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Effects of mechanical alloying on the characteristics of a nanocrystalline Ti–50 at.%Al during hot pressing consolidation
Materials Science and Engineering B 168 (2010) 136–141
Contents lists available at ScienceDirect
Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb
Effects of mechanical alloying on the characteristics of a nanocrystalline Ti–50 at.%Al during hot pressing consolidation Mohammad Reza Farhang ∗ , Ali Reza Kamali, Masoud Nazarian-Samani Advanced Materials Research Center, Department of Materials Science and Engineering, Malek Ashtar University of Technology (MUT), Shahin-Shahr, Isfahan, Iran
a r t i c l e
i n f o
Article history: Received 30 July 2009 Received in revised form 21 October 2009 Accepted 23 October 2009 Keywords: Titanium aluminide Reactive sintering Non-reactive sintering Mechanical alloying Nanostructured materials
a b s t r a c t Two different powder metallurgy processes, i.e., reactive and non-reactive sintering, were investigated for the production of titanium aluminides. Ti–Al intermetallics have been successfully produced by mechanical alloying and hot pressing of powders with a nominal composition of Ti–50 at.%Al. A Ti(Al) solid solution was formed at the early stage of milling that transformed to an amorphous phase at longer milling times. On further milling, the amorphous structure transformed to a supersaturated hcp-Ti(Al) solid solution with trace amounts of TiAl after 80 h of milling, which was completely transformed to TiAl, Ti3 Al, and TiAl3 intermetallic compounds after additional milling up to 100 h milling with particles of about 200 nm. Blended elemental powders and 100 h MA-ed powders were used for the reactive and non-reactive sintering processes, respectively. The results showed that the HP-ed pre-alloyed powder had better properties in terms of density, hardness, homogeneity of microstructure, yield stress, and ductility than that produced by reactive sintering. The major contribution to the yield stress of the sintered pre-alloyed powder comes from the nanometer-sized grains of intermetallics in accordance with the Hall–Petch relation. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Titanium aluminide intermetallics are recognized as compounds with such desirable properties as low density, high strength, high modulus, and good oxidation resistance. Many efforts have been devoted to the fabrication of Ti–Al alloys using powder metallurgy (PM) processes, especially mechanical alloying (MA) [1–7], which has the additional advantage of giving rise to a good combination of strength and ductility. The process appears to be a very promising technique as it offers the opportunity not only for microstructural refinement but also for obtaining materials of high structural homogeneity. This process, a non-equilibrium method, helps produce different amorphous phases, solid solutions, or nanocrystalline compounds, the latter being potentially attractive for many applications due to the improved mechanical properties of the material which result from its reduced crystallite size down to the nanometer scale [8,9]. However, consolidation that consists of an array of sintering methods is a necessary step in the mechanical alloying of powders for any possible practical application. The sintering processes include reactive and non-reactive sintering. In the former, chemical reactions result in the production of new materials during sintering, while in the latter, no chemical
∗ Corresponding author. Tel.: +98 312 522 2508; fax: +98 312 522 8530. E-mail address: [email protected] (M.R. Farhang). 0921-5107/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2009.10.032
reaction occurs during sintering. The best consolidation technique to produce bulk, full-density, and yet nanocrystalline compacts is hot pressing (HP). Production of Ti–50 at.%Al has been reported by several researchers, all of whom introduced the supersaturated hcp-Ti(Al) solid solution as their final MA product [10–15]. The aim of this study is to evaluate the formation of intermetallic compounds during MA and HP processes. In addition, the structural evolution and morphology of the powders at different milling times are investigated. Finally, the effects of MA process on the microstructure and on the physical and mechanical properties of the samples produced will be investigated.
