstructure and the one observed in Alloy 3 after 1723 K / W Q suggests that

Fig. 18. Microstructures observed in Alloys 2, 3, 4 and 49 after 1823 K5 h/FC. (a) Alloy 2; (b) Alloy 3; (c) Alloy 4; and (d) Alloy 49.

WEAVER and KAUFMAN:

TERNARY AI Ti Ta SYSTEM

2635

Fig. 19. BFTEM micrograph showing the "; + a morphology in regions A and B of Alloy 2 after 1823 K5 h/FC. Microstructures of this type were consistently observed in Alloys 3, 4 and 49 after 1823 K/5 h/FC.

Fig. 20. Microstructures of Alloy 46 after 1823 K/5 h/FC showing lamellar ), and a without a continuous intercolony layer and coarse a grains.

during cooling this alloy passes first through an ~ + fl phase field. The c~ and fi phases then transform via different mechanisms during further solid-state cooling. The strands and precipitates of a observed in the % + 7 lamellar regions likely formed via the following reaction sequence

correspond with those of the ~ and fl phases reported for the ST and W Q alloys. The coarse ~ + 7 colonies evolved from the ,q phase independent of the reaction sequence described above, as discussed in greater detail below. Similar microstructures were observed in Alloys 4 and 49 after A C from 1823 K. A single phase fi microstructure was observed in Alloy 26 after AC which is an indication of the greater stability of the 13 phase and/or the sluggishness of its decomposition due to its higher tantalum content. 3.1.3.2. Samples air cooled from 1723 K. Optical, X R D and S E M observations. Alloy 2 was air cooled from 1723K to further elucidate the transformation sequence from the high temperature phase field and the resulting microstructure, which is shown in Fig. 17, consisted of two phases which were compositionally consistent with the 7 phase observed after W Q from 1723 K and the a phase observed after W Q from 1623 K. The presence of a and ~, were confirmed via room-temperature X R D . Optically, the cr phase was found to be distributed along prior c~ grain boundaries as precipitates and as strands

~rrC( ~-'}' ~-O'p----+,3{2-1-3' -}- O'p +), % + 7 ---' ~_, + 7 + % where the ~ ---+c~ + 7 transformation is similar to that observed in binary Ti AI alloys [3, 5] and ~p represents the ~ strands and precipitates. In the polycrystalline regions all evidence would suggest transformation via an intermediate 7 phase field as follows ~ - - , ~ + 7 - - - , 7 - - + 7 + a vPerhaps the most convincing evidence in support of these hypotheses is provided in Fig. 16(d-f) and in the results for specimens of Alloy 2 air and furnace cooled from 1723 K (described below) which clearly show the nucleation and growth of O'p in the higher Ta ~ (or ~:) lamellae and from the prior c~ grain boundaries. It is not known if the % forms prior to the ordering of the ~ to ~2. E P M A analysis revealed the average compositions of the partially transformed matrix and the coarse +7 colonies to be 51AI 27Ti 23Ta and 4 5 A l ~ 5 T i - 3 0 T a respectively. These compositions Table 4. Compositions of regions A and B for alloys FC from 1823 K Alloy 2 3 4 49

Region

%AI

%Ti

%Ta

A B A B A B A B

47.76 51.97 44.62 50.56 41. I I 46.35 47.20 50.30

18.35 28.83 25.10 26.75 22.97 31.71 28.46 30.63

33.89 19.20 30.28 22.69 35.92 21.94 24.34 19.07

Fig. 21. BFTEM micrograph of Alloy 4 after 1723 K5 h + FC to 1623 K + WQ showing a and y growing into ft.

2636

WEAVER and KAUFMAN: TERNARY A1-Ti Ta SYSTEM

Fig. 22. Secondary electron image of Alloy 26 after 1823 K5 h/FC. The arrow indicates a multiphase mixture of ~, 7, and ~2. growing out into the ~ grains. The prior :~ grains exhibited a coarser version of the massive transformation product observed in Alloy 2 after WQ from 1723 K. This observation would suggest that during AC, the high temperature c~ phase first transforms massively to 7. This massive transformation product then coarsens with time. Finally the cr phase nucleates and grows during the latter stages of this transformation at prior ~ phase boundaries and grows into the 7.

