Experimental study and thermodynamic modelling of the ternary Al–Ta–Ti system

Intermetallics 19 (2011) 234e259

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Experimental study and thermodynamic modelling of the ternary AleTaeTi system V.T. Witusiewicz a, *, A.A. Bondar b, U. Hecht a, V.M. Voblikov b, O.S. Fomichov b, V.M. Petyukh b, S. Rex a a b

ACCESS e.V., Intzestr. 5, D-52072 Aachen, Germany Frantsevich Institute for Problems of Materials Science, Krzhyzhanovsky Str. 3, 03680, Kyiv, Ukraine

a r t i c l e i n f o

a b s t r a c t

Article history: Received 30 July 2010 Accepted 2 October 2010 Available online 20 November 2010

In the present paper a thermodynamic description of the entire ternary AleTaeTi system is proposed, being obtained by CALPHAD modelling. Dedicated experiments were performed in order to complement literature data on phase equilibria reported mainly for the temperature interval from 1273 to 1723 K. Ascast and annealed samples were investigated by means of DTA, pyrometry after PiranieAlterthum, SEM/ EDS, SEM/EBSD and XRD techniques in order to prove the phase relations in ranges of invariant equilibria at melting, a2 þ g alloys and low-temperature ternary phases. The experiments revealed that two ternary phases exist in the AleTaeTi system: the orthorhombic O-phase (Ti2.17Ta0.77Al1.06, Pearson symbol oC16) and the ordered B82-phase (Ti3.00Ta0.89Al2.11, Pearson symbol hP6) are thermodynamically stable below 1123 K and 1171 K, respectively. The obtained and literature data were employed throughout the optimization work, using the PARROT module of Thermo-Calc. The elaborated thermodynamic description was applied to calculate selected phase equilibria as to provide a comparison between calculated and experimental data. The calculations are shown to reproduce correctly experimental data. They revealed the presence of a miscibility gap for the ternary (bTi,Ta,Al) solid solution in the temperature interval from 1461 to 813 K. Ó 2010 Elsevier Ltd. All rights reserved.

Keywords: A. Aluminide B. Phase diagrams B. Thermodynamic and thermochemical properties E. Phase diagram, prediction

1. Introduction The AleTaeTi system attracted recent interest, since tantalum containing titanium aluminide alloys, e.g. Tie46Ale8Ta (at.%), can be grain-refined via massive transformation upon air cooling [1e3]. Alloy processing by casting and heat treatment and various aspects of microstructure stability at envisaged service temperatures can benefit from the knowledge of phase equilibria in the wide temperature range that spans from the liquid state down to room temperature. A comprehensive analysis of the earlier thermodynamic description for the AleTaeTi system [4] proved that it does not correctly reflect the present status of experimental knowledge. This regards liquidus temperatures, fields of primary solidification, invariant reactions, element partition between different phases, etc. The efforts presented here aimed to improve the thermodynamic description of the entire AleTaeTi system. The work encompasses experimental investigations and the thermodynamic modelling. Published experimental data on phase equilibria for the AleTaeTi system were assessed and complemented by own experiments in

* Corresponding author. Tel.: þ49 241 8098007; fax: þ49 241 38578. E-mail address: [email protected] (V.T. Witusiewicz). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.10.002

critical composition-temperature ranges. Modelling was performed based on the thermodynamic description of the binary constituent systems AleTa [5], AleTi [6] and TieTa adopted from Ref. [7]. The present paper is structured as follows: in Section 2 experimental data are summarized with main emphasis on the measurements performed within the frame of this work. Section 3 contains a brief description of the thermodynamic models used for expressing the Gibbs free energy of the individual phases of the alloy system and presents the optimization procedure. The resulting thermodynamic database is given in Appendix. Section 4 contains a large number of calculations performed with the elaborated database, including liquidus, solidus and b-transus projections, isothermal sections and selected isopleths. The calculations are discussed in comparison to experimental data regarding the phase equilibria of major interest.

2. Experimental information for the ternary AleTaeTi system 2.1. Literature data 2.1.1. Solid phases Crystallographic data of all solid phases in the AleTaeTi system are listed in Table 1. The AleTaeTi system was reviewed in

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235

Table 1 Phase designations most often used in literature for the AleTaeTi system along with crystal structure data [8e13] and thermodynamic models used in the present description. Phase (Designation)

Pearson symbol

Space group

Struktur bericht designation

Prototype

Model used in the present description

La, liquid (Al),(aAl), fcc_A1 a, (aTi), hcp_A3 a2, Ti3Al b, (bTi), (Ta), bcc_A2 b0, b0 , bcc_B2 g, gTiAl, TiAl z, Ti2þxAl5x,Ti5Al11 h, TiAl2 Ti3Al5 3, (Ti1-xTax)Al3, TiAl3(h), TaAl3 3(l), TiAl3(l) k, d,Ta39Al69 4, Ta48Al38,TaAl s, Ta2Al, Ti1xyTaxAly sb, Ti3Ta1xAl2þx at x ¼ 0.11 Ob, Ti2.17Ta0.77Al1.06

cF4 hP2 hP8 cI2 cI2 tP4 tP28 tI24 tP32 tI8 tI32 cF432 mP86 tP30 hP6 oC16

Fm3m P63/mmc P63/mmc Im3m Pm3m P4/mmm P4/mmm I41/amd P4/mbm I4/mmm I4/mmm F43m P21/c P42/mnm P63/mmc Cmcm

A1 A3 D019 A2 B2 L10 e e e D022 e

Cu Mg Ni3Sn W CsCl AuCu Ti2Al5 HfGa2 Ti3Al5 TiAl3(h) TiAl3(l) Ta39Al69 Ta48Al38 sCrFe Ni2In NaHg

[Al,Ta,Ti] [(Al,Ta,Ti)1:(Va)1] [(Al,Ta,Ti)1:(Va)0.5] [(Al,Ta,Ti)3:(Al,Ta,Ti)1] [(Al,Ta,Ti)1:(Va)3] [(Al,Ta,Ti)1:(Va)3] þ [(Al,Ta,Ti)0.5:(Al,Ta,Ti)0.5:(Va)3] [(Al%,Ta,Ti)1:(Al,Ta,Ti%)1] [(Al,Ta,Ti)5:(Al,Ta,Ti)2] [(Al,Ta,Ti)2:(Al,Ta,Ti)1] [(Al)5:(Ta,Ti)3] [(Al, Ti)3:(Al,Ta,Ti)1] [(Al,Ti)3:(Al,Ta,Ti)1] (AL%,Ta,Ti)0.6389:(AL,Ta%) 0.3611 (AL,Ta,Ti) 0.8837:(AL,Ta,Ti) 1.1163 [(Al,Ta,Ti)0.533:(Al,Ta,Ti)0.333:(Ta,Ti)0.134] [(Ti)3:(Al,Ta,Ti)1:(Al)2] [(Al,Ta,Ti)0.75:(Al,Ta,Ti)0.25]

a b

D8b B82 e

Designations given in bold letters are used throughout the present work. Ternary phase discovered within the frame of the present work.

[10,11,13e15]. The authors of Refs. [11,13] provide a critical evaluation of experimental data on phase relationships in the AleTaeTi system and a recommended set of internally consistent phase diagrams, accepting the constituent binary AleTa, AleTi and TaeTi

phase diagrams from Refs. [16,6,8], respectively. According to their assessments, no ternary compounds exist in the system. As reported, there are solid solution phases, b or (bTi,Ta,Al) above 1155 K and 3 or (Ti,Ta)Al3 above 997 K [17e23]. The bcc b-phase extends up

Table 2 Sample compositions and phase identification in as-cast and annealed conditions. Sample no.

Composition, at.% Nominal

XRD, SEM/EDS and SEM/EBSD phase identification EDS data

As-cast condition

Annealed condition

Ti

Ta

Ti

Ta

Primary phase

Phases

T, K

Time, h

Phases

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15

2.5 4.2 10.2 12.0 15.0 17.5 19.0 20.5 22.2 24.9 29.9 30.0 33.3 41.0 43.0

33.5 36.6 25.6 34.0 45.0 31.0 31.0 31.0 18.8 18.8 4.0 19.0 33.4 12.0 12.0

2.8 4.4 10.5 13.1 16.0 17.7 20.1 21.4 22.4 23.7 28.5 28.2 32.8 41.3 43.4

32.4 36.4 25.7 34.7 45.7 30.9 29.6 30.0 18.7 19.1 3.4 19.1 32.1 11.8 11.8

sa s k s ba ba ba ba aa aa g ba b ba ba

3þk kþs 3þk gþ3þs s gþs gþs gþs g g g þ (z?) g b a2 þ g þ (b?) a2 þ g þ (b?)

