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Structural defects and thermal stability of Ti(Al) solid solution obtained by mechanical alloying
Materials Science and Engineering A242 (1998) 230 – 234
Structural defects and thermal stability of Ti(Al) solid solution obtained by mechanical alloying V.I. Fadeeva a, A.V. Leonov a, E. Szewczak b, H. Matyja b,* a
b
Department of Chemistry, Moscow State Uni6ersity, Vorobio6y Gory, 119889 Moscow-234, Russia Department of Materials Science and Engineering, Warsaw Uni6ersity of Technology, Narbutta 85, 02 -524 Warsaw, Poland Received 8 January 1997; received in revised form 18 June 1997
Abstract The structure of a metastable Ti(Al) hexagonal close packed (h.c.p.) solid solution obtained by mechanical alloying of the Al50Ti50 powder, was investigated using X-ray diffractometry (XRD). An analysis of the profiles shows physical broadening of (100), (002) and (101) diffraction lines. The probability of stacking faults on the basal (0001) and prismatic {101( 0} planes was determined from the differences between the physical broadening of the (100) and (101) lines which are influenced by the stacking faults and the (002) line which is not affected by the stacking faults. The total concentration of stacking faults is comparable to the values obtained earlier for Al–Ti solid solutions deformed by filing. Thermo-desorption and mass-spectrometery showed that, during mechanical alloying, about 1 at.% of hydrogen is dissolved in the alloy. A significant amount of this hydrogen contributes to the formation of stacking faults on the basal and prismatic planes. The temperature range within which the metastable h.c.p. structure is transformed into the equilibrium AlTi (L10) structure was determined by DSC calorimetry. The transition of the Ti(Al) solid solution into the AlTi intermetallic proceeds through an intermediate stage of the metastable f.c.c. phase with a lattice parameter a= 0.4012 nm. © 1998 Elsevier Science S.A. Keywords: Structural defects; Thermal stability; Mechanical alloying
1. Introduction A considerable number of investigations recently performed on mechanical alloying of Ti – Al powder mixtures, were inspired by the necessity to develop a technology which would enable us to obtain ductile intermetallics of high strength. In order to describe the solid state reaction mechanism that occurs when Al and Ti powders are milled it is important to not only trace the changes of phase composition, but also determine the influence of prolonged milling on the structure of the phases formed. Phase analysis by X-ray diffraction (XRD) of mechanically alloyed equiatomic Ti – Al alloys shows that the solid state reaction of Al and Ti can take place in various ways. Thus, in a series of investigations it was ascertained that an increase of the milling time only resulted in a decrease of the Ti and Al particle sizes, which was manifested by diffraction lines broadening * Corresponding author. Fax: +48 22 484875. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 0 3 - 0
and a decrease in their intensity. After a certain milling time, an amorphous phase was formed [1–5]. In various types of mills, full amorphization has been obtained after different milling times—from 40 h when using a high energy Fritsch Pulverisette planetary ball mill [5] up to 500 h when using a low energy conventional horizontal ball mill [3]. According to other investigators [6–8] a metastable solid solution of Al in Ti with a hexagonal close packed (h.c.p.) lattice was formed prior to the amorphous phase. The XRD analysis showed that even a considerable prolongation of milling time did not change the structure of the Ti(Al) solid solution. One should however suppose that, at the same time, defects resulting from gliding and dissociation of dislocations will be formed in the crystal lattice of the solid solution. Defects, such as ‘chaotic’, correlated and dissociated dislocations, cause the diffraction lines to broaden. How the ‘chaotic’ dislocations (microstrains), correlated dislocations (small-angle boundaries of coherently diffracting domains) and dissociated dislocations
V.I. Fadee6a et al. / Materials Science and Engineering A242 (1998) 230–234
(stacking faults) contribute to physical broadening of the diffraction peaks was described by Warren [9]. According to the Warren theory, the stacking fault (s.f.) probability for some structures can be determined from the shifts of diffraction lines. This method has recently been used by Gayle et al. [10], who estimated the s.f. probability in a mechanically alloyed f.c.c. Cu –Co solid solution. The effect of deformation produced by filing and by alloying on the formation of s.f. in Ti and its alloys was described in [11 – 13]. It was found that if component, (in this case) Al, is dissolved in the a-Ti lattice, the s.f. energy decreases. The present paper is concerned with the results of measurements of s.f. probability in the supersaturated Ti–Al solid solutions produced by mechanical alloying of Al and Ti powders. The influence of the s.f. on the structural transformations during subsequent heating is discussed.