The ~r +~, transformation product observed in region A was surrounded by a continuous layer of a. In order to determine the nature of this layer and to further delineate the decomposition behavior of the fl phase, two additional experiments were conducted. First, Alloy 46, which lies in the fl + e region at 1723 K (described later) was annealed at 1723 K for 108 h and furnace cooled• The resulting microstructure (Fig. 20) consists of coarse a grains and of lamellar 7 + e colonies similar to those observed in regions A and B in Alloys 2, 3, 4, and 49 (1823 and 1723 K/FC). The coarse colonies observed in the other alloys were not observed here. Compositional analysis of the phases in this alloy showed the a and y phases to be equivalent to those observed in Alloy 2 after FC from the c~ phase field which is consistent with transformation occurring through an intermediate ~ phase field. In a second experiment, Alloy 4 was annealed at 1723 K for 5h, cooled to 1623 K in l0 min and WQ. Subsequent T E M analysis revealed the presence of ~, fl, ~, and ~ phases with the a and

3. 1.4. Furnace cooled samples 3.1.4.1. Samples" .[urnace cooled.fi'om 1823 K. XRD, SEM and T E M observations. The microstructures observed in Alloys, 2, 3, 4, and 49 after 1823 K/5 b/FC are presented in Fig. 18. X R D and T E M analysis indicated the presence of o-, :~2 and ,,' after cooling and the a and 7 phases were distributed either in coarse ~r-rich colonies (region A in Fig. 19) or in fine ?,-rich regions (region B in Fig. 19). Interestingly, the transformation product in region B of the FC specimen appeared slightly different from that observed in the AC specimen. In the FC case, region B consisted predominantly of a fine lamellar mixture of ~r and 7- Only in limited cases were :~2+ ?' lamellar structures observed. This is most likely due to the greater time for the :~ + "~,(or e2 + 7) structure to transform during FC to an equilibrium mixture of a and 7 whereas AC is too rapid for this transformation to reach completion. The compositions of regions A and B for Alloys 2, 3, 4 and 49 are listed in Table 4 from which it can be seen that region A corresponds compositionally with the high temperature fl phase observed after W Q from 1723 K. As noted in the previous section, these results suggest that the high temperature fl phase partitions during cooling into a mixture of ~ and fl of different composition. The resulting ~ and fl phases next transform via different transformation sequences (suggested above) eventually forming mixtures of a and 7.

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WEAVER and KAUFMAN: TERNARY A1 Ti Ta SYSTEM (a) iagram for

I I

Iloy 2

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FC

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Fig. 24. Schematic transformation diagrams for Alloys 2, 3, 4, 49 and 26. (a) Alloy 2: (b) Alloys 3, 4, and 49; and (c) Alloy 26.

2637

7 phases, Fig. 21, growing into and consuming the fl phase. Adjacent to these regions were lamellar ~2 + 7 regions which again indicates that the high temperature [~' phase transforms first during cooling to a mixture of ~ and fl with the new [I having a different composition than the original ft. The resulting ~ and [~'phases then decompose via two different transformation sequences into fine mixtures of a and y. The microstructure observed in Alloy 26 after FC l¥om 1823 K is shown in Fig. 22. TEM analysis revealed the presence of both or, 7 and ~2 after cooling. Thus, it appears that the er precipitated in the fl grains rejecting lower Z solute into the D'. The remaining fl then transformed into the multiphase mixture indicared by the arrow in Fig. 22. 3. 1.4.2. Samph's furnace cooled Jorm 1723 K. SEM, XRD and TEM observations. In order to gain a more detailed understanding of the phase transformations taking place in region B of Fig. 18, Alloy 2 was annealed at 1723 K (in the ~ field) and FC to room temperature. The resulting microstructure is presented in Fig. 23(a). XRD and TEM analyses confirmed the presence of both o- and 7 [Fig. 23(b and c)]. Significantly, the coarse transformation product observed in region A was absent in this specimen thereby supporting the precipitation reactions mentioned above for region B. Similar to the observations of McCullough et al. [11, 12], the oand 7 phases were related by the same orientation relationship described above. By combining these results, it is possible to generate schematic transformation diagrams for Alloys 2, 3, 4, 26, and 49 (Fig. 24). The diagram for Alloy 2 [Fig. 24(a)] represents cooling from the ~ phase field while the remaining diagrams [Fig. 24(b and c)] represent cooling from the fl field. As already shown, the [I phase always decomposed and the resulting microstructures depended strongly on cooling rate and composition. For Alloy 2, :t transforms massively during water quenching from the ~ field while furnace cooling results in the formation of coarse lamellar :¢ + ' / structures. Alloys 3, 4, 26, and 49 lie in the fl phase field at 1823 K while Alloy 2 lies in the two-phase [ / + liquid region. This [/ orders to B2 in Alloy 26 during WQ but transforms martensitically to a Widmanstfitten structure consisting of hexagonal ~2 and B2 phases in Alloys 3, 4 and 49 during WQ. During FC, Alloy 26 decomposes into a coarse two-phase mixture of ~ and ~. while Alloys 2, 3, 4 and 49 transform through an intermediate ~ phase to a coarse mixture of c, and y. These observations confirm the results of Boettinger el al. [16] who observed similar structures in an alloy of composition 50AI 25Ti-25Ta after WQ from 1798 K.