2. Experimental procedure Ti (60–200 m, 99.9% purity) and Al (40–100 m, 99.9% purity) powders were mixed to give the composition Ti50 Al50 (at.%) which was then charged into a WC vial with WC balls under a high purity argon atmosphere. The ball to powder weight ratio (BPR) was 10:1. The milling was performed in a Fritsch P6 planetary ball mill for periods varying from 0 to 100 h at a speed of 300 rpm. Dispersants, such as Hexane [16], which prevent agglomeration of elemental powders were not used in order to keep contamination of the mixtures to a minimum [3,5]. For the purposes of reactive and non-reactive sintering, HP was performed on two different series of powders: (1) 100 h-milled
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(pre-alloyed) powders for the non-reactive sintering process, and (2) blended elemental powders for the reactive sintering process. The powders were hot pressed under 50 MPa for 10 min at 1000 ◦ C under a high purity argon atmosphere. The HP-ed samples were 20 mm in diameter and about 12 mm in thickness. After hot pressing, the consolidated samples were ground and polished. The milled and HP-ed samples were characterized by means of X-ray diffraction (XRD) using a Seifert 3003TT diffractometer with Cu K␣, operating at 40 kV and 30 mA. The crystallite size of the as-milled powders was determined by X-ray line broadening and calculated using the Scherrer equation [17]: d=
0.9 ˇ cos
(1) 1/2
2 − ˇ2 ) , ˇM is the full width at half maximum where ˇ = (ˇM I (FWHM), ˇI is the correction factor for instrument broadening, is the angle of the peak maximum, and is the Cu K␣ weighted wavelength ( = 0.15406 nm). Using a Vega© Tescan machine, scanning electron microscopy (SEM) was carried out to characterize the morphology of the milled powder, the surfaces of the consolidated samples, and the fracture surface topography. Energy-dispersive X-ray spectroscopy (EDS) coupled with SEM was used for the semi-quantitative investigation of the microstructure of the HP-ed samples. The density of the HP-ed samples was determined using the immersion method in distilled water based on Archimedes principle. Open porosity of HP-ed samples was calculated using the mass measurements performed during density determination. Hardness measurements were conducted using a Koopa model UV1 machine. Compressive tests were carried out on cylinder samples (D: 4 mm, L: 6 mm) both at room temperature and at 700 ◦ C using an Instron model 8503 tensile/compressive testing machine at a constant strain rate of 10−2 s−1 . All the data reported are averaged out of at least three results.
3. Results and discussion 3.1. MA process 3.1.1. XRD analysis The evolution of the transformations occurring during milling was followed by XRD. Fig. 1 shows the diffraction pattern for Ti50 Al50 . It can be seen that the XRD pattern of as-received powder is almost similar to that of 10–40 h-milled powders. However, peak broadening and decreasing intensity were observed with milling time due to the decrease in crystallite size from 26 nm after 10 h to about 13 nm after 40 h milling. Shifting of the main reflexion of Ti peaks towards higher angles was also observed, which was due to the reduction in the lattice parameter attributed to a distortion of the Ti lattice by Al diffusion. From the widely accepted Ti–Al phase diagram [18], it can be seen that Al has an extended solubility in Ti, while Ti has a limited solubility in Al at different temperatures. So as expected, Al diffuses into Ti predominately during MA up to a certain value beyond which the system collapses into the amorphous state. Peak broadening at this stage led to the overlapping of peaks so that the peaks could not be easily distinguished. Further milling caused the integrated intensity of peaks to decrease which coincided with the appearance of an amorphous halo until the amorphization of the powders was complete. It is noteworthy that the contribution of the grain boundary energy to the overall energy increases with decreasing grain size. This can lead to amorphization when the free energy of the intermetallic becomes higher than that of the amorphous phase [19]. According to Fig. 1, after 60 h of MA, the peaks completely broadened between angles of 35–40◦ ; this is an indication of the
Fig. 1. XRD patterns of as-received and milled powders for different times.
amorphization of the powders. The XRD results were verified by transmission electron diffraction (TED) (not shown). It is interesting to note that the free energy curve of the amorphous phase in the Ti–Al system is lower than those of the solid solution and intermetallic phases [14]. Hence, the first phase to form is the amorphous one. Further milling up to 80 h led to the formation of a supersaturated solid solution hcp-Ti(Al) due to the crystallization of an amorphous phase or disordered Ti3 Al, because these two phases present the same unit cell [14]. Furthermore, the diffraction pattern after 80 h MA indicates the nucleation onset of the TiAl intermetallic compound, which is complete after additional milling by up to 100 h of MA. The diffraction pattern for the 100-h MA-ed powder exhibits the formation of the TiAl3 and Ti3 Al intermetallic compounds. According to the results obtained in this section, it may be concluded that the formation of intermetallic compounds is possible during MA. The enthalpy values for the formation of TiAl, Ti3 Al,
Fig. 2. Variations of lattice parameters and c/a ratio of ␣-Ti at different milling times.
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Fig. 3. SEM micrographs of as-received and milled powders for different milling times. (a) As-received powder, (b) 10 h MA-ed powder, (c) 20 h MA-ed powder, (d) 40 h MA-ed powder, (e) 60 h MA-ed powder, (f) 80 h MA-ed powder and (g) 100 h MA-ed powder.
and TiAl3 are −75, −73 and −142 kJ/mol, respectively [20,21]. Thus, the intermetallics are produced during MA due to such large negative values of enthalpy, while further heat treatment is generally required to form the intermetallics after MA [19]. In addition, the local temperature rise during MA may help promote the nucleation of the intermetallic compound. It should be noted that the nucleation of an intermetallic compound substantially alters the evolution of the alloying process by lowering the system energy [14].