2. Experimental procedure Elemental powders of Al with purity of 99.9% and particle size of 40 – 50 mm, and Ti with purity of 99.0% and particle size of 150 – 200 mm were used for preparing the Al–50at.% Ti mixture. Mechanical alloying was carried out in a high energy planetary ball mill in an argon atmosphere. The mass of the powder was 5 g and the ball-to-powder ratio was 20:1. The XRD examinations were performed with Philips 1830 and DRON-4 diffractometers using the CuKa radiation. Diffraction patterns were recorded at a stepscanning D(2u)= 0.05°. The profiles of the (100), (002), (101) lines of the h.c.p. phase (the indices of the a-Ti structure) were measured in the step-scanning mode at the steps D(2u)= 0.02° in order to evaluate the s.f. probability. The overlapping (002) and (101) lines were separated by fitting the line profiles using the Lorentz function [14]. The profiles were corrected for the Ka doublet by the Rachinger method [15]. The instrumental broadening was corrected using an annealed standard Ti sample. The lattice parameters were determined with the errors Da = 93 × 10 − 4 nm, Dc = 9 5×10 − 4 nm. A thermal analysis was carried out in a Perkin-Elmer DSC-7 calorimeter using a heating rate of 40 K min − 1 between 150 and 720°C. The gas contamination content in mechanically alloyed powders was measured using the thermo-desorption method, under continuous heating conditions, with a heating rate of 10 K min − 1 from 100 to 750°C. The released gas was collected in a closed gas collector calibrated according to the ‘gas amount-gas pressure’ principle. The composition of the gas mixture was determined in the static mode by mass spectrometry.
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3. Results and discussion The deformation behaviour of the material strongly depends on the milling time. During the first 3 h of milling the powder coated the balls. The inner vial walls were coated after 6 h of milling, but the powder was easily detachable and after low intensity grinding in a mortar it seemed homogeneous. Changes in the phase composition of the mechanically alloyed Al50Ti50 alloy at various milling times are shown in Fig. 1. The shift of a-Ti diffraction lines toward larger u angles and the increase in their intensity is observed after 3 h of milling, because the h.c.p. phase becomes the main phase in the alloy. After 6 h of milling the XRD pattern indicates that the alloy is entirely single-phase. The structure of the alloy does not change after 10 h of milling. Assuming the similarity between the structure of the phase obtained and a-Ti structure we found its lattice parameters to be: a= 0.2865 nm, c= 0.4605 nm. These values are considerably lower than a-Ti parameters (a= 0.2950, c = 0.4686 nm) and indicate that Al formed a solid solution in Ti. As seen in Fig. 2, the values of the parameters are consistent with data reported in the literature (the diagram representing the variation of the parameters a and c with the Ti(Al) solid solution concentration extrapolated to 50 at.% of Al) [16]. This confirms that all the aluminium is dissolved in Ti as early as after 6 h of milling. The range of metastable Ti(Al) solid solutions obtained by mechanical alloying is considerably wider than that reported in [8]. Taking into account the fact that the Ti(Al) solid solution is supersaturated and that it is deformed by milling, one can presume that the Ti(Al) lattice is highly distorted due to the presence of dislocations, including dissociated dislocations known as deformation stacking
Fig. 1. X-ray diffraction patterns of Al50Ti50 initial (a) and mechanically alloyed powders after milling for (b) 2 h, (c) 6 h and (d) 10 h.
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232
Fig. 2. Dependence of a and c parameters of Ti(Al) solid solution vs. aluminium content [16].
faults. It is well known, that deformation s.f. in a h.c.p. structure are formed when the dislocations glide on the basal (0001) [112( 0] planes and the prismatic {101( 0} 112( 0 planes. According to the Warren model, no interaction between stacking faults of different types occurs and they contribute to the diffraction line broadening independently. This enables the probabilities of the stacking faults formed during dissociation of basal (a1) and prismatic (a2) dislocations to be determined separately. The stacking faults of both types (a1 and a2) are deformation faults. Assuming that, in a mechanically alloyed Ti(Al) solid solution, the s.f. probability due to crystal growth is negligible, the effect of these s.f. on the diffraction line broadening will be ignored. For a h.c.p. lattice, the physical broadening of diffraction lines, bhkl, due only to the a1 and a2 probabilities, can be written as [13]: bhkl =
2dhkl tg u (G(0001) · a1 +G{101( 0} · a2) pa
G(0001) =
3dhkla l c2
G{101( 0} =
)
2dhkl 1 % h+ k ap j 2
)
Fig. 3. The broadening of different diffraction lines of h.c.p. phases: (1a, 2a) Al50Ti50 after MA for 6 h; (1b, 2b) Al50Ti50 after MA for 6 h and annealing at 550°C; (3) annealed Ti.