3.2. Phase Equilibria In order to provide a more detailed understanding of the microstructural evolution during cooling, information concerning the stable phases present at

2638

WEAVER and KAUFMAN: TERNARY A1-Ti Ta SYSTEM Ti

e7

7 8"'.

20 3 /

Ta

90

o

80

70

60

\

~

/ ~ in

50 40 Atomic%Ta

30

70 S0

20

10

A1

Fig. 25. Experimentally determined AI Ti Ta isotherm at 1723 K. The dotted lines indicate phase boundaries approximated from Ref. [7 12, 16] while the solid lines indicate the findings of this study. various temperatures and compositions are needed. However, the reactions that occur upon cooling make traditional methods of phase diagram determination difficult. Therefore, the phase transformation information, divulged above, was used in conjunction with the following observations for 23 additional alloys (see Table 1) to construct ternary isotherms at 1623 and 1723 K. TEM, EPMA and XRD were the primary methods used to determine the phases present in this portion of the study and the results are presented below. Where possible, the various transformations observed in the five alloys described above were used as a guide to characterize the structures of the remaining alloys. 3,2.1. Phase equilibria at 1723 K Optical microscopy, room temperature XRD, and TEM analysis of Alloys 6, 7, 8, 32, 37, and 38 revealed the existence of an extensive fl phase field at 1723 K. Specifically, the microstructures observed in Alloys 6, 7, 8, 37 and 38 consisted of large equiaxed grains (mm size) indicating the existence of a single phase region at 1723 K. TEM/SADP analysis indicated that Alloys 6, 7, 8, 37, and 38 had the ordered B2 (CsCI) structure after quenching. Furthermore, all of the alloys contained thermal APBs indicating that these alloys were disordered at 1723 K and ordered during WQ. These findings are consistent with those of Das and Perepezko [7-10] who indicated the presence of a B2 phase field near the composition Ti 25A1 25Ta which existed up to 1473 K but disordered to b.c.c, upon heating above this temperature. Alloy 32 transformed during water quenching yielding a Widmanstfitten-type microstructure consisting of ~ and B2 which were crystallographically oriented via the Burgers orientation relationship, The microstructure observed in Alloy 50 was used to construct the single phase ~ region near the composition AIeTa. TEM analysis by Weaver and

Kaufman [22] revealed that this phase had an f.c.c. crystal structure (space group Fm3m) with a lattice parameter of 1.9286 nm. A more detailed description of this phase is provided in Ref. [22]. Finally, Alloy 18 was found to consist of the tetragonal r/AI3(Ti,Ta) phase while single phase cr was observed in Alloys 33 and 52 after quenching from 1723 K. The two-phase e + fl field was established based on analysis of Alloys 3, 16, 41, 49, and 51 after quenching from 1723 K. The microstructure observed in Alloy 16 consisted of large primary particles (5-15 ~m) and regions containing a fine transformation product. A similar microstructure was observed in Alloy 41. Room temperature XRD indicated the presence of ~2 in both alloys and subsequent TEM analysis was used to confirm the presence of az in Alloy 16. Significantly, however, both the coarse particles and the fine transformed regions contained thermal APBs indicating an ordering type reaction (i.e. ~ - - ~ 2 ) upon cooling. In addition, a small fraction of retained B2 was observed in the transformed region consistent with a transformation from a high temperature fl phase. Based upon this analysis, the phases present in Alloy 16 at 1723 K are ~ and/3. Although TEM was not conducted on Alloy 41, EPMA indicated the presence of two distinct phases which likely correspond to ~ and fl at temperature. The microstructure observed in Alloy 51 consisted of two phases, the EPMA compositions of which were consistent with those of the ~ and fl phases observed in Alloys 3 and 49 which suggest that e and /3 are the stable phases present at 1723 K. /3 and a were found to be the stable phases in Alloy 46 at 1723 K while 7 and ~ were found to be stable in Alloy 1. In Alloy 30, room temperature XRD and EPMA analyses suggested the presence of 6-A12Ta and q. TEM analysis (see Ref. [14]) confirmed these observations indicating that r/ and ,~ are the stable Ti

/I\o ~" ~o,~

Ta

90

80

70

II I~,~o

60

50 40 Atomic%Ta

30

20

10

AI

Fig. 26. Experimentally determined A1 T i T a isotherm at 1623 K. The dotted lines indicate phase boundaries approximated from Refs [7 -12, 16] while the solid lines indicate the findings of this study.