16 17

45.0 46.0

8.0 8.0

44.9 46.5

7.7 7.7

ba ba

a2 þ g a2þg

18 19 20 21 22 23

47.0 49.0 51.0 55.0 53.0 49

8.0 4.0 4.0 e e 28

46.9 48.7 50.8 55.6 53.2 49.8

7.7 3.9 3.9 e e 28.0

ba ba ba ba ba b

a2 þ g a2 þ g a2 þ (g?) a2 þ g a2 þ g a2 þ b

24

50

12.5

50.1

11.9

b

a2 þ b

25

50

25

51.0

24.3

b

a2 þ b

26

54

22

54.5

21.8

b

a2 þ b

1773 1773 1713 1773 1773 1773 1833 1833 1713 1713 1663 1773 2073 1703 1703 1023 1703 1703 1373 1023 1703 1703 1703 1703 1703 1973 1123 1023 1873 1123 1023 1973 1123 1023 2023 1123 1023

2 2 2 2 2 2 2 2 2 2 2 2 2 10 10 1082 10 10 24 1082 10 10 10 10 10 7 192 708 7 192 708 7 192 708 7 192 708

3þk kþs gþ3þk L þ sa; 3 þ k þ s bþs L þ a þ sa; a þ s aþs aþbþs gþk aþgþs g a b gþb a2 þ g þ b a2 þ g þ s þ s a2 þ g a2 þ g a2 þ g a2 þ g þ s a2 þ g a2 þ g a2 þ g a2 þ g a2 þ g b a2 þ b þ O b a2 þ b þ O b b a2 þ g þ sb a2 þ g þ sb b a2 þ b þ O b a2 þ b þ O b b a2 þ b þ O b a2 þ Ob

a b

Identified based on SEM observations of morphology of dendrites or primary particles and EDS measurements of their compositions. SEM/EBSD result.

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Fig. 1. SEM/BSE micrographs of as-cast microstructures with different primary phases: (a) s-phase in Ti4.4Ta36.4Al59.2, (b) k-phase in Ti10.5Ta25,7Al63.8, (c) a-phase in Ti22.4Ta18.7Al58,9, (d) g-phase in Ti28.5Ta3.4Al68.1 and (e) b-phase (transformed in the course of cooling) in Ti46.9Ta7.7Al45.4.

to about 40 at.% Al at 1373 K and from 42 to 50 at.% Al at 1623 to 1723 K [22e25]. In Refs. [22,23] also the A2/B2 ordering of the bphase was revealed. In Ref. [25] the ordering was analyzed based on anti-phase boundaries observed in ternary alloys quenched from 1623 to 1723 K. According to Refs. [22,23,26] the A2/B2 ordering of the b-phase at 1373 K takes place within the composition range of about 40e60 at.% Ti and 15 to 60 at.% Al.

A single-phase region around the Ti4TaAl3 composition is assumed by Refs. [10,22,23] to be a new ternary compound. It was observed both in bulk alloys annealed at 1373 K and also in some of the diffusion couples [10]. However, it was not taken into account in the assessments [11,13,15]. Recently, the ternary compound close to the stoichiometric composition Ti4TaAl3 has been observed by Lapin et al. [27] in Tie46Ale8Ta (at.%) after long-time aging at

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Fig. 2. SEM/BSE micrographs of selected samples solidified with primary b-phase showing its decomposition in the course of cooling in the as-cast (left column) and annealed (right column) samples.

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Table 3 XRD data of as-cast AleTaeTi alloys melted to study invariant equilibria at solidus temperatures. Sample no.

Phase constituents from XRD

1

kþ3

2 3

kþs kþ3

4

sþgþ3

5 6

s gþs

7

gþs

8

gþs

9 10 11 12 13

g g g þ (z?) g b

Lattice parameters, pm Phase

a

c

k 3 k k 3 s g 3 s g s g s g s g g g g b

1916.9 385.0 1917.7 1919.3 387.0 981.6 398.8 386 986.8 398.8 980.9 398.8 981.6 399.3 983.4 396.7 398.5 393.7 399.4 322.6

e 854.0 e e 852.0 519.9 410.2 858 517.4 410.9 520.7 410.2 519.9 410.1 519.0 411.8 410.4 413.4 409.4 e

Table 4 Average composition of phases in multiphase samples measured by EDS. Sample no. History

Phase

Composition (at.%) Al

1 2 3

4 5 6

7 8

9 10

15 17 24

25

3 k s k g 3 k

73.8 61.6 46.2 1773 K/2 h 58.9 57.7 1713 K/2 h 70.2 58.0 1773 K/2 h Liquid 56.9 s 43.3 As-cast Eutectic b þ s 53.8 As-cast Eutectic a þ s 57.0 1773 K/2 h Liquid 55.0 a 51.5 s 41.7 a 50.3 1833 K/2 h s 44.2 a 50.5 1833 K/2 h b e s 43.7 g 57.3 1713 K/2 h k 57.1 a 52.1 1713 K/2 h g 57.3 s 45.4 49.0 1023 K/1082 h g s 35.7 48.3 1023 K/1082 h g s 35.2 37.1 1123 K/192 h a2 g 44.9 s 34.5 1023 K/708 h a2 37.3 g 46.9 s 34.9 27.0 1123 K/192 h a2 b 23.1 O 26.5

1773 K/2 h

Ti                 

0.2 0.2 0.4 0.4 0.4 0.2 0.3 0.3 0.3 1.1 1.3 0.6 0.3 1.0 0.4 0.6 0.8

                  

0.4 0.3 0.3 0.5 0.2 0.6 0.7 0.6 0.4 1.1 2.8 0.4 0.6 0.7 0.5 0.5 0.2 0.2 0.3

3.7 2.1 3.0 4.3 22.0 11.4 8.7 18.1 8.1 29.9 24.9 19.8 23.2 9.5 22.9 7.5 22.7 e 8.1 22.3 8.4 23.5 24.2 5.3 43.5 49.6 45.1 50.5 52.9 47.3 50.5 56.2 49.3 50.6 60.4 47.4 54.2

Ta                 

0.2 0.1 0.3 0.2 0.3 0.3 0.3 0.3 0.4 0.8 1.2 0.3 0.3 0.7 0.8 0.3 1.0

                  

0.6 0.3 0.2 0.5 0.4 0.3 0.8 0.5 0.4 0.9 2.2 0.3 0.5 0.6 0.6 0.6 0.2 0.3 0.9

22.5 36.4 50.8 36.8 20.3 18.4 33.3 22.3 48.2 16.3 18.1 25.2 25.3 48.8 26.8 48.3 26.8 e 48.2 20.4 34.5 24.4 18.5 49.3 7.5 14.7 6.6 14.7 10.0 7.8 15.0 6.5 3.8 14.5 12.6 29.5 19.3

                

0.3 0.3 0.7 0.3 0.3 0.3 0.2 0.2 0.5 0.8 0.9 0.4 0.4 1.6 0.5 0.7 0.7

                  

0.9 0.4 0.3 0.6 0.3 0.4 0.7 0.5 0.4 0.5 1.3 0.3 0.4 0.5 0.4 0.5 0.3 0.2 0.9

1023 K. As it was established by XRD, the crystal structure of this phase corresponds to B82 [27] and is identical to the s-phase of the AleNbeTi system [12]. Within the frame of the present work, in the AleTaeTi alloys annealed at 1023 and 1123 K, this compound was also experimentally confirmed to exist as Ti3.00Ta0.89Al2.11, and a second ternary phase with orthorhombic structure (the O-phase of NaHg crystal structure type) of Ti2.17Ta0.77Al1.06 composition was revealed for the first time. This phase was found to be identical to the O1-phase (Ti2NbAl) of the AleNbeTi system [12].

Fig. 3. XRD diffractograms of samples with 11.8 at.% Ta before and after annealing at sub-solidus temperature: (a) sample Ti43.4Ta11.8Al44.8 in the as-cast condition; (b) sample Ti43.4Ta11.8Al44.8 after annealing at 1703 K for 10 h and cooling with a rate of 3 K s1; (c) sample Ti41.3Ta11.8Al46.9 in the as-cast condition and (d) sample Ti41.3Ta11.8Al46.9 after annealing at 1703 K for 10 h and subsequent cooling with a rate of 3 K s1. For clarity, the curves have been shifted vertically.

2.1.2. Reaction scheme and liquidus surface No experimental information concerning the invariant equilibria in the AleTaeTi system is available from literature. A tentative reaction scheme was proposed by Ref. [15] based on limited experimental data on the univariant L þ a þ b, L þ a þ g and L þ g þ 3 liquidus lines [28e30], isothermal sections from 1713 K to 1273 [23e25,31,32] and the constituent binary phase diagrams accepted from Ref. [8]. Information on the liquidus surface was limited to the abovementioned univariant lines obtained by Refs. [28e30] and more recent by Ref. [33]. Based on the liquidus lines of Refs. [28e30] and data on the phase equilibria at 1713 K of Ref. [23], six invariant reactions were proposed by Ref. [15] and used to provide a tentative liquidus projection for the entire system. Recently Shuleshova et al. [34] reported the results of in situ synchrotron X-ray diffraction experiments that shed light on high-temperature phase transformations in Tie43.9Ale7.2Ta, Tie45.2Ale7.0Ta, Tie48.9Ale30.6Ta, Tie55.6Ale17.1Ta, Tie58.4Ale14.7Ta (all in at.%). The b-phase was

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Fig. 4. SEM/BSE micrographs of samples Ti51.0Ta23.4Al25.6 (no. 25), annealed at 1123 K for 192 h (a) and at 1023 K for 708 h followed by water-quenching (b), and Ti50.1Ta11.9Al38.0 (no. 24), annealed at 1123 K for 192 h (c) and at 1023 K for 708 h followed by water-quenching (d): the samples of Ti51.0Ta23.3Al24.7 composition exhibit plate-like precipitates of the Ophase embedded in a b-matrix; in Ti50.1Ta11.9Al38.0 there are plate-like g-precipitates in a s-matrix.

shown to be the primary solidification phase for the four first alloys, while the last alloy solidified to yield primary a-phase. It was also shown that the location of the L þ a þ b univariant line agrees well to the data from Ref. [33] and somewhat differs from Refs. [30] or [15].