Some studies have reported the formation of a stable fcc-TiN during the milling of Ti and Al [3,5] due to the impure protective gas used in the milling media. This phase was not, however, formed in our experiments because we used high purity argon as the protective gas. The lattice parameters of ␣-Ti (Fig. 2), determined from the XRD patterns according to Bragg’s relation, decreased gradually up to 40 h of milling and an abrupt decrease was observed between
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Table 1 The values for density, porosity, and hardness of Ti–50 at.%Al produced by reactive and non-reactive processes. HP-ed powder
Density (g/cm3 )
Porosity (%)
Hardness (VHN)
Elemental Pre-alloyed
3.78 3.85
5.50 3.75
417 1256
ary energy during annealing, straightness of grain boundaries, and large number of grain intersections, the latter effectively pinning the grain boundaries [19]. It can be seen in Fig. 4b that TiAl, Ti3 Al, and TiAl3 were produced after HP of elemental powders for reactive sintering.
Fig. 4. XRD patterns of HP-ed samples by: (a) non-reactive and (b) reactive processes, respectively.
40 and 80 h of milling; this finding confirms the formation of the supersaturated Ti(Al) solid solution after 80 h of milling due to the dissolution of large amounts of Al into the Ti lattice. This phase can also be confirmed by vanishing of the Al peaks. According to Fig. 2, the ratio of c/a decreased with milling time and a dramatic decrease was observed after 40 h of milling. 3.1.2. Microstructure The morphology of as-received and MA-ed powders at different milling times was investigated by SEM analysis, as shown in Fig. 3. According to Fig. 3a, Ti powders appear in irregular shapes and in various sizes, but the particles of Al are mainly round-shaped. The early stages of MA (before 10 h of MA) result in the formation of powder agglomerates (Fig. 3b) due to the ductility of Al. Further milling leads to increased deformation and work hardening, and the agglomerated powders then disintegrate into fragments producing finer particles after 10 h or after up to 80 h of milling (Fig. 3c–f). Upon further milling, particle size abruptly reduces down to about 200 nm after 100 h of milling (Fig. 3g). This is because of the formation of TiAl, Ti3 Al, and TiAl3 intermetallic compounds. It is known that the titanium aluminides as intermetallic compounds are brittle phases. Formation of these compounds leads to a sharp increase in the reduction of particle size, as shown in Fig. 3f and g. 3.2. HP processes 3.2.1. XRD analysis Fig. 4a and b present the XRD patterns of the HP-consolidated pre-alloyed and elemental powders, respectively. Comparing Fig. 4a with that of the MA-ed powder before HP, one can see the reduction in peak widths. This reduction is due to the increasing crystallite size and decreasing lattice strain; but nanostructured grains are retained after HP. It has been postulated that the higher than expected stability of nanostructured grains is the result of a number of factors such as uniformity of the crystallite size, solute segregation to grain boundaries, achievement of a low grain bound-
3.2.2. Physical and mechanical properties The density and amount of porosities obtained during reactive and non-reactive sintering have been shown in Table 1. The density of the HP-ed sample made with reactive sintering is lower than that of the HP-ed pre-alloyed powder. Both kinds of the HP-ed samples possess relatively adequate densities to exhibit considerable mechanical properties. Nanocrystalline powders have a high potential for agglomeration during MA [8]. Hence, the presence of agglomerated powders increases the possibility for large porosities to form in the material during consolidation. Moreover, finer particles can provide shorter diffusion distances during sintering; they also contribute to high concentrations of grain boundaries which act as preferable nucleation sites or as preferential diffusion paths and increase the transport rate of the alloy components in the presence of chemical potential gradients. Therefore, the density of the HP-ed pre-alloyed powder is higher than that of the HP-ed elemental one. The hardness of both HP-ed samples has also been shown in Table 1. The hardness value of the HP-ed pre-alloyed powder is almost three times higher than that of the HP-ed elemental one, which is in consistence with the density results. On the other hand, the increased density induces increased hardness. Moreover, as can be observed in the SEM micrographs (Figs. 3g and 5a), the intermetallic particles in the HP-ed pre-alloyed powder are very fine compared to those of the HP-ed elemental one (Fig. 5b). So refinement of the non-reactively sintered sample is associated with a decrease in the mean intermetallic inter-particle distance. The harmful effect of pores and porosities is indicated by a proportional relationship between density and hardness. Porosities lead to the weakening of the sample because the available stress bearing area is reduced, thereby lowering the amount of stress that a sample is able to withstand. The results of compressive tests at room temperature and at 700 ◦ C are presented in Table 2. All the tests revealed relatively brittle fracture characteristics. The compressive yield strength of the non-reactively sintered sample is higher than that of the reactively sintered one. In addition, the ductility of the first one is also greater than the second one, which reveals the improvement of ductility as a result of using the non-reactive sintering process. Enhanced compressive yield strength and improved ductility were mainly caused by nano-particle powders. This is because the volume of grain boundaries increased which naturally gave rise to fast diffusion paths. These fast paths made a few initial tiny cracks which Table 2 Compressive yield strength and ductility of Ti–50 at.%Al produced by hot pressing of elemental and pre-alloyed powders. Characteristics
Yield strength (MPa) Ductility (%)
HP-ed elemental powder HP-ed pre-alloyed powder 25 ◦ C
700 ◦ C
25 ◦ C
700 ◦ C
540 2.13
450 3.45
650 2.50
600 5.90
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Fig. 5. SEM micrographs of HP-ed samples by: (a) non-reactive and (b) reactive processes, respectively.
