Phase equilibria in niobium rich Nb-Al-Ti alloys

Scripta METALLURGICA et MATERIALIA

Vol. 27, pp. 265-270, 1992 Printed in the U.S.A.

Pergamon Press Ltd. All rights reserved

PHASE EQUILIBRIA IN NIOBIUM RICH Nb-AI-Ti ALLOYS

E. S. K. Menon, P. R. Subramanian¢ and D. M. Dimiduk § *Systran Corporation, 4126 Linden Avenue, Dayton, OH 45432 Universal Energy Systems, Inc., Dayton, Ohio 45432 § WL/MLLM, WPAFB, Ohio 45433. (Received May 4, 1992) (Revised May 26, 1992)

Introduction Recently, niobium based alloys have attracted much attention as potential high temperature materials [1]. Nb-A1 alloys have been examined, especially with a view to producing intermetallic matrix composite materials with a distribution of ductile bcc precipitates in the NbaA1 matrix [2,3]. The NbaAI phase exhibits the A15 crystal structure and is extremely brittle at temperatures below 1475 K. Attempts at alloy development are being carded out by several researchers and studies on Ti and other alloying additions to NbA1 alloys are being done [4]. At the same time, several studies on Nb additions to Ti-A1 alloys are available [5-7] and phase stability in this ternary system is still being critically examined [8-10]. However, most of the phase boundaries in the Nb-rich comer of the ternary involving the bcc, NbaA1 and the Nb2A1 phases, shown in reported isothermal sections of the Ti-A1-Nb phase diagram [8,11,12], are only speculative, based on judicious extrapolations of the binary Nb-A1 phase diagram proposed by Jorda et al [13] a number of years ago. In the present study, phase boundaries involving the bcc and the NbaA1 phase were experimentally determined and isothermal sections of the Nb rich comer of the Ti-A1-Nb system established at 1923 K, 1473 K and 1273K. Experimental Procedures The alloys were prepared by argon arc meking the pure metals into 250g-size cigars. The cigars were melted four to six times to ensure homogeneity and were further homogenized by heat treating the alloys at 1923 K for 50 hours in argon atmosphere. The samples were wrapped in tantalum and a continuous flow of the inert gas was maintained during heat treatment; upon termination of the heat treatment, the alloys were quenched by switching off the furnace and increasing the rate of inert gas flow. The compositions of the homogenized alloys as determined by chemical analysis (metallic elements by atomic absorption spectroscopy) are given in Table I.

265 0956-716X/92 $5.00 + .00 Copyright (c) 1992 Pergamon Press Ltd.

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Table I

Chemical Composition of Alloys AHoy No.

at% Nb

at% Al

at% Ti

!1 2 3 4 5

90.11 82.53 73.95 65.35 73.78

9.89 17.47 15.60 13.25 16.82

10.45 21.40 9.40

carbon (ppm) 129 350 145 219 259

loxygen (ppm) 180 30 460 690 400

nitrogen (ppm) 60 40 50 130 160

The homogenized alloys were subsequently equilibrated at 1473 K and 1273 K for 14 days and 30 days, respectively. Samples cut from the heat treated alloys were examined by optical and scanning electron microscopy, x-ray diffraction (XRD) and transmission electron microscopy (TEM). Chemical composition of the individual phases present in these samples was determined by electron microprobe analysis (EPMA). Results and Discussion

Metallographic examination of the homogenized alloys showed that alloys 2, 3 and 5 contained a distribution of second phase precipitates ( Figs. 1 (a) and (b)) which was identified to be the NbaA1 phase by XRD, TEM and EPMA. Alloys 1 and 5 showed a single phase structure and XRD patterns from these could be consistently indexed to be bcc. TEM examination of these samples as well as the matrix phase in homogenized alloys 2, 3 and 5

Fig. 1 : Optical micrographs illustrating the typical appearance of Nb3A1 plates in heat treated samples (a) Nb-17.47at% Al, aged at 1923 K for 50 hours and (b) Nb-13.25at%Al21.40at%Ti, aged at 1273K for 30 days subsequent to homogenization at 1923 K.

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showed that the crystal structure of this phase was B2. Typical selected area diffraction patterns obtained from these samples are shown in Figs. 2 (a) and Co). Notice that there are no superlattice spots at ~ [222]~ positions in the [110] pattern in Fig. 2 (b) and thus it is clear that the structure is B2 and not DO3. All of the previous structure determinations on Nb-A1 alloys by XRD failed to notice the ordered bee structure [13-15] while all of the electron diffraction work on these alloys identified the B2 structure [16-18]. In this study, all the alloys heat treated at 1473 K and 1273 K showed a distribution of the Nb3A1 phase with the A15 crystal structure in a B2 matrix. Details of the microstructural evolution in these alloys will be presented in a later publication.