2.1.3. Isothermal sections and isopleths The isothermal sections at 1723, 1713, 1623, 1603, 1573, 1473, 1373 and 1273 K were experimentally studied in Refs. [17,18,20e23,25,31,32]. Despite of considerable scatter regarding

Table 5 XRD data of AleTaeTi samples annealed at 1023 K for 708 h followed by water-quenching or annealed at 1123 K for 192 h followed by furnace cooling. Sample no.

History

Composition, at.% Ti

Ta

Phase constituents from XRD

Phase

23

1123 K/192 h/FC

49.8

28.0

b0 þ b00 þ a2 þ Oa

24

1023 K/708 h/WQ 1123 K/192 h/FC

50.1

11.9

b þ O þ a2 s þ g þ a2b

25

1023 K/708 h/WQ 1123 K/192 h/FC

51.0

24.3

s þ g þ a2 b0 þ b00 þ O þ a2

O þ b þ a2b

1023 K/708 h/WQ

26

1123 K/192 h/FC

54.5

21.8

1023 K/708 h/WQ a b

Minor content of phase. Only reflexes inherent in both a and a2 phases were identified.

Lattice parameters, pm

b0 þ b00 þ a2 þ Oa O þ a2b

b0 b00 a2

a

328.2 325.5 579.4 (XRD taken from section) s 457.2 g 402 a2 575 (XRD taken from section) b0 328.3 b00 325.5 O 607 a2 579.8 O 606.7 b 328.7 a2 581.3 b0 328.2 b00 326.2 a2 579.6 (XRD taken from section)

b

c

e e e

e e 467.3

e e e

552.8 406 474

e e 960 e 960.2 e e e e e

e e 466 469.5 465.9 e 468.0 e e 467.3

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Fig. 5. XRD diffractograms of samples Ti51.0Ta23.4Al25.6 (no. 25) annealed at 1023 K for 708 h followed by water-quenching (a), Ti50.1Ta11.9Al38.0 (no. 24) annealed at 1123 K for 192 h followed by furnace cooling (b) and Ti49.8Ta28.0Al22.2 (no. 23) annealed at 1123 K for 192 h followed by furnace cooling (c).

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Fig. 6. EBSD maps for sample no. 25 with composition Ti51.0Ta24.3Al24.7 after low-temperature annealing at 1123 K for 192 h: (a) SE-image of mapping area; (b) phase map with three distinct phases bo (yellow/white), a2 (red/black) and O (gray); (c) orientation map of bo-phase in normal direction; (d) orientation map a2-phase in normal direction and (e) orientation map of O-phase in normal direction. For purpose of illustration a typical, indexed Kikuchi pattern is given for each phase (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).

phase compositions along tie-lines and tie-triangles, the results from all above publications are in tolerable agreement. To our knowledge no isopleths were experimentally studied in the AleTaeTi ternary system. Thus, the above reviewed literature distinctly represents the evolution of solid-state equilibria in the temperature interval from 1273 to 1723 K. At the same time the experimental data on liquidsolid equilibria as well as information concerning stability of low-temperature ternary phases, isostructural with Ti2NbAl and Ti4NbAl3, are lacking and ought to be accomplished.

2.2. Complementary experiments 2.2.1. Preparation of alloys Aiming to investigate phase equilibria in the AleTaeTi system a number of 26 samples with different alloy composition were prepared by arc melting with a non-consumable tungsten electrode on a water-cooled copper hearth under purified argon. The samples, with a mass of 15 or 20 g each, solidified at cooling rates of about 100 K s1. The initial materials were bulk Al (99.99 wt.% Al), Ta (99.9 wt.% Ta) and Ti (99.9 wt.% Ti). The sample compositions are

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Fig. 7. EBSD maps for sample no. 25 with composition Ti50.1Ta11.9Al38.0 after low-temperature annealing at 1123 K for 192 h: (a) SE-image of the mapping area; (b) phase map with three distinct phases a2 (gray), g (red/black) and s (yellow/white); (c) orientation map of the a2-phase in normal direction; (d) orientation map of the g-phase in normal direction; (e) orientation map of the s-phase in normal direction. For purpose of illustration a typical, indexed Kikuchi pattern is given for each phase (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).

listed in Table 2 and pertain to three groups: (i) samples for the determination of invariant reactions and primary solidification phases were prepared on the basis of calculated liquidus and solidus surfaces using preliminary thermodynamic modelling (samples nos. 1e13); (ii) samples containing from 0 to 12 at.% Ta and from 45 to 47 at.% Al that form a þ g or a2 þ g lamellar structures (samples nos. 14e22); (iii) samples located around compositions where the formation of orthorhombic O and hexagonal B82 ternary phases were expected in analogy to the AleNbeTi system (samples nos. 23e26). As measured by reducing extraction in a Ni bath followed by chromatography, the oxygen content in the samples ranged from 0.02 to 0.04 wt.% (200e400 wt-ppm), and the contamination by N and H was lower than the detection limit (about 0.001 wt.% N and 0.003 wt.% H). The samples were chemically characterized and studied in the as-cast state and after annealing at selected temperatures (see Table 2) by DTA, XRD and SEM. Annealing was performed in a resistance furnace with a tungsten heater in argon

atmosphere gettered by Ti, Zr or Hf cuttings depending on an annealing temperature. The cooling rate after annealing ranged around 3 K s1. 2.2.2. SEM/EDS and XRD analysis of as-cast samples Scanning electron microscopy (SEM) and energy dispersive Xray analysis (EDS) were carried out in a SEM type Gemini 1550. For quantitative analysis of EDS spectra an own standard of Tie45.8 at% Ale7.8 at%Ta was used, cut from a sample with known chemical composition measured by wet chemical analysis. XRD measurements were carried out in a DRON-3 diffractometer on powder samples or metallographic sections for unbreakable alloys with an angular step of 2q ¼ 0.05 and an exposure of 8 s. Fig. 1 shows back scatter SEM micrographs of typical as-cast structures. The microstructure in the as-cast state of samples nos. 1, 2 and 4 (see Table 2) was similar to the one shown in Fig. 1(a) where the white large grains are primary s-phase. The samples differed only with respect to the phase fraction of s. In a sample no. 3 the

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Fig. 8. Precipitation of equilibrium phases a2 and/or a from massive gamma gm as function of the local temperature after annealing for 15.5 h in a Bridgman furnace and finally quenching. The annealing temperatures are (a) 1290 K, (b) at 1338 K, (c) at 1401 K, (d) at 1452 K, (e) at 1511 K and (f) at 1538 K. Table 6 Average composition of equilibrium phases in the alloy no. 17 (Ti46.5Ta7.7Al45.8) as function of annealing temperature and phase fractions measured by image analysis. T [K]

Composition of a2 or

a (at.%) Al

Ta

Composition of g (at.%)

Fraction of phase (area%)

Al

g

a2 or a

Ta

a2-g phase field 1290 1315 1338 1361 1373a 1381 1401 1419

38.1 38.3 38.5 38.9 38.3 39.4 39.7 40.4

       

0.7 0.8 0.7 0.8 0.5 0.5 0.3 0.6

9.1 9.0 9.3 9.4 8.7 9.2 9.5 9.6

       

0.4 0.5 0.3 0.3 0.2 0.3 0.2 0.2

47.3 47.2 46.8 47.1 47.6 47.0 47.7 47.3

       

0.4 0.6 0.5 0.4 0.3 0.4 0.6 0.3

7.5 7.5 7.6 7.8 6.7 7.7 7.5 7.8

       

0.8 0.2 0.3 0.4 0.3 0.3 0.4 0.3

e e 92.08 e 87.7  2.1 e 89.27 e

e e 7.92 e 12.3  2.1 e 10.73 e

0.3 0.3 0.2 0.3 0.3 0.2 0.2 0.2 0.2 0.2 0.3 0.3

9.6  9.2  9.1  9.1  9.0  8.8  8.7  8.5  8.5  8.5  8.4  8.3 

0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2

47.8 47.8 48.3 48.3 48.7 49.1 49.2 49.5 49.6 49.7 49.8 50.1

           

0.2 0.4 0.2 0.2 0.2 0.2 0.3 0.3 0.2 0.2 0.2 0.2

7.6 7.1 7.1 7.0 6.7 6.7 6.7 6.4 6.4 6.3 6.3 6.3

           

0.2 0.3 0.2 0.2 0.4 0.3 0.4 0.2 0.2 0.2 0.2 0.2

e 65.52 e 53.75 e e 46.64 e 32.39 e e 20.73

e 34.48 e 46.25 e e 53.36 e 67.61 e e 79.27

a-g phase field 1436 1452 1466 1479 1491 1502 1511 1519 1526 1531 1536 1538 a

40.7  42.0  42.3  42.9  43.1  43.6  43.7  44.1  44.2  44.4  44.7  44.9 

Reference experiment - isothermal annealing at 1373 K for 24 h.

primary solidification phase was the k-phase based on Ta39Al69, as shown in Fig. 1(b). The as-cast samples nos. 9 and 10 displayed primary a-dendrites with the morphology related to the hexagonal crystal structure of a, as shown in Fig. 1(c). A sample no. 11 solidified with primary g-phase dendrites as shown in Fig. 1(d). All other samples listed in Table 2 displayed the b-phase as primary solidification phase. Typical as-cast microstructures formed through primary b-solidification in vicinity L-a-b and L-b-s univariant lines are shown in Figs. 1(e) and 2 (left column): the most prominent characteristic of primary b-dendrites is the pronounced segregation, with Al being rich in interdendritic regions and Ta being rich in dendrite core (see Figs. 1(e), 2(a) and (c)). However, b-dendrites commonly decompose during cooling, following specific solid-state transformations (see Fig. 2(d) and (e)). Because the b-phase cannot be preserved untransformed down to room temperature in a number of alloys, XRD measurements cannot help identifying it. Microstructure analysis by SEM was therefore mandatory. This can easily be seen from Table 3 that comprises the XRD measurement results from as-cast samples nos. 1 to 13. XRD diffractograms of as-cast and annealed samples from the second group contain the reflections of the g-phase (gTiAl) and a2phase (Ti3Al) (Fig. 3). The lattice parameters were found to remain almost unaffected by the alloy composition, ranging from

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Table 7 Phase transformation temperatures measured in the AleTaeTi samples by DTA upon heating and cooling with a rate of 20 K min1. All samples were homogenized at subsolidus temperature. The calculated phase transformation temperatures are given in the last column. Sample no.