Fig. 6. The fracture surfaces of HP-ed samples by: (a) non-reactive and (b) reactive processes at 25 ◦ C.
quickly closed during the deformation process and, therefore, crack growth and propagation were avoided to a certain extent, contributing to the fine grain toughness. It is worth noting that the mechanical properties of metallic materials are influenced by grain boundaries, solute atoms, dislocations, second-phase particles, dispersoids, etc. Out of these, only the refinement of grain size strengthening mechanism is relevant to the present work. Grain size is known to have a significant effect on the mechanical behavior of materials and on yield stress, in particular. The dependence of yield stress on grain size in metals is well established. Yield stress, , for materials with the grain size d is found to follow the Hall–Petch relation:
3.2.3. Microstructure Fig. 5a and b illustrates the microstructure of the HP-ed samples by both reactive and non-reactive processes, respectively. As expected, the particle size of the blended elemental powder after HP is obviously larger than that of the pre-alloyed powder. In addition, more homogeneous phase distributions are observed in the latter Figure. It should be mentioned that the increase in milling time increases the lattice strain and the reduction in the crystallite size, which can improve the sinterability of the 100 h-milled powder. The EDS analysis was used to identify the three different areas revealed in both HP-ed samples shown in Fig. 5. The results of this analysis are presented in Table 3. According to these results, the likely phases are TiAl, Ti3 Al, and TiAl3 corresponding to areas 1, 2, and 3, respectively. It is obvious from Fig. 5a that the TiAl inter-
= 0 + kd−1/2
(2)
where 0 is the friction stress and k is the Hall–Petch constant [22]. This equation describes the behavior of many materials. Huang [23], Dimiduk et al. [24], Liu et al. [25], Morris and Leboeuf [26], and Bohn et al. [27] have found that the yield strength of Ti–Al alloys apparently obeys the Hall–Petch relation. This mechanism is based on dislocation pile up at the grain boundaries or dislocation slip being impeded at the grain boundaries. Therefore, the grain boundary dislocations form a Taylor dislocation forest, which inhibits dislocation motion and provides strengthening [28].
Table 3 EDS analysis of HP-ed elemental and pre-alloyed powders. HP-ed powder
Area
Ti (at.%)
Al (at.%)
Likely phase
Elemental
1 2 3
49.75 68.52 24.56
50.25 31.48 75.44
TiAl Ti3 Al TiAl3
Pre-alloyed
1 2 3
48.38 75.74 26.38
51.62 24.26 73.62
TiAl Ti3 Al TiAl3
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metallic compound has the greatest amounts of the phases. Ti3 Al is only found in some separate islands in the TiAl matrix, but TiAl3 surrounds the TiAl phase. This distribution of intermetallics during reactive sintering is due to the non-homogenous powders that consolidate during the HP process. The above-mentioned observations are not made in the HP-ed pre-alloyed powder in which a good particle distribution is obtained. The EDS results obtained are in good agreement with XRD analysis of both HP-ed samples (Fig. 4) that reveal the formation of TiAl, Ti3 Al, and TiAl3 intermetallics after both reactive and non-reactive sintering processes. Fig. 6a and b illustrates the fracture surface of the samples made by reactive and non-reactive sintering processes. Both figures exhibit a mixture of intergranular and transgranular fractures. But, as observed in Table 2, the reactively sintered sample is more brittle and, so the cleavage surfaces are evidently more conspicuous. 4. Conclusion The results of this study show that titanium aluminide powders can be produced by mechanical alloying of elemental Ti and Al at times longer than 80 h. Increasing MA time led to a dramatic decrease in the particle size of titanium aluminide powders down to about 200 nm after 100 h of milling. By reactive and non-reactive sintering processes, titanium aluminide samples could be produced from Ti and Al elemental powders and titanium aluminide powders, respectively. According to results obtained, the non-reactive sintering process promotes better physical and mechanical properties. References [1] A. Takasaki, Y. Furuya, Nanostruct. Mater. 11 (1999) 1205–1217.