(2)
In accordance with Warren’s assumption [9] concerning the isotropy of microstrains and of coherently diffracting domains, bhkl cos u should exhibit a linear increase vs. sin u if the stacking faults are absent. As shown in Fig. 3, bhkl cos u vs. sin u for the (100), (002) and (101) lines of the Ti(Al) solid solution is not linear, in contrast to the analogous diagram obtained for annealed a-Ti. The considerably higher values of (100) and (101) lines broadening compared to the broadening of (002) line, indicate that both types of s.f. are present in the lattice. All the three lines, i.e. (100), (002) and (101) run close to each other and their broadenings due to microstrains and coherent domains size can be considered to be equal at the initial stage of calculations. Therefore, the differences b(100) − b(002) and b(101) − b(002) can be attributed to the presence of s.f.
(3)
Table 1 Stacking fault probabilities in h.c.p. Ti(Al) solid solutions
(1)
where a and c are h.c.p. lattice parameters, dhkl is the interplanar distance, p is the multiplicity factor, G(0001)and G{101( 0} can be obtained from Eq. (2) and Eq. (3) for the diffraction lines that satisfy the condition h−k = 3N 91 (N =0,1,2…) and are broadened by s.f. In this case, for the (100) line —G(0001) = 0 and G{101( 0} =0.577; for the (101) line —G(0001) =0.508 and G{101( 0} =0.886. At the same time, no s.f. of any type contributes to the broadening of lines, such as the (002) line, for which h −k = 3N.
S.f. probability
Alloy treatment MA 6h
MA 6h+annealing 550°C
Filinga
13×10−3 [12] 14×10−3 [13]
a1
3×10−3 (1) 0 (2)
11×10−3 (1) 5×10−3 (2)
a2
19×10−3 (1) 20×10−3 (2)
23×10−3 (1) 26×10−3 (2)
— —
a Stacking fault probabilities were measured in Ti – Al solid solutions containing 3.5 – 12.5 at.% Al.
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233
Fig. 5. DSC curve for Al50Ti50 alloy after MA for 6 h.
Fig. 4. Integral (1) and differential (2) curves of H2 thermodesorption for Al50Ti50 alloy after MA 6 h.
Table 1 gives measured s.f. probabilities a1 and a2 for the Ti(Al) solid solution. In order to confirm the results of s.f. measurements, the experiment was conducted twice. Samples 1 and 2 are the samples obtained after 6 h of milling from each of these experiments. The line broadenings are measured for MA alloys and for samples heated in a calorimeter up to a temperature of 550°C at which the h.c.p. solid solution is still present. One can observe that, for MA alloys, the s.f. probability on the prismatic planes is about 2%, whereas s.f. on the basal planes either do not exist or their probability a1 is very low. The total value of the s.f. probability corresponds to values of 1×10 − 2 – 1.8 × 10 − 2 determined for the Ti – 3.5% Al solid solution, which has been deformed by filing [9,17]. After the mechanically alloyed powder is annealed, the s.f. probability a1 due to basal gliding increases. This can be attributed to an increase in the dislocation mobility and to the obstacles on basal planes (0001) being removed due to the rearrangement or removal of impurities. The latter could probably accumulate in the lattice of the solid solution during prolonged milling. Thermo-desorption examinations of the mechanically alloyed Al50Ti50 alloy show that in the 400–700°C temperature range, gas release takes place (Fig. 4). The maximum thermo-desorption rate occurs at a temperature of 550°C. Mass-spectrometry analysis shows, that the released gas is hydrogen. The amount of evacuated hydrogen related to the powder mass corresponds to 1 at.% H2 dissolved in the alloy. The amount of hydrogen in the starting Ti powder was below 0.16 at.%.