Composition

Phase transformation

Temperature, K Measured

1

Ti2.8Ta32.4Al64.8

2

Ti4.4Ta36.4Al59.2

3

Ti10.5Ta25.7Al63.8

4

Ti13.1Ta34,7Al52.2

5

Ti16.0Ta45.7Al38.4

6

Ti17.7Ta30.9Al51.4

7

Ti20.1Ta29.6Al50.3

8

Ti21.4Ta30.0Al48.6

9

Ti22.4Ta18.7Al58.9

10

Ti23.7Ta19.1Al57.2

11

Ti28.5Ta3.4Al68.1

3þ4þk43þk 3þk4Lþ3þk Lþ3þk4Lþk Lþk4Lþsþk Lþsþk4Lþs Lþs4L 4þs43þ4þs 3þ4þs43þs 3þs43þsþk 3þsþk4sþk sþk4Lþsþk Lþsþk4Lþs Lþs4L gþ3þ44gþ3þs gþ3þs4gþ3þk gþ3þk4Lþ3þk Lþ3þk4Lþk Lþk4L gþ3þ44gþ3þs gþ3þs4gþsþk gþsþk4aþsþk aþsþk4Lþaþs Lþaþs4Lþs Lþs4L s4bþs bþs4Lþbþs Lþbþs4Lþb Lþb4L gþ 4 4 g þ 4 þ s gþ4þs4gþs gþs4aþgþs aþgþs4aþsþk aþsþk4Lþaþs Lþaþs4Lþbþs Lþbþs4Lþb Lþb4L gþ44g þ4þs gþ4þs4gþs gþs4aþgþs aþgþs4aþs aþs4Lþaþs Lþaþs4Lþbþs Lþbþs4Lþb Lþb4L gþs4aþgþs aþgþs4aþs aþs4aþbþs aþbþs4Lþbþs Lþbþs4Lþb Lþb4L g4Lþg Lþg4Lþaþg Lþaþg4Lþa Lþa4L gþ4þs4gþs gþs4g g4Lþg Lþg4Lþa Lþa4Lþb Lþb4L 3þhþz4gþz gþz4z z4Lþz Lþz4Lþg Lþg4L

Calculated

Heating

Cooling

1565 1805 e 1863 e 1889 1320 e e 1687 1803 1847 1954 1302 1626 1714 1761 1814 1343 1616 1675 1743 1798 1935 1653 1937 2105 2220a 1133 1210 1667 1736 1763 1808 1878 1917 914 1045 1658 1755 1784 1807 1855 1967 1639 e 1727 1800 1837 2023 1733 1742 e 1826 1347 1586 1753 e 1787 1818 1259 1404 1705 e 1742

e 1772 e 1824 e 1875 e e e 1648 1724 1807 1879 e 1555 1594 1761 1796 e 1607 1649 1730 1809 1931 1552 1820 2002 e e e 1599 1639 1685 e 1805 1873 e e e 1641 1719 e 1785 1923 1610 e 1710 e 1810 1905 1687 1742 e 1807 e e 1640 1740 e 1798 e e 1663 e 1714

1540 1800 1803 1866 1867 1883 1471 1534 1555 1612 1817 1849 1956 1327 1623 1746 1770 1801 1327 1623 1728 1763 1774 1942 1569 1945 2070 2241 1090 1237 1674 1728 1763 1770 1870 1945 914 1105 1621 1685 1761 1770 1831 1984 1582 1632 1700 1770 1828 2030 1754 1769 1770 1784 1285 1615 1766 1770 1796 1802 1096 1428 1704 1710 1722

V.T. Witusiewicz et al. / Intermetallics 19 (2011) 234e259

245

Table 7 (continued) Sample no.

Composition

Phase transformation

Temperature, K Measured

12

Ti28,2Ta19.1Al52.6

13

Ti32.8Ta32.1Al35.1

14

Ti41.3Ta11.8Al46.9

15

Ti43.4Ta11.8Al44.8

16

Ti44.9Ta7.7Al47.4

17

Ti46.5Ta7.7Al45.8

18

Ti46.9Ta7.7Al45.4

19

Ti48.7Ta3.9Al47.4

20

Ti50.8Ta3.9Al45.3

gþs4aþgþs aþgþs4aþg aþg4a a4Lþa Lþa4Lþb Lþb4L sþs4gþsþs gþsþs4gþs gþs4bþgþs bþgþs4bþs bþs4b b4Lþb Lþb4L a2 þ g þ s 4 b þ g þ s bþgþs4bþg b þ g 4 bo þ g bo þ g 4 bo þ a þ g bo þ a þ g 4 a þ g aþg4a a4aþb aþb4Lþaþb Lþb4L g þ s þ s 4 a2 þ g þ s a2 þ g þ s 4 a2 þ b o þ g a2 þ b o þ g 4 b o þ g bo þ g 4 a þ bo þ g a þ bo þ g 4 a þ g aþg4a a 4 a þ bo a þ bo 4 a þ b aþb4b b4Lþb Lþb4L g4bþg bþg4aþg aþg4a a4aþb aþb4Lþaþb Lþaþb4Lþb Lþb4L g þ s 4 a2 þ g þ s a2 þ g þ s 4 a2 þ g a2 þ g 4 a2 þ bo þ g a2 þ bo þ g 4 a þ bo þ g a þ bo þ g 4 a þ g aþg4a a4aþb aþb4Lþaþb aþbþL4Lþb Lþb4L a2 þ g 4 a þ bo þ g a þ bo þ g 4 a þ g aþg4a a4aþb aþb4Lþb Lþb4L a2 þ g 4 g g4aþg aþg4a a4Lþa Lþa4Lþb Lþb4L a2 þ g 4 a þ g aþg4a ? (metastable transition)

a4aþb aþb4Lþb Lþb4L

Calculated

Heating

Cooling

1602 1617 1713 1798 1802 1911 e 1093 e 1419 1670 2103, 2068a e 1316 e 1493 1537 e 1599 1724 1805 1940 1129 e 1365 1492 e 1569 1614 1674 1764 1836 1952 1412 e 1615 1743 1783 e 1879 1125

1521 1581 1600 1778 1813 1905 e e e 1360 1478 2124 e e e e e e 1516 1677 1801 1914 e 1278 e 1413 e e 1568 e 1682 1763 1933 1348 e 1610 1698 e 1794 1870 e

1384 e e 1583 1717 1785  6a 1813 1901 1415 e 1576 1698 1810 1901 1392 e 1620 1763 1810 1850 1400 1582 1686 1757 1799 1871

e e e 1571 1692 e 1781 1887 1369 e 1487 1674 1782 1892 e 1397 1610 e 1779 1833 1363 1483 1681 1741 1760 1860

1556 1571 1714 1798 1804 1917 1064 1084 1365 1374 1590 2136 2282 1321 1327 1502 1534 1543 1596 1649 1812 1944 1132 1321 1340 1509 1526 1527 1530 1682 1779 1829 1967 1410 1411 1622 1754 1802 1803 1880 1127 1129 1388 1391 1398 1597 1720 1801 1802 1889 1391 1392 1563 1674 1800 1901 1283 1441 1631 1785 1789 1825 1404 1568 e 1731 1786 1848 (continued on next page)

246

V.T. Witusiewicz et al. / Intermetallics 19 (2011) 234e259

Table 7 (continued) Sample no.