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Contents lists available at ScienceDirect
Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb
Effects of mechanical alloying on the characteristics of a nanocrystalline Ti–50 at.%Al during hot pressing consolidation Mohammad Reza Farhang ∗ , Ali Reza Kamali, Masoud Nazarian-Samani Advanced Materials Research Center, Department of Materials Science and Engineering, Malek Ashtar University of Technology (MUT), Shahin-Shahr, Isfahan, Iran
a r t i c l e
i n f o
Article history: Received 30 July 2009 Received in revised form 21 October 2009 Accepted 23 October 2009 Keywords: Titanium aluminide Reactive sintering Non-reactive sintering Mechanical alloying Nanostructured materials
a b s t r a c t Two different powder metallurgy processes, i.e., reactive and non-reactive sintering, were investigated for the production of titanium aluminides. Ti–Al intermetallics have been successfully produced by mechanical alloying and hot pressing of powders with a nominal composition of Ti–50 at.%Al. A Ti(Al) solid solution was formed at the early stage of milling that transformed to an amorphous phase at longer milling times. On further milling, the amorphous structure transformed to a supersaturated hcp-Ti(Al) solid solution with trace amounts of TiAl after 80 h of milling, which was completely transformed to TiAl, Ti3 Al, and TiAl3 intermetallic compounds after additional milling up to 100 h milling with particles of about 200 nm. Blended elemental powders and 100 h MA-ed powders were used for the reactive and non-reactive sintering processes, respectively. The results showed that the HP-ed pre-alloyed powder had better properties in terms of density, hardness, homogeneity of microstructure, yield stress, and ductility than that produced by reactive sintering. The major contribution to the yield stress of the sintered pre-alloyed powder comes from the nanometer-sized grains of intermetallics in accordance with the Hall–Petch relation. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Titanium aluminide intermetallics are recognized as compounds with such desirable properties as low density, high strength, high modulus, and good oxidation resistance. Many efforts have been devoted to the fabrication of Ti–Al alloys using powder metallurgy (PM) processes, especially mechanical alloying (MA) [1–7], which has the additional advantage of giving rise to a good combination of strength and ductility. The process appears to be a very promising technique as it offers the opportunity not only for microstructural refinement but also for obtaining materials of high structural homogeneity. This process, a non-equilibrium method, helps produce different amorphous phases, solid solutions, or nanocrystalline compounds, the latter being potentially attractive for many applications due to the improved mechanical properties of the material which result from its reduced crystallite size down to the nanometer scale [8,9]. However, consolidation that consists of an array of sintering methods is a necessary step in the mechanical alloying of powders for any possible practical application. The sintering processes include reactive and non-reactive sintering. In the former, chemical reactions result in the production of new materials during sintering, while in the latter, no chemical
∗ Corresponding author. Tel.: +98 312 522 2508; fax: +98 312 522 8530. E-mail address: [email protected] (M.R. Farhang). 0921-5107/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2009.10.032
reaction occurs during sintering. The best consolidation technique to produce bulk, full-density, and yet nanocrystalline compacts is hot pressing (HP). Production of Ti–50 at.%Al has been reported by several researchers, all of whom introduced the supersaturated hcp-Ti(Al) solid solution as their final MA product [10–15]. The aim of this study is to evaluate the formation of intermetallic compounds during MA and HP processes. In addition, the structural evolution and morphology of the powders at different milling times are investigated. Finally, the effects of MA process on the microstructure and on the physical and mechanical properties of the samples produced will be investigated.