The DSC curve in Fig. 5 exhibits an exothermic peak within the 590–650°C temperature range, associated with the transformation of the metastable solid solution into the equilibrium AlTi phase. In the XRD pattern of the sample after heating up to 720°C (Fig. 6), we can see the diffraction lines characteristic of the tetragonal AlTi (L10) phase and one weak line of the h.c.p. phase. The alloy becomes entirely single-phase after isothermal anealing at 800°C for 1 h. The phase transition of the metastable Ti(Al) solid solution into equilibrium TiAl intermetallic proceeds through the formation of an intermediate f.c.c. phase. F.c.c. phase lines can be observed in the XRD pattern obtained for the sample heated to 615°C during DSC measurements (Fig. 6). For a heating rate of 40°C min − 1 this temperature corresponds to the initial part of the exothermic peak on the calorimetric curve. The f.c.c. phase is coherent
Fig. 6. The change of X-ray diffraction pattern of the mechanically alloyed Al50Ti50 alloy after subsequent annealing at: (a) 550°C, (b) 615°C, (c) 635°C, (d) 720°C.
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with the h.c.p. phase and the orientation relationship between those phases is: (0001)h.c.p. (111)f.c.c., [112¯0]h.c.p. [110]f.c.c.. The lattice parameter of the f.c.c. phase is a =0.4012 nm. The coherent nucleation of the cubic phase and its growth on the base of the h.c.p. structure is easy to explain if one considers the structural correlation: the h.c.p. lattice B = \ f.c.c. stacking faults. Stacking faults in a hexagonal phase formed in basal (0001) planes during heating have a cubic configuration and can be used as cubic phase nuclei in the metastable Ti(Al) solid solution. An analogous transformation mechanism from the h.c.p. to L10 structures in Ti–Al alloys was examined by Zhang et al. and Abe et al. [18,19]. The metastable f.c.c. phase transforms into an ordered L10 phase with tetragonal distortion. The distortion results in the diffraction lines undergoing splitting, such as, e.g. (200)f.c.c. (002)(200)L10 and (220)f.c.c. (202)(220)L10, etc. We can see this splitting in the diffraction patterns obtained for the alloy annealed at 635 and 720°C (Fig. 6c and d).
4. Conclusions Aluminium was completely dissolved in the h.c.p.-Ti lattice when an equiatomic elemental powder mixture was subjected to mechanical alloying. The lattice of this solid solution contains deformation stacking faults which form due to dissociation of dislocations on the prismatic {101( 0} and basal (0001) planes. The phase transformation sequence during heating of a mechanically alloyed alloy from the metastable Ti(Al) solid solution to the equilibrium AlTi intermetallic compound is: h.c.p. f.c.c.L10. The intermediate f.c.c. phase is metastable and nucleates on the base of stacking faults. The latter are formed in the hexagonal structure by MA followed or not by heating. One can suppose that elements like H and C that contaminate
.
the alloy during milling, affect the gliding of dislocation and the formation of stacking faults.
Acknowledgements The authors wish to thank Dr A.N. Streletskii for thermo-desorption experiment. The financial support of the Science Research Committee, contract number 7 T08D 043 10 is gratefully acknowledged.