Fig. 2 : Selected area diffraction patterns illustrating that the crystal structure of the matrix phase in microstructures such as Fig. 1 is ordered CsCl type. (a) [100]~c pattern from as cast Nb-13.25at% Al-21.40at% Ti and (b) [110]~ pattern from as cast Nb-9.87at% Al. Results of the composition analysis by EPMA from all five alloys are consolidated in the phase diagrams shown in Figs. 3 and 4. The results from the present work from alloys 1 and 2 are included in the Nb-A1 phase diagram shown in Fig. 3. The full lines are the phase boundaries established by Jorda et al [13]. The dashed lines drawn on the basis of the present work suggest that the solubility of A1 in Nb is significantly lower than that determined by Jorda et al [13], especially at lower temperatures, and may be due to the longer annealing periods employed in the present study*. The phase boundaries and their variation with temperature in Nb-rich Nb-A1-Ti alloys are ploUed as isothermal sections in Fig. 4. The tie line compositions were determined from the average of at least 5 to 8 composition determinations. Volume fraction measurements by point counting technique agreed very well with the phase boundaries shown in these phase diagrams. A portion of the isothermal section * In Ref. [12], the samples were annealed for 7 days.

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at 1473 K as determined by Perepezko et al [8] is shown by the dotted lines, and here again, it appears that the bcc/(bcc+Nb3A1) phase boundary is shifted to lower A1 concentrations similar to that observed in the binary Nb-A1 system. These results suggest that A1 diffusion in Nb is quite sluggish and is substantiated by the observation of very slow growth of precipitates from supersaturated Nb-A1 and Nb-A1-Ti solid solutions [our unpublished work]. It may also be pointed out that the (bcc+Nb3Al+Nb2A1) region in [8] and [12] are positioned differently and our results would match better with that of [12]. Phase composition analysis on hot worked and equilibrated samples from the alloys studied here are in progress to confirm the phase boundaries reported in this paper.

2100

"e

B~

Vol.

30

Fig. 3 : Nb-rich end of the binary Nb-Al phase diagram. The full lines are from Ref. [8] and the dashed lines are from the present work.

40

An important observation made in the present study is the identification of the crystal structure of the matrix phase in quenched Nb-A1 and NbA1-Ti alloys to be B2. The B2 structure has been found in quenched Nb-A1 [16-18] as well as in several Ti-A1-Nb alloys [8-11, 19-22] before. The question that remains unanswered is whether or not a bcc ~ B2

ordering transformation occurred during quenching. In order to answer this question, homogenized samples of alloys 1, 3 and 4 were aged at various temperatures between 673 K and 1573 K for less. than 16 hours and subsequently water quenched. The aging time was adjusted so that Nb3A1 precipitation did not occur. An examination of the SAD pattems showed that the intensity of the 100 superlattice spots in samples of alloy 1 aged at 1073 K and lower temperatures was much stronger than those quenched from higher temperatures. Similarly, the supedattice reflections in samples from alloy 4 quenched from all temperatures from 1473 K and below were stronger in comparison to those from samples quenched from 1573K or higher. These results suggest that the ordering temperatures in alloys 1 and 4 are close to 1073 K and 1473 K, respectively. Figure 5 illustrates the typical antiphase domain structure observed in the samples that were held below the ordering temperature. These results clearly indicate that the bcc ~ B2 ordering occurs in a metastable fashion in these compositions since the experimentally determined phase boundaries lie above the observed order-disorder transformation temperatures. Also, it is clear that the ordering temperature rises drastically with increasing Ti concentrations. Now, the dashed lines in Fig. 4 trace out the expected phase boundary corresponding to bcc --* B2 ordering in these alloys.

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+ ,,?

Nb /

I

I

I

10

20

30

40

at% TI

(b)

7A15

~

7C

/

/ Nb

I

I

10

20

at%

/1

I

TI

30

40

boundaries observed in the B2 phase. Nb16.82at% AI-9.4Oat% Ti, as cast. This dark field micrograph was obtained with a 100 superlattice reflection and the zone axis is [llO]~c.

40

(c)

3O

Nb

!

10

I

20

at%TI

Fig. 5 : Typical anti-phase domain

I

30

40

• ALLOY COMPOSITION X EPMA DATA

Fig. 4 : Nb-rich corner of the ternary NbAI-Ti phase diagram. Isothermal sections at (a) 1923 K, (b) 1473 K and (c) 1273 K. Tie lines constructed on the basis of EPMA data are shown. Dashed lines are from the Ref. [13] and the dotted lines indicate the bcc/B2 phase boundary.