Composition

Phase transformation

Temperature, K Measured

21

Ti55.6Al44.4

22

Ti53.2Al46.8

23

Ti49.8Ta28.0Al22.2

a2 þ g 4 a þ g aþg4a a4aþb aþb4b b4Lþb Lþb4L a2 þ g 4 a þ g aþg4g aþg4a a4Lþa Lþa4Lþb Lþb4L b þ O 4 a2 þ b þ O a2 þ b þ O 4 a2 þ b a2 þ b 4 b ? (metastable transition)

24

Ti50.1Ta11.9Al38.0

25

Ti51.0Ta24.3Al24.7

26

Ti54.5Ta21.8Al23.7

bo 4 b b4Lþb Lþb4L g þ s 4 a2 þ g þ s a2 þ g þ s 4 a2 þ s a2 þ s 4 a2 a2 4 a2 þ b o a2 þ b o 4 a2 þ b o þ g a2 þ b o þ g 4 a þ a2 þ b o a þ a2 þ b o 4 a þ b o a þ bo 4 bo bo 4 b b4Lþb Lþb4L bo þ s þ O 4 a2 þ bo þ s a2 þ b o þ s 4 a2 þ b o a2 þ bo 4 bo bo 4 b b4Lþb Lþb4L bo þ O 4 a2 þ bo þ O a2 þ b o þ O 4 a2 þ b o ? (metastable transition) a2 þ bo 4 bo bo 4 b b4Lþb Lþb4L

a b c

Calculated

Heating

Cooling

1379 1602 1678 1747 1770 1834 1231 1466 1644 1757 e 1813 1159b, 1123c e 1175, 1209b, 1149c 1282, 1293b, 1181c 1410, 1437c 2193 2327 1030, 1008b 1124, 1171b, 1165c 1170, 1219b, 1215c 1343, 1343b, 1342c e 1413, 1445b, 1408c e e 1793 1939 2048 1117b, 1126c 1170b, 1192b 1258b, 1302c 1515, 1518b, 1522c 2163 2277 1126, 1124b, 1130c 1164, 1124b, 1142c 1215, 1250b 1331, 1340b, 1274c 1536 2124 2237

1323 1483 1669 1733 1770 1813 e 1357 1608 1697 1755 1793 e e 1209b, 1180c 1411c 2179 2321 e 1011c 1034c 1337b, 1282c e e e 1452b, 1456b e 1930 2034 1012c e 1167b, 1163c 1505, 1477c 2136 2267 1110b e e 1340b e 2114 2229

1392 1566 1708 1759 1765 1817 1298 e 1628 1761 1762 1792 1121 1122 1168 e 1409 2199 2330 969 1128 1148 1336 1367 1391 1392 1424 1799 1930 2026 1120 1146 1229 1505 2148 2274 1122 1123 e 1246 1477 2116 2242

Temperature of incipient melting measured by the PiranieAlterthum method. Data for the sample annealed at 1123 K for 192 h. Data for the sample annealed at 1023 K for 708 h.

a ¼ 400 O 403 pm and c ¼ 407 O 409 pm for the g-phase; a ¼ 575 O 577 pm and c ¼ 462.0 O 463.5 pm for the a2-phase. In samples nos. 14 and 15, containing 12 at.% Ta, a small fraction of the bcc b-phase (presumably it is bo of the CsCl crystal structure type) was identified and the lattice parameter was determined to be a ¼ 322 O 323 pm. 2.2.3. SEM/EDS, SEM/EBSD and XRD analysis of annealed samples All prepared samples were annealed at sub-solidus temperature and cooled to room temperature at an approximate rate of 3 K s1. As example, few selected SEM/BSE micrographs of microstructures of the heat-treated samples are presented in right column of Fig. 2. The constituent phases were also identified by XRD as listed in Table 3. 14 samples were taken for microstructure analysis and phase compositions measurements in the annealed condition, where the above given cooling rate was sufficient to preserve the phase composition equilibrated at annealing temperature. Microstructure constituents and EDS results on phase composition are summarized in Table 4. Several annealed samples were analyzed in more detail and gave valuable information, as described below.

The as-cast samples nos. 4 and 6 were annealed at 1773  10 K in the solideliquid state. This allowed measuring the composition of the eutectic phase constituent in equilibrium with the s-phase (see for instance Fig. 2(f)) and determining the composition of the individual phases inside the eutectic by EDS. In this way, the tie triangle pertinent to the univariant equilibria L þ a þ s at 1773  10 K was experimentally accessible. The composition of the eutectic (liquid at this annealing temperature) was Tie57.0Ale18.1Ta (at.%). From the last samples the L þ s and L þ a tie-lines were also determined (see Table 4). Slices cut from samples nos. 23 through 26 after sub-solidus homogenizing were subjected to low-temperature annealing at 1123 K for 192 h and at 1023 K for 708 h, aiming to investigate the formation of low-temperature ternary phases O and B82. For the annealing treatments at 1023 K the samples and additional Ti chips were sealed in quartz tubes and filled with argon after repeated evacuation cycles. On annealing the samples were water quenched. The 1123 K annealing was performed in a chamber of resistance furnace under Ar and samples were inserted in Ti cuttings. Fig. 4(a) and (b) display the microstructure observed in samples no. 25 (Ti51.0Ta23.4Al24.7) after above annealing treatments. Fine, plate-like

V.T. Witusiewicz et al. / Intermetallics 19 (2011) 234e259

6

1743

1412

4 no.16 1343

2

1615 1510 16151675

no.4

1890 1743 1742 1798

0 -2 no.9 1259

-4

Δ T, K

1879

1733 1705

1408

1826

no.14

-6

1339

-8 -10

no.14:

Al68.1Ta3.4Ti28.5

-12

no.9: Al58.9Ta18.7Ti22.4

-14

no.16:

no.4:

1000

1200

Al52.2Ta34.7Ti13.1 Al47.4Ta7.7Ti44.9 1400

1742

1600

1800

Temperature (K) Fig. 9. DTA curves recorded upon heating with a rate of 20 K min1 for selected AleTaeTi samples after sub-solidus annealing. Table 7 contains further information regarding the phase transformations behind individual endothermal peaks. For clarity, the curves have been shifted vertically.

precipitates were found throughout the samples, the precipitates being much finer after lower temperature annealing, regardless of the fact that the annealing time was longer. The similar microstructures were distinctive also for the alloys nos. 23 and 26. Fig. 4 (c) and (d) show the microstructure of samples no. 24 (Ti50.1Ta11.9Al38.0) annealed at 1123 and 1023 K. Different precipitates with distinct shape and size were observed.

a 1873 K / 7 h

4

1123 K / 192 h

-1

Tnorm. (K g )

1023 K / 708 h

0 Alloy no. 24: Ti50.1Ta11.9Al38.0

b -4

1973 K / 7 h 1123 K / 192 h 1023 K / 708 h

-8 Alloy no. 25: Ti51.0Ta24.3Al24.7

800

1000

1200 Temperature (K)

1400

1600

Fig. 10. Comparison of DTA curves obtained upon heating with rate 20 K min1 from homogenized at sub-solidus temperatures, annealed at 1123 K for 192 h and 1023 K for 708 h samples: (a) alloy no. 24 and (b) alloy no. 25.

247

XRD (Table 5) and EBSD examinations (see Table 2) were performed to identify phases and to reveal their crystallographic nature. EDS measurements of local phase composition were performed, wherever the size of precipitates was large enough, e.g. at triple junctions of matrix grains and grain boundaries (see Table 4). Fig. 5 displays the obtained XRD diffractograms of three samples, no. 25 (Ti51.0Ta23.4Al25.6) annealed at 1023 K for 708 h followed by water-quenching (WQ), no. 24 (Ti50.1Ta11.9Al38.0) and no. 23 (Ti49.8Ta28.0Al22.2) both annealed at 1123 K for 192 h followed by furnace cooling (FC). The first sample contains the O-phase as being evidenced by EBSD mapping (Fig.6). Its experimental XRD diffractogram was compared with calculated spectra performed by using site occupancies of both O1 and O2 modifications corrected for the result of EDX measurements of composition as Ti54.2Ta19.3Al26.5 (see Table 4) that corresponds to Ti2.17Ta0.77Al1.06. As shown in experimental [35] and theoretical [36] works, the high-temperature O1 and low-temperature O2 modifications in the AleNbeTi system differ by site occupations in (4c2) and (8g) Wyckoff positions, while the (4c1) positions are occupied only with Al atoms in both cases. The O1-phase is characterized by a random occupancy of Ti and Nb atoms in the both sites. In the O2-phase Nb atoms tend to occupy (4c2) and Ti atoms to be in (8g). The difference between the fractions of Ti atoms in these positions was found to reach about 0.5. Thus, the site occupations in Ti2.17Ta0.77Al1.06 were estimated based on its composition as follows: 2.89Ti þ 1.03Nb þ 0.08Al in (4c2) and 5.78Ti þ 2.06Nb þ 0.16Al in (8g) for O1 and 1.56Ti þ 2.36Nb þ 0.08Al in (4c2) and 7.11Ti þ 0.73Nb þ 0.16Al in (8g) for O2. The experimental XRD diffractogram of Ti2.17Ta0.77Al1.06 matches well with the calculations (Fig. 5(a)) for both O1 and O2. Nevertheless, the O1 calculated XRD pattern provides better R-factors (Rp ¼ 6.8 and Rwp ¼ 9.0 against Rp ¼ 9.2 and Rwp ¼ 12.3). The authors of Ref. [37] reported that in the s-phase of the AleNbeTi system according to single crystal XRD a mixture of atoms 1.00Ti þ 0.75Nb þ 0.25Al occupies (2a) Wyckoff positions, while in (2c) and (2d) there are only one kind of atom, Al and Ti, respectively. Taking into account the crystal structure type, it should be Ti3Nb1xAl2þx with x ¼ 0.25 that exactly corresponds to conventional Ti4NbAl3 stoichiometry. This composition somewhat contradicts with EDX measurements of Ref. [37] (Ti52.5Nb13Al34.5) and TEM X-ray microanalysis of Ref. [38] (Ti51.7Nb13.1Al35.2). Drawing an analogy, we assume that the s-phase (i.e. Ti3.00Ta0.89Al2.11) includes the mixture of atoms 1.00Ti þ 0.89Ta þ 0.11Al in (2a) positions to meet its composition. As seen in Fig. 5(b), the calculated XRD diffractogram of Ti3Ta1xAl2þx at x ¼ 0.11 in the sample Ti50.1Ta11.9Al38.0 (no. 24) is well consistent with the experimental one. It should be noted that our attempts to describe this XRD data using crystallographic parameters of a trigonal C6 incomplete u-phase (CdI2 crystal structure type) [37,39] failed, since the characteristic for the phase (111) peak at 2Q z 43 is absent in the experimental diffractogram (see Fig. 5(b)). EBSD maps from the sample Ti51.0Ta23.4Al25.6 (no. 25) annealed at 1123 K for 192 h are shown in Fig. 6. The matrix consists of large grains of the transformed b-phase and the precipitates mainly consist of O-phase with a small fraction of a2-platelets inside the core of O and at the grain boundary. This indicates that the O-phase forms through a peritectoid reaction b þ a2 / O, similar to the AleNbeTi system [40]. It is worth looking at Fig. 4(a) to observe the envelope-like layer of the light gray O-phase around darker gray a2precipitates. EBSD maps from the sample Ti50.1Ta11.9Al38.0 (no. 24) annealed at 1123 K for 192 h are shown in Fig. 7. Here the matrix phase consists of small grains of the s-phase with B82 structure, having strict crystallographic relations with the parent b-grains, as it was previously shown for the AleNbeTi system [37,38,41,42]. The precipitates are mainly g-phase and minor