WEAVER and KAUFMAN: TERNARY A1 Ti Ta SYSTEM phases in Alloy 30 at 1723 K. 5 and a were observed in Alloy 53 (1723 K/WQ). Alloys 39 and 47 (1723 K/WQ) consisted of large spherical particles inside a continuous matrix which exhibited clear evidence of melting. XRD analysis of Alloys 39 and 47 revealed that the majority phase present was r/after quenching while the low volume fraction of the prior liquid regions prevented its detection. Consequently, it can be concluded that the stable phases present at 1723 K are t1 and liquid. Alloy 25 (1723 K/WQ) also exhibited clear signs of grain boundary liquidation. The grain interiors, however, exhibited transformed morphologies which were similar to those observed in binary AI -Ti alloys WQ from the ~ field. In addition, room temperature XRD analysis revealed the majority phase to be 7 after quenching which suggests that this alloy lies in the + L phase field at 1723 K, and that the ~ transforms during quenching to ? while the liquid solidifies in a manner analogous to binary Ti AI alloys (see McCullough et al. [11, 12]). The phase compositions were determined using EPMA. The remaining two-phase regions (i.e, ~. + q, ~ + 7, Y + L, and r / + 7) were not observed in the samples examined yet were approximated based upon the low temperature observations of McCullough et al. [11, 12], Boettinger et al. [16], and Das and Perepezko [7 10] and upon the results of the present study. The three-phase :~+~1+,' and ~ + ~ 1 + ~ fields were established based on EPMA. XRD, and LOM analysis of Alloys 40 and 54 respectively. The remaining three-phase regions (i.e. : + ;' + L, ~ + fi + o-, + 5 + er, and 7 + r/ + L) were estimated based upon previous literature [7 16, 23, 24]. The various results described above have been plotted on the ternary isotherm shown in Fig. 25 where the solid dark lines indicate the findings of this study. The remaining lines indicate approximate phase boundaries based upon computer models [23] and the lower temperature work of McCullough et al. [11, 12], Das and Perepezko [7 10], and Boettinger et al. [16]. It is important to note that no evidence of the binary cAITa phase reported by Subramanian et al. [6], was observed at this temperature. 3.2.2.

1623 K phase equilibria

Alloys 1, 2, 3, 4, 26, 45, 49, and 50 were annealed at 1623 K followed by WQ or AC to estimate the phase equilibria at 1623 K. None of the alloys held at 1623 K were single phase. However, it was possible to approximate the locations of the single phase fields based upon the tie-line and tie triangle locations for the eight alloys investigated. Five single phase regions (i.e. :< fl, 7, a, and r/) have been identified using these methods. The ~ + 7 phase field has been established based on EPMA and TEM analysis of Alloy 2 (1623 K/ 43h/WQ) and Alloy 45 (1623K/24h/WQ) which consisted of coarse distributions of these phases. In both alloys, the a and 7 lacked any of the afore-

2639

mentioned transformation products which indicates that they are in equilibrium at 1623 K. In Alloys 1 and 50, duplex microstructures consisting of r/and ~ were observed after 1623 K/48 h/WQ. None of the alloys studied contained more than two phases. Consequently, the three-phase regions were estimated based on the lower temperature observations of Boettinger et al. [16], McCullough et al. [l l, 12], and Das et al. [7 10]. The results described above have been plotted on the ternary isotherm shown in Fig. 26. Again, the dark lines indicate the findings of this study and the dotted lines indicate phase boundaries approximated from Ref. [7-11, 16]. 4. DISCUSSION This study has demonstrated that Alloys 2, 3, 4, 26, and 49 solidify through a double peritectic reaction involving the ~ and [3 phases with fl being the primary phase. This is in agreement with McCullough et al. [11, 12] who reported similar observations for an alloy near the composition 45A1-28Ti-27Ta. The [~ phase in these alloys was also observed to be heavily cored and transformed during cooling with the resulting microstructure dependent upon the local Ta content. In addition, the alloys slowly cooled from the fi field were consistently observed to transform by passing through two three-phase regions (i.e. [~' + ~ + ~ and ~ + a +7)4.1. [J Decomposition