2. Experimental procedure Ti (60–200 m, 99.9% purity) and Al (40–100 m, 99.9% purity) powders were mixed to give the composition Ti50 Al50 (at.%) which was then charged into a WC vial with WC balls under a high purity argon atmosphere. The ball to powder weight ratio (BPR) was 10:1. The milling was performed in a Fritsch P6 planetary ball mill for periods varying from 0 to 100 h at a speed of 300 rpm. Dispersants, such as Hexane [16], which prevent agglomeration of elemental powders were not used in order to keep contamination of the mixtures to a minimum [3,5]. For the purposes of reactive and non-reactive sintering, HP was performed on two different series of powders: (1) 100 h-milled
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(pre-alloyed) powders for the non-reactive sintering process, and (2) blended elemental powders for the reactive sintering process. The powders were hot pressed under 50 MPa for 10 min at 1000 ◦ C under a high purity argon atmosphere. The HP-ed samples were 20 mm in diameter and about 12 mm in thickness. After hot pressing, the consolidated samples were ground and polished. The milled and HP-ed samples were characterized by means of X-ray diffraction (XRD) using a Seifert 3003TT diffractometer with Cu K␣, operating at 40 kV and 30 mA. The crystallite size of the as-milled powders was determined by X-ray line broadening and calculated using the Scherrer equation [17]: d=
0.9 ˇ cos
(1) 1/2
2 − ˇ2 ) , ˇM is the full width at half maximum where ˇ = (ˇM I (FWHM), ˇI is the correction factor for instrument broadening, is the angle of the peak maximum, and is the Cu K␣ weighted wavelength ( = 0.15406 nm). Using a Vega© Tescan machine, scanning electron microscopy (SEM) was carried out to characterize the morphology of the milled powder, the surfaces of the consolidated samples, and the fracture surface topography. Energy-dispersive X-ray spectroscopy (EDS) coupled with SEM was used for the semi-quantitative investigation of the microstructure of the HP-ed samples. The density of the HP-ed samples was determined using the immersion method in distilled water based on Archimedes principle. Open porosity of HP-ed samples was calculated using the mass measurements performed during density determination. Hardness measurements were conducted using a Koopa model UV1 machine. Compressive tests were carried out on cylinder samples (D: 4 mm, L: 6 mm) both at room temperature and at 700 ◦ C using an Instron model 8503 tensile/compressive testing machine at a constant strain rate of 10−2 s−1 . All the data reported are averaged out of at least three results.
3. Results and discussion 3.1. MA process 3.1.1. XRD analysis The evolution of the transformations occurring during milling was followed by XRD. Fig. 1 shows the diffraction pattern for Ti50 Al50 . It can be seen that the XRD pattern of as-received powder is almost similar to that of 10–40 h-milled powders. However, peak broadening and decreasing intensity were observed with milling time due to the decrease in crystallite size from 26 nm after 10 h to about 13 nm after 40 h milling. Shifting of the main reflexion of Ti peaks towards higher angles was also observed, which was due to the reduction in the lattice parameter attributed to a distortion of the Ti lattice by Al diffusion. From the widely accepted Ti–Al phase diagram [18], it can be seen that Al has an extended solubility in Ti, while Ti has a limited solubility in Al at different temperatures. So as expected, Al diffuses into Ti predominately during MA up to a certain value beyond which the system collapses into the amorphous state. Peak broadening at this stage led to the overlapping of peaks so that the peaks could not be easily distinguished. Further milling caused the integrated intensity of peaks to decrease which coincided with the appearance of an amorphous halo until the amorphization of the powders was complete. It is noteworthy that the contribution of the grain boundary energy to the overall energy increases with decreasing grain size. This can lead to amorphization when the free energy of the intermetallic becomes higher than that of the amorphous phase [19]. According to Fig. 1, after 60 h of MA, the peaks completely broadened between angles of 35–40◦ ; this is an indication of the
Fig. 1. XRD patterns of as-received and milled powders for different times.
amorphization of the powders. The XRD results were verified by transmission electron diffraction (TED) (not shown). It is interesting to note that the free energy curve of the amorphous phase in the Ti–Al system is lower than those of the solid solution and intermetallic phases [14]. Hence, the first phase to form is the amorphous one. Further milling up to 80 h led to the formation of a supersaturated solid solution hcp-Ti(Al) due to the crystallization of an amorphous phase or disordered Ti3 Al, because these two phases present the same unit cell [14]. Furthermore, the diffraction pattern after 80 h MA indicates the nucleation onset of the TiAl intermetallic compound, which is complete after additional milling by up to 100 h of MA. The diffraction pattern for the 100-h MA-ed powder exhibits the formation of the TiAl3 and Ti3 Al intermetallic compounds. According to the results obtained in this section, it may be concluded that the formation of intermetallic compounds is possible during MA. The enthalpy values for the formation of TiAl, Ti3 Al,
Fig. 2. Variations of lattice parameters and c/a ratio of ␣-Ti at different milling times.
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Fig. 3. SEM micrographs of as-received and milled powders for different milling times. (a) As-received powder, (b) 10 h MA-ed powder, (c) 20 h MA-ed powder, (d) 40 h MA-ed powder, (e) 60 h MA-ed powder, (f) 80 h MA-ed powder and (g) 100 h MA-ed powder.
and TiAl3 are −75, −73 and −142 kJ/mol, respectively [20,21]. Thus, the intermetallics are produced during MA due to such large negative values of enthalpy, while further heat treatment is generally required to form the intermetallics after MA [19]. In addition, the local temperature rise during MA may help promote the nucleation of the intermetallic compound. It should be noted that the nucleation of an intermetallic compound substantially alters the evolution of the alloying process by lowering the system energy [14].