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Fig. 11. Calculated with the present description site fractions of Al, Ti and Ta in the L10 g-phase at 1173 K as function of Ta dissolved in this phase.

amounts of a2. The precipitate orientations indicate that at the given annealing temperature the initial b-phase (in fact b0) first precipitated g and a2 and then transformed almost completely to the s-phase. As seen in Table 5, the O-, a2- and b-phase were identified by XRD in samples no. 23 (Ti49.8Ta28.0Al22.2) and 25 (Ti51.0Ta23.4Al25.6) annealed at 1123 and 1023 K, as well as in sample no. 26 (Ti54.5Ta21.8Al23.7) annealed at 1123 K. It is worth noting that after annealing at 1123 K followed by furnace cooling, the XRD diffractograms of samples nos. 23, 25 and 26 contain split peaks of the b-phase, which were described as b0 and b00 having different lattice parameters (Fig. 5(c)). They do not correspond to the ordered b0 since the characteristic peak (111) of the CsCl crystal structure types was not observed at 2Q z 48.2 (b0 ) and at 2Q z 48.6 (b00 ). It seems to be well consistent with a miscibility gap in the b(Ti,Ta,Al) solid solution, inevitably resulted from thermodynamic modelling. The same samples annealed at 1023 K follow by quenching in water did not show two b-phases. Apparently, the demixing occurred at temperatures below 1023 K during furnace cooling of the samples. 2.2.4. Annealing experiments starting from massive gamma in alloy Ti46.5Ta7.7Al45.8 A special annealing experiment was performed for alloy Ti46.5Ta7.7Al45.8, aiming to investigate the evolution of equilibrium phases, starting from the metastable massive gamma (gm) [3]. The main focus was to measure the phase fractions and the chemical composition of the equilibrium phases a or a2 and g grown from gm during holding periods at different annealing temperatures. The experiment was performed in a Bridgman furnace [43e45], which offers well-defined conditions to achieve and maintain a stationary temperature gradient along the longitudinal axis of the sample and allows freezing-in the resulting microstructure by a final quenching operation. A cylindrical sample, ⌀ 7 mm and length of 121 mm, from alloy Ti46.5Ta7.7Al45.8 (sample no. 17 in Table 2), previously transformed to the massive gm state, was subjected to annealing in a defined temperature gradient that covered the temperature interval of interest from 1273 to 1573 K. The axial temperature field was measured in an identical, but separate experiment using thermocouples placed inside the cylindrical sample. This allowed assigning the local temperature to each position along the sample axis, with a measurement error of 10 K. The Bridgman annealing experiment itself encompassed heating, holding for a duration of 15.5 h, followed by quenching at

cooling rates >50 K s1. In the region of interest that spans from about 1283 to 1553 K, the sample was analyzed in an SEM Gemini 1550 within local observation areas, each corresponding to a known annealing temperature. These areas were targeted successively by moving the sample with a step of 2 mm underneath the electron beam. Within each observation area the phase composition of the equilibrium phases a or a2 and g grown from gm were measured by EDS: at least 10 individual spectra were acquired per phase, mainly from grain boundaries, where precipitates were larger than 1 mm. Fig. 8 displays a series of micrographs that illustrate the precipitation of a or a2 from the massive gm as function of the local annealing temperature. One can distinguish (i) precipitation at the former boundaries of massive gamma grains and (ii) intragranular precipitation of a and/or a2 laths on the (111) planes of g. Both, phase fractions and the size of a2 and/or a increase with increasing annealing temperature. Table 6 summarizes the average composition of equilibrium phases and the volume fraction of the coexisting phases. The latter was measured by image analysis. The measured phase compositions were used as input data for thermodynamic modelling, along with all other available data on phase equilibria in the AleTaeTi system. 2.2.5. DTA analysis and incipient melting of the alloys Temperatures of solid-state and solideliquid phase transformations were determined by DTA using Sc2O3 and Al2O3 crucibles and W/W-20Re string thermocouples designed by Kocherzhinskiy et al. [46,47]. The DTA measurements were performed under high purity He with heating and cooling rate of 20 K min1. The thermocouple was calibrated using the IPTS-90 reference points of Al, Au, Pd, Pt, Rh and additional reference points of Fe and Al2O3. DTA measurements were performed first for the ascast samples. This allowed to select an appropriate temperature for sub-solidus annealing, in order to homogenize segregation inherited from casting. The DTA results for the homogenized samples are summarized in Table 7. During the evaluation of phase transformation temperatures from the recorded DTA curves the practical guide on differential thermal analysis of metals and alloys recommended NIST [48] was followed. The temperatures of incipient melting were also determined for samples nos. 2, 3, 7 to 11 using the method devised by Pirani and Alterthum [49]. Bar-shaped specimens, which were clamped between two water-cooled copper electrodes through tungsten inserts (tips of electrodes and plate washers), were resistively

V.T. Witusiewicz et al. / Intermetallics 19 (2011) 234e259

249

Fig. 12. Reaction scheme down to room temperature (RT) for the AleTaeTi system calculated using the present thermodynamic description.

heated under high purity Ar (slightly higher than the atmospheric pressure). The temperature was measured on the background of a black body hole (the diameter to depth ratios were about 1 to 4) with a disappearing filament-type pyrometer “EOP-68” of standard

quality level. Its maximum instrumental errors amount to 2.8 K for the temperature region 1170e1670 K and 4 K for 1670e2270 K. It was calibrated and certificated by the National Scientific Centre “Institute of Metrology” (http://www.metrology.kharkov.ua/eng/

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Fig. 12. (continued).

centres/NC-1.htm). The temperature of incipient melting was read in the moment of formation of a spot on the background of the black body hole. This moment also corresponds to a change of the sample shape at the bottom (which becomes half-spherical) and (or) to a change of the reflection at the sample surface [50]. The latter turns smooth due to filling the surface roughness by melt. A number of 5 measurements were carried out for each sample and afterwards the average value as well as the standard deviation was determined.

Table 7 contains the phase transformation temperatures evaluated from DTA measurements and by the PiranieAlterthum method in all investigated samples, together with the values calculated with the proposed thermodynamic description of the AleTaeTi system. For illustration, Fig. 9 depicts heating curves obtained by DTA for selected samples. As can be seen from these examples the majority of the alloys undergo several solid-state transformations before melting.

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251

Table 8 Invariant equilibria in the AleTaeTi system involving the liquid phase. Invariant equilibrium 41 þ 42 þ 43 þ 44

Type

T, K

Composition of phases, at.%

41 L L L L L L L L L L

þb4a þa4g þb4aþs þs4aþk þa4gþk 4gþ3 4gþ3þk þg4z þg43þz þ 3 þ 3(l) 4 (aAl)

p4 p5 U1 U2 U3 e1 E1 p8 U5 P

1815 1775 1770 1763 1755 1750 1746 1712 1704 938.7

42

43

44

Al

Ta

Al

Ta

Al

Ta

Al

Ta

54.7 61.0 57.6 58.5 59.8 64.3 62.2 67.8 71.8 99.9

11.0 13.4 24.6 24.3 23.2 17.7 21.8 1.8 4.1 1  104

46.2 55.6 44.2 44.9 53.9 60.2 58.8 62.8 64.3 76.6

16.2 15.4 29.1 49.3 22.7 18.3 20.0 3.4 6.9 6  106

51.5 57.6 52.3 53.0 57.1 71.3 70.8 66.0 73.4 75.7

13.1 14.2 24.0 23.5 20.9 17.7 19.9 2.5 6.7 3  106

e e 44.3 56.4 57.7 e 59.8 e 68.4 99.2

e e 49.4 34.0 33.6 e 33.2 e 4.5 9  104

Fig. 13. Calculated liquidus projection in the AleTaeTi system compared to experimental data on primary solidification phases (a) and liquidus projection with isotherms (b).