The important conclusions from this study are: (a) the solubility of AI in Nb is considerably less than that previously established; (b) isothermal sections and tie lines at 1923 K, 1473 K and 1273 K have been established; (c) the high temperature bcc phase in a wide range of Nb-A1-Ti alloys undergoes an ordering transformation to produce the B2 structure and (d) plausible bcc --, B2 ordering curves have also been shown in the phase diagrams. Acknowledgements The authors gratefully acknowledge the assistance provided by T. Campbell and M. Dodd in heat treatments, S. Apt in TEM

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foil prepration and C. Begg in microprobe work. ESKM acknowledges the support provided by the AFOSR contract No. F33615-90-C-5944 with the Materials Directorate, Wright Laboratory, Wright-Patterson Air Force Base. References

1. F.H. Froes, JOM, 4 1 , 6 (1989). 2. D.L. Anton and D. M. Shah, in "High Temperature Ordered Intermetallic Alloys IV", MRS Proceedings, 213, Eds. L. A. Johnson, D. P. Pope and J. O. Stiegler, Pittsburgh, PA, 1991, p. 733. 3. T. Kummagai, S. Hanada and S. Saito, in "Intermetallic Compounds - Structure and Mechanical Properties, JIMIS - 6 ", Eds. O. Izumi, The Japan Institute of Metals, Sendai, Japan, 1991, p. 1039. 4. T. Fujiwara, K. Yasuda and H. Kodama, ibid, p. 633. 5. D.T. Hoelzer and F. Ebrahimi, in "Intermetallic Matrix Composites", MRS Proceedings, 194, Eds. D. L. Anton, P. L. Martin, D. B. Miracle and R. McMeeking, Pittsburgh, 1990, p. 393. 6. D.M. Dimiduk, D. B. Miracle and C. H. Ward, To appear in Mater. Sci. Tech., April, 1992. 7. R.G. Rowe, Advanced Materials & Processes, 141 No. 3, 33 (1991). 8. J.H. Perepezko, Y. A. Chang, L. E. Seitzman, J. C. Lin, N. R. Bonda, T. J. Jewett and J. C. Mishurda, in "High Temperature Aluminides and Intermetallics", Eds. S. H. Whang, C. T. Liu, D. P. Pope and J. O. Stiegler, The Minerals, Metals and Materials Society, Warrendale, PA, 1990, p. 19. 9. C. Suryanarayana and D. S. Lee, Scripta Metall. et Mater., 26, 919 (1992). 10. A. G. Jackson and D. S. Lee, Submitted to Scripta Metall. et Mater. 11. G. Inden and A. Hellwig • Private Communication. 12. K. Kaltenbach, S. Gama, D. G. Pinatti, K. Schulze and E-T. Henig, Zeit. Metallkde, 80, 535 (1989). 13 J.L. Jorda, R. Flukiger and J. Muller, J. of the Less Comm. Met., 7 5 , 2 2 7 (1980). 14. C. E. Lundin and A. S. Yamamoto, TMS-AIME, 236 863 (1968). 15. G. W. Webb, IEEE Transactions on Magnetics, vol. MAG-15, 616 (1979). 16. V. M. Pan, V. I. Latysheva, Ye. N. Khusid, V. N. Minakov and Ye. V. Turtsevich, Physics of Metals and Metallography, 40, 52 (1975). 17. T. N. Marieb, A. D. Kaiser, S. R. Nutt, D. L. Anton and D. M. Shah, in "High Temperature Ordered Intermetallic Alloys IV", MRS Proceedings, 213, Eds. L. A. Johnson, D. P. Pope and J. O. Stiegler, Pittsburgh, PA, 1991, p. 329. 18. M. Aindow, J. Shyue, T. A. Gasper and H. L. Fraser, Phil. Mag. Lett., 64, 59 (1991). 19. L. A. Bendersky and W. J. Boettinger, in "High Temperature Ordered Intermetallic Alloys IIr', MRS Proceedings, 1~ , Eds. C. T. Liu, A. I. Taub, N. S. Stoloff and C. C. Koch, Pittsburgh, PA, 1989, p. 45. 20. R. Strychor and J. C. Williams, in "Proceedings of the International Conference on Solid --) Solid Phase Transformations", Eds. H. I. Aaronson, D. E. Laughlin, R. F.Sekerka and C. M. Wayman, TMS-AIME, New York, p. 249. 21. R. Strychor, J. C. Williams and W. A. Soffa, Met. Trans., 19A, 225 (1988). 22. H. L. Fraser, Private Communication.