References [1] L. Schultz, Mater. Sci. Eng. 97 (1988) 15. [2] M.S. El-Eskandarany, K. Aoki, K. Suzuki, Materials Science Forum 88/90 (1992) 81. [3] J.H. Ahn, H.S. Chung, R. Watanabe, Y.H. Park, Materials Science Forum 88/90 (1992) 347. [4] M. Qi, M. Zhu, G.B. Li, D.Z. Yang, K.H. Kuo, Materials Science Forum 88/90 (1992) 356. [5] B.S. Murty, M.D. Naik, M. Mohan Rao, S. Ranganathan, Materials Science Forum 16 (1992) 19. [6] C. Suryanarayana, Guo-Hao Chen, Frefer Abdulbaset, F.H. Froes, Mater. Sci. Eng. A158 (1992) 93. [7] T. Itsukaichi, K. Masuyama, M. Umemoto, I. Okane, J.G. Cabanas-Moreno, J. Mater. Res. 8 (1993) 1817. [8] N. Burgio, W. Guo, M. Magini, F. Padella, S. Martelli, I. Soletta, Proc. ASM Int. Conf. Structural Applications of Mechanical Alloying, Myrth Beach, South Carolina, SC, 27–29 March, 1990. [9] B.E. Warren, Prog. Metall. Phys. 8 (1958) 147. [10] F.W. Gayle, F.S. Biancaniello, Nanostr. Mater., in press. [11] R. Krishnan, Z. Metallkunde 58 (1967) 811. [12] E.A. Metzbower, Metall. Trans. 2 (1971) 3099. [13] A.S. Shishmakov, D.A. Mirzaev, J.F. Kmelinin, R.A. Adamesku, Fiz. Met Metalloved (in Russian) 37 (1974) 313. [14] S. Sato, J. Appl. Phys. 1 (1962) 210. [15] W.A. Rachinger, J. Sci. Instr. 25 (1948) 254. [16] W.B. Pearson, in: G.V. Raynor (Ed.), A Handbook of Lattice Spacings and Structures of Metals and Alloys, vol. 4, Pergamon Press Oxford, 1958. [17] R.K. Gupta, T.K. Namboodire, P. Ramarao, Proc. Symp. Mater. Sci. Res. Nat. Aeronaut. Lab. Bangalore 2 (1970) 479. [18] Y.G. Zhang, F.D. Tichelaar, F.W. Schapink, Q. Xu, C.Q. Chen, Scripta Metall. Mater. 32 (1993) 981. [19] E. Abe, T. Kumagai, M. Nakamura, Phil. Mag. Lett. 72 (1995) 291.
Structural defects and thermal stability of Ti(Al) solid solution obtained by mechanical alloying V.I. Fadeeva a, A.V. Leonov a, E. Szewczak b, H. Matyja b,* a
b
Department of Chemistry, Moscow State Uni6ersity, Vorobio6y Gory, 119889 Moscow-234, Russia Department of Materials Science and Engineering, Warsaw Uni6ersity of Technology, Narbutta 85, 02 -524 Warsaw, Poland Received 8 January 1997; received in revised form 18 June 1997
Abstract The structure of a metastable Ti(Al) hexagonal close packed (h.c.p.) solid solution obtained by mechanical alloying of the Al50Ti50 powder, was investigated using X-ray diffractometry (XRD). An analysis of the profiles shows physical broadening of (100), (002) and (101) diffraction lines. The probability of stacking faults on the basal (0001) and prismatic {101( 0} planes was determined from the differences between the physical broadening of the (100) and (101) lines which are influenced by the stacking faults and the (002) line which is not affected by the stacking faults. The total concentration of stacking faults is comparable to the values obtained earlier for Al–Ti solid solutions deformed by filing. Thermo-desorption and mass-spectrometery showed that, during mechanical alloying, about 1 at.% of hydrogen is dissolved in the alloy. A significant amount of this hydrogen contributes to the formation of stacking faults on the basal and prismatic planes. The temperature range within which the metastable h.c.p. structure is transformed into the equilibrium AlTi (L10) structure was determined by DSC calorimetry. The transition of the Ti(Al) solid solution into the AlTi intermetallic proceeds through an intermediate stage of the metastable f.c.c. phase with a lattice parameter a= 0.4012 nm. © 1998 Elsevier Science S.A. Keywords: Structural defects; Thermal stability; Mechanical alloying
1. Introduction A considerable number of investigations recently performed on mechanical alloying of Ti – Al powder mixtures, were inspired by the necessity to develop a technology which would enable us to obtain ductile intermetallics of high strength. In order to describe the solid state reaction mechanism that occurs when Al and Ti powders are milled it is important to not only trace the changes of phase composition, but also determine the influence of prolonged milling on the structure of the phases formed. Phase analysis by X-ray diffraction (XRD) of mechanically alloyed equiatomic Ti – Al alloys shows that the solid state reaction of Al and Ti can take place in various ways. Thus, in a series of investigations it was ascertained that an increase of the milling time only resulted in a decrease of the Ti and Al particle sizes, which was manifested by diffraction lines broadening * Corresponding author. Fax: +48 22 484875. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 0 3 - 0
and a decrease in their intensity. After a certain milling time, an amorphous phase was formed [1–5]. In various types of mills, full amorphization has been obtained after different milling times—from 40 h when using a high energy Fritsch Pulverisette planetary ball mill [5] up to 500 h when using a low energy conventional horizontal ball mill [3]. According to other investigators [6–8] a metastable solid solution of Al in Ti with a hexagonal close packed (h.c.p.) lattice was formed prior to the amorphous phase. The XRD analysis showed that even a considerable prolongation of milling time did not change the structure of the Ti(Al) solid solution. One should however suppose that, at the same time, defects resulting from gliding and dissociation of dislocations will be formed in the crystal lattice of the solid solution. Defects, such as ‘chaotic’, correlated and dissociated dislocations, cause the diffraction lines to broaden. How the ‘chaotic’ dislocations (microstrains), correlated dislocations (small-angle boundaries of coherently diffracting domains) and dissociated dislocations