Fig. 14. Calculated solidus projection in the AleTaeTi system with experimental data obtained in the present work (a) and with isotherms (b).

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elements are adopted from the SGTE database compiled by Dinsdale [51]. The designation of individual phases, their crystal structure and the models employed in the present thermodynamic description are summarized in Table 1. The expressions for the molar Gibbs energy as function of temperature and composition for these models are given in Refs. [52,53]. The models were kept simple, involving two or three sublattices only. For some phases, e.g. s and g, more sublattices may be requested to correctly represent lattice occupation and to distinguish between specific configurational contributions to their Gibbs free energy [54,55]. Since in this case the number of model parameters increases drastically, preference was given to the simpler models. Literature provides experimental data [56,57] and first-principle calculations [58e60] regarding the site occupation of Ta in the L10 g-phase. Above room temperature Ta preferably partitions to the Ti sublattice, independent of the g-phase composition. Fig. 11 shows the calculated site fractions of Al, Ti and Ta in two sublattices of the g-phase at 1173 K as function of dissolved Ta in this phase. The results calculated with the present description are fully consistent with the experimental and ab initio findings [56e60]. The composition of the B82 s-phase in AleTaeTi was found to have a narrow solubility range close to Ti3Ta1xAl2þx at x ¼ 0.11 (see Section 2.2.3). Taking into account above information on the occupancy of Wyckoff positions this phase was modelled using a three sublattice model as [(Ti)3:(Al,Ta,Ti)1:(Al)2]. For the O-phase we used two sublattice model analogous to that for the O1-phase in the AleNbeTi system [12]. The model parameters of the phases listed in Table 1 were evaluated by searching for the best fit to available experimental data on phase equilibria using the PARROT optimiser of the software Thermo-Calc [52]. During the optimization process the experimental data obtained in the present work, the data on tielines and tie-triangles at 1373 K by Ref. [22] and at 1473 K by Ref. [25] as well as the data on the L þ a þ b univariant reaction by Ref. [33] were accorded a weight of 1.5. A weight of 2 was assigned to solid phase transformation data obtained by DSC [61] and data on tie-lines and tie-triangles at 1273, 1473, and 1573 K measured by Kainuma et al. [32]. All other data used were assigned a weight equal 1. The gas phase was included in order to allow extrapolation up to temperatures of 6000 K. The complete thermodynamic database in Thermo-Calc format [52] is included in the Appendix. Fig. 15. The b-transus surface in the AleTaeTi system projected on the Gibbs triangle (a) and a partial b-transus surface in the region of the misibility gap of the b-phase (b). Within the shadowed area in (a) the b-phase is ordered, i.e. bo.

4. Results of modelling and discussion

Fig. 10 shows the DTA curves obtained for samples nos. 24 and 25 before and after low-temperature annealing at 1023 K for 708 h and at 1123 K for 192 h. These samples were shown in Section 2.2.3 to contain O-phase (no. 25) and s-phase (no. 24) formed during the low-temperature annealing treatments. The DTA peaks associated to O-phase or s-phase transformations in heating are clearly visible in the DTA curves after low-temperature annealing. Without this additional annealing they are sluggish and hardly detectible. The corresponding onset temperatures are 1121  5 K and 1168  5 K, respectively.

The reaction scheme was calculated with the present thermodynamic description down to room temperature. It is presented in Fig. 12, including the binary constituent systems AleTa and AleTi reported earlier in Refs. [5,6]. According to the present description, the ternary system AleTaeTi contains six invariant four-phase equilibria involving the liquid phase and four quasi-binary three-phase invariant equilibria with the liquid phase (e1, p4, p5 and p8). The compositions of coexisting phases involved in invariant equilibria with the liquid phase are summarized in Table 8. A number of 23 fourphase invariant reactions and 4 quasi-binary invariant reactions take place in solid state down to room temperature. They are given in the reaction scheme. The calculations yielded the miscibility gap in the b(Ti,Ta,Al) solid solution for compositions close to the Ta-rich corner. This agrees well to XRD results (see Section 2.2.3). The b-phase demixing originates from two critical points marked as C1 (1461 K) and C2 (813 K) and extends down to

3. Thermodynamic models and optimization procedure In the present work we used the thermodynamic models for the individual phases stemming from descriptions of the constituent binary systems AleTa [5], AleTi [6] and TieTa [7]. The Gibbs energy descriptions for the stable and metastable structures of the pure

4.1. Reaction sequence

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Fig. 16. Isothermal sections at elevated temperatures: (a) 1770 K, (b) 1723 K, (c) 1623 K, (d) 1573 K and (e) 1473 K. The lines are calculated equilibria using the proposed thermodynamic description and symbols are experimental data. The dotted lines show the A2/B2 orderedisorder transformation of the b-phase.

the eutectoid reaction E3 at 787 K. A similar demixing of the bphase of the ternary AleNbeTi system resulted from recent calculations with Thermo-Calc version S using the thermodynamic description of Ref. [12], which was found for a temperature interval from 866 K down to 725 K. Since previous versions of Thermo-Calc software were not capable to detect such features

properly, it was omitted in reaction scheme and discussion. A number of other alloys systems with Ti or Ta (TieW, TieMo, TieSc, TaeHf, TaeZr, etc. [8]) display b-demixing. Also some ternary systems with A2/B2 ordering, for instance AleFeeNi [62], show it. Future work will hopefully shed more light on this specific behaviour also for AleTaeTi alloys.

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Fig. 17. Isothermal sections at moderate and ambient temperatures showing demixing in the b-phase and the occurrence and evolution of equilibria with ternary O- and s-phases at (a) 1373 K, (b) 1323 K, (c) 1273 K, (d) 1123 K, (e) 1023 K and (f) 298 K. Lines are calculated equilibria using the proposed thermodynamic description and symbols are experimental data. The dotted lines show the A2/B2 orderedisorder transformation of the b-phase.

4.2. Liquidus, solidus and b-transus surfaces The calculated liquidus surface is displayed in Fig. 13 as projection on the Gibbs triangle along with available experimental data on primary phase solidification. Compared to the AleNbeTi system [12], the primary a and b fields are more extended towards

the Al-rich corner, while the field of primary g is less extended. The primary s-phase field is wider than in the liquidus surface projection tentatively proposed by Ref. [15] due to the incorporation of the new binary AleTa description [5]. The present thermodynamic description can reproduced the boundary between a and b primary solidification fields evaluated by Johnson et al. [33].

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255

Fig. 18. Selected isopleths in the AleTaeTi system showing calculated equilibria (lines) compared to experimental data on phase transformation temperatures. The dotted lines show the A2/B2 orderedisorder transformation of the b-phase.

The data on primary phase fields reported in Refs. [25,28,29,34] also agree well with the calculated liquidus surface. The liquidus surface is characterized by four U-type (Übergangsreaktion) reactions marked as U1, U2, U3 and U5, one eutectic reaction marked as E1 and one peritectic reaction marked as P (the latter being in the Al-corner). None of these reactions was distinctly known before. The reactions given in the tentative scheme of Ref. [15] differ with regard to the type and/or the phases involved. The L-a-b, L-a-g and L-g-z univariant reaction lines display maximum temperatures at points p4, p5 and p8, which corresponds to the three-phase quasi-binary reactions Lþb4a, L þ a 4 g and L þ g 4 z of peritectic nature, respectively. Like in the AleNbeTi system [12], the univariant reaction line L-g-3 displays one maximum point e1 corresponding to the quasi-binary eutectic reaction L 4 g þ 3. The calculated solidus surface is shown in Fig. 14 in comparison to experimental data and with calculated isotherms. The solidus

projection includes 6 tie-triangles corresponding to the four-phase invariant reactions with liquid phase described above. Four maximum folds on the solidus projection are marked by thick solid lines. They correspond to the four quasi-binary reactions with the liquid phase. The solidus surface shows that seven phases have wide homogeneity ranges, with the largest range being that of b, and only for 3(l) and (aAl) they are narrow. The calculated b-transus surface is shown in Fig. 15(a). The invariant reactions are labeled as in the reaction scheme in Fig. 12. The calculated isotherms are in the range from 800 K to 2347 K (temperature of the L þ b 4 s reaction, p1). Along the line p1-U1p4-p6 the b-transus surface connects to the solidus surface. Within the shadowed area the b-phase is ordered, i.e. bo. The field E2-U8U10 corresponds to a narrow range where precipitation of g occurs from bo. From the figure one can clearly distinguish the large composition ranges related to the precipitation of a-, a2-, O- or sphases. The calculated miscibility gap of the ternary b(Ti,Ta,Al) solid

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Fig. 19. Investigated isopleths that pass through or near to the alloy composition Tie46Tie8Ta. Lines are calculated equilibria and symbols are experimental data on phase transformation temperatures obtained by DTA measurements and reported by Refs. [22,61]. The dotted lines show the A2/B2 orderedisorder transformation of the b-phase. The thin dotted lines in the first isopleth show shift of phase boundaries when the contribution of B2 ordering is eliminated.

solution (marked as b0 þ b00 or bacehf þ babdgf) is shown in some detailed in Fig. 15(b). The binodal surface splits the fields of a2- and O-phase precipitation into two domains and gives rise to a cascade of four-phase reactions U16, U17 and E3 on both b-d-g and c-e-h margins.