Some studies have reported the formation of a stable fcc-TiN during the milling of Ti and Al [3,5] due to the impure protective gas used in the milling media. This phase was not, however, formed in our experiments because we used high purity argon as the protective gas. The lattice parameters of ␣-Ti (Fig. 2), determined from the XRD patterns according to Bragg’s relation, decreased gradually up to 40 h of milling and an abrupt decrease was observed between
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Table 1 The values for density, porosity, and hardness of Ti–50 at.%Al produced by reactive and non-reactive processes. HP-ed powder
Density (g/cm3 )
Porosity (%)
Hardness (VHN)
Elemental Pre-alloyed
3.78 3.85
5.50 3.75
417 1256
ary energy during annealing, straightness of grain boundaries, and large number of grain intersections, the latter effectively pinning the grain boundaries [19]. It can be seen in Fig. 4b that TiAl, Ti3 Al, and TiAl3 were produced after HP of elemental powders for reactive sintering.
Fig. 4. XRD patterns of HP-ed samples by: (a) non-reactive and (b) reactive processes, respectively.
40 and 80 h of milling; this finding confirms the formation of the supersaturated Ti(Al) solid solution after 80 h of milling due to the dissolution of large amounts of Al into the Ti lattice. This phase can also be confirmed by vanishing of the Al peaks. According to Fig. 2, the ratio of c/a decreased with milling time and a dramatic decrease was observed after 40 h of milling. 3.1.2. Microstructure The morphology of as-received and MA-ed powders at different milling times was investigated by SEM analysis, as shown in Fig. 3. According to Fig. 3a, Ti powders appear in irregular shapes and in various sizes, but the particles of Al are mainly round-shaped. The early stages of MA (before 10 h of MA) result in the formation of powder agglomerates (Fig. 3b) due to the ductility of Al. Further milling leads to increased deformation and work hardening, and the agglomerated powders then disintegrate into fragments producing finer particles after 10 h or after up to 80 h of milling (Fig. 3c–f). Upon further milling, particle size abruptly reduces down to about 200 nm after 100 h of milling (Fig. 3g). This is because of the formation of TiAl, Ti3 Al, and TiAl3 intermetallic compounds. It is known that the titanium aluminides as intermetallic compounds are brittle phases. Formation of these compounds leads to a sharp increase in the reduction of particle size, as shown in Fig. 3f and g. 3.2. HP processes 3.2.1. XRD analysis Fig. 4a and b present the XRD patterns of the HP-consolidated pre-alloyed and elemental powders, respectively. Comparing Fig. 4a with that of the MA-ed powder before HP, one can see the reduction in peak widths. This reduction is due to the increasing crystallite size and decreasing lattice strain; but nanostructured grains are retained after HP. It has been postulated that the higher than expected stability of nanostructured grains is the result of a number of factors such as uniformity of the crystallite size, solute segregation to grain boundaries, achievement of a low grain bound-
3.2.2. Physical and mechanical properties The density and amount of porosities obtained during reactive and non-reactive sintering have been shown in Table 1. The density of the HP-ed sample made with reactive sintering is lower than that of the HP-ed pre-alloyed powder. Both kinds of the HP-ed samples possess relatively adequate densities to exhibit considerable mechanical properties. Nanocrystalline powders have a high potential for agglomeration during MA [8]. Hence, the presence of agglomerated powders increases the possibility for large porosities to form in the material during consolidation. Moreover, finer particles can provide shorter diffusion distances during sintering; they also contribute to high concentrations of grain boundaries which act as preferable nucleation sites or as preferential diffusion paths and increase the transport rate of the alloy components in the presence of chemical potential gradients. Therefore, the density of the HP-ed pre-alloyed powder is higher than that of the HP-ed elemental one. The hardness of both HP-ed samples has also been shown in Table 1. The hardness value of the HP-ed pre-alloyed powder is almost three times higher than that of the HP-ed elemental one, which is in consistence with the density results. On the other hand, the increased density induces increased hardness. Moreover, as can be observed in the SEM micrographs (Figs. 3g and 5a), the intermetallic particles in the HP-ed pre-alloyed powder are very fine compared to those of the HP-ed elemental one (Fig. 5b). So refinement of the non-reactively sintered sample is associated with a decrease in the mean intermetallic inter-particle distance. The harmful effect of pores and porosities is indicated by a proportional relationship between density and hardness. Porosities lead to the weakening of the sample because the available stress bearing area is reduced, thereby lowering the amount of stress that a sample is able to withstand. The results of compressive tests at room temperature and at 700 ◦ C are presented in Table 2. All the tests revealed relatively brittle fracture characteristics. The compressive yield strength of the non-reactively sintered sample is higher than that of the reactively sintered one. In addition, the ductility of the first one is also greater than the second one, which reveals the improvement of ductility as a result of using the non-reactive sintering process. Enhanced compressive yield strength and improved ductility were mainly caused by nano-particle powders. This is because the volume of grain boundaries increased which naturally gave rise to fast diffusion paths. These fast paths made a few initial tiny cracks which Table 2 Compressive yield strength and ductility of Ti–50 at.%Al produced by hot pressing of elemental and pre-alloyed powders. Characteristics
Yield strength (MPa) Ductility (%)
HP-ed elemental powder HP-ed pre-alloyed powder 25 ◦ C
700 ◦ C
25 ◦ C
700 ◦ C
540 2.13
450 3.45
650 2.50
600 5.90
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Fig. 5. SEM micrographs of HP-ed samples by: (a) non-reactive and (b) reactive processes, respectively.