V.I. Fadee6a et al. / Materials Science and Engineering A242 (1998) 230–234
(stacking faults) contribute to physical broadening of the diffraction peaks was described by Warren [9]. According to the Warren theory, the stacking fault (s.f.) probability for some structures can be determined from the shifts of diffraction lines. This method has recently been used by Gayle et al. [10], who estimated the s.f. probability in a mechanically alloyed f.c.c. Cu –Co solid solution. The effect of deformation produced by filing and by alloying on the formation of s.f. in Ti and its alloys was described in [11 – 13]. It was found that if component, (in this case) Al, is dissolved in the a-Ti lattice, the s.f. energy decreases. The present paper is concerned with the results of measurements of s.f. probability in the supersaturated Ti–Al solid solutions produced by mechanical alloying of Al and Ti powders. The influence of the s.f. on the structural transformations during subsequent heating is discussed.
2. Experimental procedure Elemental powders of Al with purity of 99.9% and particle size of 40 – 50 mm, and Ti with purity of 99.0% and particle size of 150 – 200 mm were used for preparing the Al–50at.% Ti mixture. Mechanical alloying was carried out in a high energy planetary ball mill in an argon atmosphere. The mass of the powder was 5 g and the ball-to-powder ratio was 20:1. The XRD examinations were performed with Philips 1830 and DRON-4 diffractometers using the CuKa radiation. Diffraction patterns were recorded at a stepscanning D(2u)= 0.05°. The profiles of the (100), (002), (101) lines of the h.c.p. phase (the indices of the a-Ti structure) were measured in the step-scanning mode at the steps D(2u)= 0.02° in order to evaluate the s.f. probability. The overlapping (002) and (101) lines were separated by fitting the line profiles using the Lorentz function [14]. The profiles were corrected for the Ka doublet by the Rachinger method [15]. The instrumental broadening was corrected using an annealed standard Ti sample. The lattice parameters were determined with the errors Da = 93 × 10 − 4 nm, Dc = 9 5×10 − 4 nm. A thermal analysis was carried out in a Perkin-Elmer DSC-7 calorimeter using a heating rate of 40 K min − 1 between 150 and 720°C. The gas contamination content in mechanically alloyed powders was measured using the thermo-desorption method, under continuous heating conditions, with a heating rate of 10 K min − 1 from 100 to 750°C. The released gas was collected in a closed gas collector calibrated according to the ‘gas amount-gas pressure’ principle. The composition of the gas mixture was determined in the static mode by mass spectrometry.
231
3. Results and discussion The deformation behaviour of the material strongly depends on the milling time. During the first 3 h of milling the powder coated the balls. The inner vial walls were coated after 6 h of milling, but the powder was easily detachable and after low intensity grinding in a mortar it seemed homogeneous. Changes in the phase composition of the mechanically alloyed Al50Ti50 alloy at various milling times are shown in Fig. 1. The shift of a-Ti diffraction lines toward larger u angles and the increase in their intensity is observed after 3 h of milling, because the h.c.p. phase becomes the main phase in the alloy. After 6 h of milling the XRD pattern indicates that the alloy is entirely single-phase. The structure of the alloy does not change after 10 h of milling. Assuming the similarity between the structure of the phase obtained and a-Ti structure we found its lattice parameters to be: a= 0.2865 nm, c= 0.4605 nm. These values are considerably lower than a-Ti parameters (a= 0.2950, c = 0.4686 nm) and indicate that Al formed a solid solution in Ti. As seen in Fig. 2, the values of the parameters are consistent with data reported in the literature (the diagram representing the variation of the parameters a and c with the Ti(Al) solid solution concentration extrapolated to 50 at.% of Al) [16]. This confirms that all the aluminium is dissolved in Ti as early as after 6 h of milling. The range of metastable Ti(Al) solid solutions obtained by mechanical alloying is considerably wider than that reported in [8]. Taking into account the fact that the Ti(Al) solid solution is supersaturated and that it is deformed by milling, one can presume that the Ti(Al) lattice is highly distorted due to the presence of dislocations, including dissociated dislocations known as deformation stacking
Fig. 1. X-ray diffraction patterns of Al50Ti50 initial (a) and mechanically alloyed powders after milling for (b) 2 h, (c) 6 h and (d) 10 h.