4.3. Isothermal sections Figs. 16 and 17 display a series of isothermal sections calculated with the database given in Appendix. Experimentally measured tielines and tie-triangles from literature [22,23,25,32] as well as those determined within the present work are included. On the whole, the agreement between calculated and experimentally measured phase equilibria is fairly good, as discussed below:

For T ¼ 1770 K (Fig. 16(a)) only few experimental data are available from the sub-solidus annealing experiments at 1773  10 K summarized in Table 4. For T ¼ 1723 K (Fig. 16(b)) the tie-lines measured by different authors for nominally the same alloy composition and temperature sometimes differ by more than 5e10 at.%. The calculated isotherm agrees fairly well with experimental data of the present work and from Ref. [23]. Authors of Ref. [25] for instance, described the microstructure of alloy Tie50Ale25Ta (at.%) annealed at 1723 K that consists of a þ b rather than b þ s as it was reported previously [24]. During optimization of the thermodynamic description it turned out that this experimental information cannot be reproduced without negative impact on the extent of the a and s primary phase domains and phase transformation temperatures measured by DTA. Our calculations show that the above-mentioned alloy is in

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257

Fig. 20. Microstructure of samples after annealing for 10 h at 1703 K followed by cooling with a rate of 3 K s1: (a) Ti50.8Ta3.9Al45.3, (b) Ti43.8Ta11.8Al44.8, (c) Ti48.7Ta3.9Al47.4 and (d) Ti41.3Ta11.8Al46.9. The b-phase appearing white in BSE-contrast is present in (a) and (b) but absent in (c) and (d).

a three-phase field a þ b þ s in the temperature interval from 1760 to 1690 K, containing few mol.% s-phase. The isothermal sections for 1623, 1573, 1473 and 1373 K in Figs. 16(c)e(e) and 17(a) show that the experimental data of Refs. [22,25] and especially the present and most recent data by Kainuma et al. [32] agree well to the calculated phase equilibria. The calculations also confirm experimental observations from Ref. [22] regarding the stability of the 4-phase (T < 1713 K) and of the kphase (T  1473 K). The former displays only minor solubility for Ti at 1723 K, but the 4 solid solution steadily extends into the ternary as the temperature decreases. The minor solubility of Ti in the kphase is observed at 1473 K and its homogeneity range extends with increasing temperature. At T ¼ 1373 K (Fig. 17(a)) the miscibility gap of the Ta-rich bphase (b / b0 þ b00 ) discussed in Section 4.2 starts to be visible. For T ¼ 1323 K (Fig. 17(b)) experimental data are few [17,32] but the reported a2-g tie-lines are perfectly reproduced. At T ¼ 1273 K (Fig. 17(c)) the extent of the single-phase fields for s, g and 3 reported in Ref. [17] are shown with dashed lines. They were determined based on XRD measurements only and accordingly small amounts of coexisting phases were hardly detected. With this in mind, the overall agreement between experimental data and calculated phase boundaries is rather good. The solubility of Ta in the (aTi) solid solution Tie6Al (wt.%) reported by Ref. [63] at 873, 1023 and 1123 K as 4e5 wt.% seems to be underestimated (Fig. 17(d)). The occurrence of the ternary O- and s-phases and the evolution of their phase fields with temperature are shown in the calculated isothermal sections in Fig. 17(d) and (e). According to the DTA measurements presented here (Fig. 8) these ternary phases start to decompose in heating at 1121  5 K and 1168  5 K, respectively.

The calculated values are 1123 K for the O-phase and 1171 K for the s-phase, following the quasi-binary reactions s / a2 þ g and O / a2 þ bo, respectively (invariant reactions p13 and p14 in Fig. 12). Phase constituents identified by XRD and EBSD reported here (Fig. 4 to 7) are in good agreement with the calculated phase boundaries (Fig. 17(d) and (e)). Both, the O-phase and s-phase are thermodynamically stable low-temperature phases down to 298 K (Fig. 17(f)). 4.4. Isopleths and selected equilibrium calculations Selected isopleths were calculated aiming to compare calculated phase transformation temperatures with experimental data from literature and from the DTA measurements reported in Section 2.25. The isopleths are displayed in Figs. 18 and 19. The experimental data are not always precisely in the plane of the calculated isopleths, but may deviate from it by maximum 2 at.%. In general, experimentally measured phase transformation temperatures agree well to the calculated phase boundaries. Few exceptions relate to low-temperature solid-state transformations where DTA measurements give slightly overestimated transformation temperatures, as it could be expected. In numerous alloys the temperature associated to A2/B2 ordering of the b-phase was clearly detected on the DTA curves and these data agree well to the calculations (see dotted lines in Figs. 18 and 19). This satisfactory agreement is achieved thanks to the fact that second order transformations are not rate limited and hence independent of applied heating rates [22]. Some differences between measured and calculated phase transition temperatures as well as additional DTA peaks in the temperature range from 1300 to 1500 K pointed by question mark (see Table 7) are perhaps related to demixing or to the A2/B2 ordering phenomena

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Alloys with composition close to Tie46Ale8Ta (at.%) received special attention, since they were at the core of alloy development efforts within the European integrated project IMPRESS [64]. The isopleths displayed in Fig. 19 pass through or close to this composition. From Fig. 19(b) and (c) it is obvious that at 1703 K the alloys with 45 at.% Al and 4e12 at% Ta exhibit a two-phase a þ b microstructure, whereas alloys with 47 at.% Al and 4e12 at% Ta are in the single a-phase field. This is fully consistent with metallographic observations on samples annealed at 1703 K, as shown in Fig. 20. From all isopleths presented in Fig. 19 it is evident that two-phase a2 þ g microstructure states are thermodynamically unstable at ambient temperatures and decompose to g þ s lower 1120 K. This aspect deserves further attention. Lapin et al. [27] presented and discussed the formation of the s-phase with B82 structure in the Ti46.5Ta7.7Al45.8 alloy after annealing at 1023 K for various times ranging from 500 to 10,000 h and found mainly grain boundary precipitates of this phase. Further investigations are needed to check, if and at what time-scale the transformation a2 þ g / s þ g would be detrimental to alloy characteristics. Phase constituents and their compositions are displayed in more detailed in Fig. 21(a) and (b) for this alloy. The measured and calculated phase compositions of a2, a and g agree well (Fig. 21(a)), and also measured area and calculated molar fractions of the phases are in reasonably good agreement (Fig. 21(b)). 5. Summary and conclusions A thermodynamic description of the full ternary system AleTaeTi has been elaborated, involving:  experimental analysis of 26 distinct alloys in both, as-cast and annealed conditions by means of XRD, SEM/EDS, SEM/EBSD, DTA and PiranieAlterthum techniques for determination of phase equilibria and phase transition temperatures (including of solideliquid transformations, phase equilibria in the ranges of: a2 þ g, ternary Ti2.17Ta0.77Al1.06 O-phase, ternary Ti3.00Ta0.89Al2.11 s-phase and (bTi,Ta,Al) miscibility gap);  the thermodynamic description the AleTaeTi system by CALPHAD modelling, taking into account experimental information from literature and from the above-mentioned measurements;  thermodynamic calculations of the reaction scheme, the liquidus, solidus and b-transus surfaces as well as selected isothermal sections, isopleths for comparison with experimental data;  thermodynamic calculations of phase constituents for the alloy Tie46Ale8Ta (at.%) Fig. 21. Composition of coexisting phases (a) and phase fractions (b) in alloy no. 17 (Ti46.5Ta7.7Al45.8) as function of temperature: open points are experimental data obtained by annealing massive g for 15.5 h in a Bridgman furnace, solid points are data from isothermal annealing experiments and lines are calculated using the proposed thermodynamic description. Dotted lines show fraction of a2 and g phases once s is suspended. Note that measured phase fractions are given in area%, while calculated phase fractions are mole%.

of the b-phase. As an example, the shift of phase boundaries that results when suppressing the ordering is shown by thin dotted lines in Fig. 19(a). It seems clear that the DTA data on the phase transition temperature correspond to the B2 and A2 mixture, rather than to disordered b or to the completely ordered bo. From this figure it is also apparent that detected by XRD demixing of the b(Ti,Ta,Al) solid solution in the samples no. 23 (Ti49.8Ta28.0Al22.2), no. 25 (Ti51.0Ta23.4Al25.6) and no. 26 (Ti54.5Ta21.8Al23.7) (see Section 2.2.3) occurred at temperatures below 945 K in the course of their furnace cooling after the 1123 K annealing.

Summarizing, one may conclude that the present thermodynamic description is consistent with the large majority of experimental data. The thermodynamic models selected for the individual phases of the system are simple, involving two or three sublattices. The simple models are convenient when aiming to integrate the AleTaeTi system into high-order systems. Acknowledgements The authors deeply appreciate Prof. T.Ya. Velikanova for fruitful discussions of results. The authors thank L.A. Duma, V.V. Garbuz, V.A. Petrova, N.I. Tsyganenko, J. Zollinger and A. Genau for technical assistance and would like to express their gratitude for financial support from the Integrated Project IMPRESS, “Intermetallic Materials Processing in Relation to Earth and Space Solidification” (Contract NMP3-CT-2004-500635) co-funded by the European

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