Fig. 6. The fracture surfaces of HP-ed samples by: (a) non-reactive and (b) reactive processes at 25 ◦ C.
quickly closed during the deformation process and, therefore, crack growth and propagation were avoided to a certain extent, contributing to the fine grain toughness. It is worth noting that the mechanical properties of metallic materials are influenced by grain boundaries, solute atoms, dislocations, second-phase particles, dispersoids, etc. Out of these, only the refinement of grain size strengthening mechanism is relevant to the present work. Grain size is known to have a significant effect on the mechanical behavior of materials and on yield stress, in particular. The dependence of yield stress on grain size in metals is well established. Yield stress, , for materials with the grain size d is found to follow the Hall–Petch relation:
3.2.3. Microstructure Fig. 5a and b illustrates the microstructure of the HP-ed samples by both reactive and non-reactive processes, respectively. As expected, the particle size of the blended elemental powder after HP is obviously larger than that of the pre-alloyed powder. In addition, more homogeneous phase distributions are observed in the latter Figure. It should be mentioned that the increase in milling time increases the lattice strain and the reduction in the crystallite size, which can improve the sinterability of the 100 h-milled powder. The EDS analysis was used to identify the three different areas revealed in both HP-ed samples shown in Fig. 5. The results of this analysis are presented in Table 3. According to these results, the likely phases are TiAl, Ti3 Al, and TiAl3 corresponding to areas 1, 2, and 3, respectively. It is obvious from Fig. 5a that the TiAl inter-
= 0 + kd−1/2
(2)
where 0 is the friction stress and k is the Hall–Petch constant [22]. This equation describes the behavior of many materials. Huang [23], Dimiduk et al. [24], Liu et al. [25], Morris and Leboeuf [26], and Bohn et al. [27] have found that the yield strength of Ti–Al alloys apparently obeys the Hall–Petch relation. This mechanism is based on dislocation pile up at the grain boundaries or dislocation slip being impeded at the grain boundaries. Therefore, the grain boundary dislocations form a Taylor dislocation forest, which inhibits dislocation motion and provides strengthening [28].
Table 3 EDS analysis of HP-ed elemental and pre-alloyed powders. HP-ed powder
Area
Ti (at.%)
Al (at.%)
Likely phase
Elemental
1 2 3
49.75 68.52 24.56
50.25 31.48 75.44
TiAl Ti3 Al TiAl3
Pre-alloyed
1 2 3
48.38 75.74 26.38
51.62 24.26 73.62
TiAl Ti3 Al TiAl3
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metallic compound has the greatest amounts of the phases. Ti3 Al is only found in some separate islands in the TiAl matrix, but TiAl3 surrounds the TiAl phase. This distribution of intermetallics during reactive sintering is due to the non-homogenous powders that consolidate during the HP process. The above-mentioned observations are not made in the HP-ed pre-alloyed powder in which a good particle distribution is obtained. The EDS results obtained are in good agreement with XRD analysis of both HP-ed samples (Fig. 4) that reveal the formation of TiAl, Ti3 Al, and TiAl3 intermetallics after both reactive and non-reactive sintering processes. Fig. 6a and b illustrates the fracture surface of the samples made by reactive and non-reactive sintering processes. Both figures exhibit a mixture of intergranular and transgranular fractures. But, as observed in Table 2, the reactively sintered sample is more brittle and, so the cleavage surfaces are evidently more conspicuous. 4. Conclusion The results of this study show that titanium aluminide powders can be produced by mechanical alloying of elemental Ti and Al at times longer than 80 h. Increasing MA time led to a dramatic decrease in the particle size of titanium aluminide powders down to about 200 nm after 100 h of milling. By reactive and non-reactive sintering processes, titanium aluminide samples could be produced from Ti and Al elemental powders and titanium aluminide powders, respectively. According to results obtained, the non-reactive sintering process promotes better physical and mechanical properties. References [1] A. Takasaki, Y. Furuya, Nanostruct. Mater. 11 (1999) 1205–1217.
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