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232
Fig. 2. Dependence of a and c parameters of Ti(Al) solid solution vs. aluminium content [16].
faults. It is well known, that deformation s.f. in a h.c.p. structure are formed when the dislocations glide on the basal (0001) [112( 0] planes and the prismatic {101( 0} 112( 0 planes. According to the Warren model, no interaction between stacking faults of different types occurs and they contribute to the diffraction line broadening independently. This enables the probabilities of the stacking faults formed during dissociation of basal (a1) and prismatic (a2) dislocations to be determined separately. The stacking faults of both types (a1 and a2) are deformation faults. Assuming that, in a mechanically alloyed Ti(Al) solid solution, the s.f. probability due to crystal growth is negligible, the effect of these s.f. on the diffraction line broadening will be ignored. For a h.c.p. lattice, the physical broadening of diffraction lines, bhkl, due only to the a1 and a2 probabilities, can be written as [13]: bhkl =
2dhkl tg u (G(0001) · a1 +G{101( 0} · a2) pa
G(0001) =
3dhkla l c2
G{101( 0} =
)
2dhkl 1 % h+ k ap j 2
)
Fig. 3. The broadening of different diffraction lines of h.c.p. phases: (1a, 2a) Al50Ti50 after MA for 6 h; (1b, 2b) Al50Ti50 after MA for 6 h and annealing at 550°C; (3) annealed Ti.
(2)
In accordance with Warren’s assumption [9] concerning the isotropy of microstrains and of coherently diffracting domains, bhkl cos u should exhibit a linear increase vs. sin u if the stacking faults are absent. As shown in Fig. 3, bhkl cos u vs. sin u for the (100), (002) and (101) lines of the Ti(Al) solid solution is not linear, in contrast to the analogous diagram obtained for annealed a-Ti. The considerably higher values of (100) and (101) lines broadening compared to the broadening of (002) line, indicate that both types of s.f. are present in the lattice. All the three lines, i.e. (100), (002) and (101) run close to each other and their broadenings due to microstrains and coherent domains size can be considered to be equal at the initial stage of calculations. Therefore, the differences b(100) − b(002) and b(101) − b(002) can be attributed to the presence of s.f.
(3)
Table 1 Stacking fault probabilities in h.c.p. Ti(Al) solid solutions
(1)
where a and c are h.c.p. lattice parameters, dhkl is the interplanar distance, p is the multiplicity factor, G(0001)and G{101( 0} can be obtained from Eq. (2) and Eq. (3) for the diffraction lines that satisfy the condition h−k = 3N 91 (N =0,1,2…) and are broadened by s.f. In this case, for the (100) line —G(0001) = 0 and G{101( 0} =0.577; for the (101) line —G(0001) =0.508 and G{101( 0} =0.886. At the same time, no s.f. of any type contributes to the broadening of lines, such as the (002) line, for which h −k = 3N.
S.f. probability
Alloy treatment MA 6h
MA 6h+annealing 550°C
Filinga
13×10−3 [12] 14×10−3 [13]
a1
3×10−3 (1) 0 (2)
11×10−3 (1) 5×10−3 (2)
a2
19×10−3 (1) 20×10−3 (2)
23×10−3 (1) 26×10−3 (2)
— —
a Stacking fault probabilities were measured in Ti – Al solid solutions containing 3.5 – 12.5 at.% Al.
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Fig. 5. DSC curve for Al50Ti50 alloy after MA for 6 h.
Fig. 4. Integral (1) and differential (2) curves of H2 thermodesorption for Al50Ti50 alloy after MA 6 h.