It has been shown that alloys cooled from the fl field do indeed transform via the initial precipitation of ~ as indicated by Boettinger et al. [16] and McCullough et al. [11, 12]. It is additionally shown that the fl and ~ phases decomposed during further FC via the following reaction sequences

~

c~---+~L + 7L ---+:~L+ 7L + aC2 --'+aC: + 7C:

[ ~ + [ , f l - - ~ f l l + a---+ a ~+ aCl +~'Cl

where ~L + YL, C*CJ+ YC~, and acz + 72 correspond to lamellar ~/~z + 7, cellular a + 7, and lamellar a + y morphologies of differing composition, respectively, which form during cooling. Evidence in support of these decomposition sequences are provided in Figs 18 and 23 which show the microstructures in Alloy 2 after FC from the single phase fl and ~ phase fields respectively. This transformation sequence is supported by the observations of Boettinger et al. [16] who performed an extensive annealing study on an alloy of composition 50.3A1-25.8Ti-25.9Ta. 4.2. Phase Equilibria

The phase equilibria at 1623 and 1723 K have been defined in the range 38 55%A1, 19 30%Ti and 20 38%Ta resulting in the identification of threephase fields involving the zq //, and a phases at

2640

WEAVER and KAUFMAN:

elevated temperatures. In addition, a discrepancy concerning the phases present at 1723 K in alloys near the composition 45AI-28Ti 27Ta has been clarified. This study has shown that a n alloy o f this c o m p o s i t i o n should lie in the single phase /7 field at 1723 K r a t h e r t h a n in the two-phase c~ + / ~ field as reported by M c C u l l o u g h et al. [11, 12]. The observations of M c C u l l o u g h were made using a combin a t i o n of high t e m p e r a t u r e X R D and E P M A where the specimen was heated stepwise in the X-ray diffractometer a n d held at temperature for 15 rain prior to analysis. A l t h o u g h this m e t h o d has been used successfully in Ti AI alloys [3], it is suggested by the authors that a holding time of 15rain prior to analysis may be insufficient to reach equilibrium due to the presence of Ta which slows the t r a n s f o r m a t i o n kinetics of the alloys. It is likely that the multi-phase alloys observed by M c C u l l o u g h are due to failure to eliminate casting segregation prior to examination. N o such two-phase alloys were observed in this c o m p o s i t i o n range. 5. CONCLUSION The results of this study were used to construct two ternary isotherms (1623 a n d 1723 K) and to deduce the following conclusions. 1. The /~ phase in high Ta alloys (e.g. ~ 3 6 % T a ) orders to B2 d u r i n g W Q or AC. However, d u r i n g F C this phase transforms to a duplex mixture of~r and/~22. The /~ phase in the Ta-lean alloys (~<25%Ta) transforms martensitically to a mixture o f ~ and B2 u p o n W Q from 1823 K. During FC a n d A C from this same temperature, ~ first precipitates within the phase. The Ta-rich [~ phase then proceeds to transform to a which subsequently decomposes cellularly into a lamellar mixture o f a and ?. The ~ phase also decomposes initially into a lamellar mixture of ~ + ), which, at low cooling rates, decomposes further into a lamellar 7 + cr mixture. This a and S' divorces locally at the prior [~/~ phase boundaries due to the dec o m p o s i t i o n of the ~ phase into 7 during cooling. 3. D u r i n g WQ, the ~ phase frequently transforms by the '+massive" ~--~ ,'m reaction. During AC and FC however, the ~ transforms to ~2 + 7 plates with a precipitating subsequently as " s t r a n d s " within this 7 at the ~/7 interface. 4. In agreement with Boettinger et al. [16], the stable phases at 1723 K in an alloy o f composition AI 25Ti 25Ta have been determined to be ~ a n d / 3 rather t h a n a a n d [~ as previously reported by Weaver et al. [14].

Acknowledgements Work supported through the DARPA-AMMP program on composites to the University of Florida under grant No. MDA972-88-J-1006 and through the NASA Lewis Research Center under grant No. NGT-50631. In addition, the authors would like to thank Robert M. Dickerson, Patricia O. Dickerson, Sandra

TERNARY A1 Ti Ta SYSTEM L. Guy, Mark A. Falvai, Ron K. Stone, and Robert E+ Williams who performed some of the sample preparations and analyses reported in this study and Joe Wagner for the DTA analysis. REFERENCES

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