Table 1 gives measured s.f. probabilities a1 and a2 for the Ti(Al) solid solution. In order to confirm the results of s.f. measurements, the experiment was conducted twice. Samples 1 and 2 are the samples obtained after 6 h of milling from each of these experiments. The line broadenings are measured for MA alloys and for samples heated in a calorimeter up to a temperature of 550°C at which the h.c.p. solid solution is still present. One can observe that, for MA alloys, the s.f. probability on the prismatic planes is about 2%, whereas s.f. on the basal planes either do not exist or their probability a1 is very low. The total value of the s.f. probability corresponds to values of 1×10 − 2 – 1.8 × 10 − 2 determined for the Ti – 3.5% Al solid solution, which has been deformed by filing [9,17]. After the mechanically alloyed powder is annealed, the s.f. probability a1 due to basal gliding increases. This can be attributed to an increase in the dislocation mobility and to the obstacles on basal planes (0001) being removed due to the rearrangement or removal of impurities. The latter could probably accumulate in the lattice of the solid solution during prolonged milling. Thermo-desorption examinations of the mechanically alloyed Al50Ti50 alloy show that in the 400–700°C temperature range, gas release takes place (Fig. 4). The maximum thermo-desorption rate occurs at a temperature of 550°C. Mass-spectrometry analysis shows, that the released gas is hydrogen. The amount of evacuated hydrogen related to the powder mass corresponds to 1 at.% H2 dissolved in the alloy. The amount of hydrogen in the starting Ti powder was below 0.16 at.%.
The DSC curve in Fig. 5 exhibits an exothermic peak within the 590–650°C temperature range, associated with the transformation of the metastable solid solution into the equilibrium AlTi phase. In the XRD pattern of the sample after heating up to 720°C (Fig. 6), we can see the diffraction lines characteristic of the tetragonal AlTi (L10) phase and one weak line of the h.c.p. phase. The alloy becomes entirely single-phase after isothermal anealing at 800°C for 1 h. The phase transition of the metastable Ti(Al) solid solution into equilibrium TiAl intermetallic proceeds through the formation of an intermediate f.c.c. phase. F.c.c. phase lines can be observed in the XRD pattern obtained for the sample heated to 615°C during DSC measurements (Fig. 6). For a heating rate of 40°C min − 1 this temperature corresponds to the initial part of the exothermic peak on the calorimetric curve. The f.c.c. phase is coherent
Fig. 6. The change of X-ray diffraction pattern of the mechanically alloyed Al50Ti50 alloy after subsequent annealing at: (a) 550°C, (b) 615°C, (c) 635°C, (d) 720°C.
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with the h.c.p. phase and the orientation relationship between those phases is: (0001)h.c.p. (111)f.c.c., [112¯0]h.c.p. [110]f.c.c.. The lattice parameter of the f.c.c. phase is a =0.4012 nm. The coherent nucleation of the cubic phase and its growth on the base of the h.c.p. structure is easy to explain if one considers the structural correlation: the h.c.p. lattice B = \ f.c.c. stacking faults. Stacking faults in a hexagonal phase formed in basal (0001) planes during heating have a cubic configuration and can be used as cubic phase nuclei in the metastable Ti(Al) solid solution. An analogous transformation mechanism from the h.c.p. to L10 structures in Ti–Al alloys was examined by Zhang et al. and Abe et al. [18,19]. The metastable f.c.c. phase transforms into an ordered L10 phase with tetragonal distortion. The distortion results in the diffraction lines undergoing splitting, such as, e.g. (200)f.c.c. (002)(200)L10 and (220)f.c.c. (202)(220)L10, etc. We can see this splitting in the diffraction patterns obtained for the alloy annealed at 635 and 720°C (Fig. 6c and d).
4. Conclusions Aluminium was completely dissolved in the h.c.p.-Ti lattice when an equiatomic elemental powder mixture was subjected to mechanical alloying. The lattice of this solid solution contains deformation stacking faults which form due to dissociation of dislocations on the prismatic {101( 0} and basal (0001) planes. The phase transformation sequence during heating of a mechanically alloyed alloy from the metastable Ti(Al) solid solution to the equilibrium AlTi intermetallic compound is: h.c.p. f.c.c.L10. The intermediate f.c.c. phase is metastable and nucleates on the base of stacking faults. The latter are formed in the hexagonal structure by MA followed or not by heating. One can suppose that elements like H and C that contaminate
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the alloy during milling, affect the gliding of dislocation and the formation of stacking faults.
Acknowledgements The authors wish to thank Dr A.N. Streletskii for thermo-desorption experiment. The financial support of the Science Research Committee, contract number 7 T08D 043 10 is gratefully acknowledged.
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