Microstructure and Mechanical Properties of High Relative Density γ -TiAl Alloy Using Irregular Pre-Alloyed Powder

: Preparing high relative density γ -TiAl alloy by pressure-less sintering at low-cost has always been a challenge. Therefore, a new kind of non-spherical pre-alloyed TiAl powder was prepared by the reaction of TiH 2 powder and Al powder at 800 ◦ C to fabricate high-density Ti-48Al alloy via pressure-less sintering. The oxygen content was controlled to below 1800 ppm by using coarse Al powder (~120 µ m). The sintered densities ranged from 92.1% to 97.5% with sintering temperature varying from 1300 ◦ C to 1450 ◦ C. The microstructure of the sintered compact was greatly inﬂuenced by the sintering temperature. The as-sintered samples had a near- γ structure at 1350 ◦ C, a duplex structure at 1400 ◦ C, and a nearly lamellar structure at 1450 ◦ C. To achieve full densiﬁcation, non-capsule hot isostatic pressing was performed on the 1350 ◦ C and 1400 ◦ C sintered samples. As a result, high compressive strengths of 2241 MPa and 1931MPa were obtained, which were higher than the existing Ti-48Al alloys.


Introduction
γ-TiAl-based alloys with low density, high mechanical properties at elevated temperatures, excellent creep performance and good corrosion resistance have been widely used in the recent decades, which are considered one of the most potential materials for aerospace and automobile industries [1][2][3][4]. In the aerospace industry, γ-TiAl can be used instead of Inconel 718 to prepare components, such as turbine blades and compressor blades. In the automotive industry, these materials are suitable for racing and high-end vehicle parts, such as engine valves, turbine wheels and connecting rods [5,6]. For example, Ti-48Al-2Cr-2Nb alloy has been used to manufacture the six-stage and seven-stage low pressure turbine blades in GEnx engines [7]. However, TiAl alloys generally suffer from low ductility at room temperature, resulting in the manufacturing difficulties. Some reports have found that adjusting the grain size in a suitable range and improving the homogeneity can be beneficial for improving the workability of this alloy [8][9][10].
In recent years, powder metallurgy (PM) technique is of special interest because of no-composition segregation, uniform microstructure, fine grain and low cost [11][12][13]. The economical cold-compact and pressure-less sintering process is one of the most attractive PM fabrication approaches for γ-TiAl [14], in which the raw powder is the most important part of pressure-less sintering.
Two types of TiAl powders are often used: Pre-alloyed (PA) powder, prepared by gas atomization and rotating electrode atomization, and the blended elemental (BE) powder composed of Ti, Al and other alloying elements. PA powder with a mean particle size

Sample Preparation
Raw materials were TiH 2 powder produced by hydrogenation dehydrogenation process and Al powder produced by gas atomization. The particle sizes of the raw powders are shown in Table 1. Al powders with different particle sizes were used to study the interstitials content of the prepared TiAl alloys, especially oxygen content. Firstly, the TiH 2 and Al powders were blended in a mixer under Ar atmosphere for 180 min with a mass ratio of 66:34. Then, the mixed powder was respectively pretreated in a vacuum rotary furnace at 500, 600, 700 and 800 • C for 300 min with a heating rate of 5 • C min −1 . Subsequently, the Ti-48Al PA powder was crushed by a vibratory milling under vacuum until the mean powder grain size was about 15 µm. The particle size distributions were measured using the BT-9300S laser particle size distribution analyzer. These details would systematically discuss later. After that, the PA powder was compacted by cold isostatic pressing (CIP) at 200 MPa for 120 s. Then, the compacts were respectively sintered at 1300, 1350, 1400, and 1450 • C for 150 min with a heating rate of 3 • C·min −1 . Accordingly, these samples were referred to simply as T1300, T1350, T1400 and T1450. Finally, HIP was performed on the sintered T1350 and T1400 samples to obtain high relative density, and the two samples were marked as H1350 and H1400. The HIP parameters were 1200 • C/140 MPa/120 min. The schematic diagram for the preparation process of PM TiAl alloy is presented in Figure 1. simply as T1300, T1350, T1400 and T1450. Finally, HIP was performed on the sintered T1350 and T1400 samples to obtain high relative density, and the two samples were marked as H1350 and H1400. The HIP parameters were 1200 °C/140 MPa/120 min. The schematic diagram for the preparation process of PM TiAl alloy is presented in Figure 1.

Testing and Characterization
Sintered densities were measured by the Archimedes method according to the ASTM standard B328. Phase constituents were analyzed by X-ray diffraction (XRD, D/max-RB, Japan) with Cu Kα radiation (λ ¼ 0.15418 nm) at 50 kV and 200 mA with a diffraction angle 2θ from 10° to 90° and scanning rate of 5°·min −1 . The oxygen content of the samples was measured by non-dispersive infrared analysis, while nitrogen and hydrogen contents were determined by the thermal conductivity method, using an Eltra ONH-2000 apparatus. Powder morphology and as-sintered microstructures were analyzed by a scanning electron microscopy (SEM, Philips LEO-1450) equipped with energy-dispersive spectroscopy (EDS). Hardness was tested with a Vickers hardness tester (EM1500) at room temperature, and the loading pressure was 1 kg. Compression tests were carried out at room temperature using an INS-TRON 4206 machine. The test pieces were cylinders of 3 mm in diameter and 6 mm in length. The nominal strain rate for the compression tests was 10 −3 s −1 . Room temperature tensile tests were conducted using an AGI-250 KN testing machine with a strain rate of 10 −3 s −1 . Besides, element distribution was characterized by Electron Microprobe (EPMA, 1720).

PA Powder Analysis
In this study, the blended TiH2 and Al powder was heated to a high temperature to form porous products, which took advantage of the Kirkendall effect between Ti and Al [11]. Then, the partially sintered porous products were crushed and milled to produce the non-spherical PA powder (~15 μm) with large specific surface area. At a high temperature, the possible reactions between TiH2 and Al are shown in Equations (1)-(4) [32]: 6TiH2→6TiH1.5 + 1.5H2→6Ti + 6H2 (1) 6Ti + 6Al→4Ti + 2TiAl3 (2)

Testing and Characterization
Sintered densities were measured by the Archimedes method according to the ASTM standard B328. Phase constituents were analyzed by X-ray diffraction (XRD, D/max-RB, Japan) with Cu Kα radiation (λ 1 4 0.15418 nm) at 50 kV and 200 mA with a diffraction angle 2θ from 10 • to 90 • and scanning rate of 5 • ·min −1 . The oxygen content of the samples was measured by non-dispersive infrared analysis, while nitrogen and hydrogen contents were determined by the thermal conductivity method, using an Eltra ONH-2000 apparatus. Powder morphology and as-sintered microstructures were analyzed by a scanning electron microscopy (SEM, Philips LEO-1450) equipped with energy-dispersive spectroscopy (EDS). Hardness was tested with a Vickers hardness tester (EM1500) at room temperature, and the loading pressure was 1 kg. Compression tests were carried out at room temperature using an INS-TRON 4206 machine. The test pieces were cylinders of 3 mm in diameter and 6 mm in length. The nominal strain rate for the compression tests was 10 −3 s −1 . Room temperature tensile tests were conducted using an AGI-250 KN testing machine with a strain rate of 10 −3 s −1 . Besides, element distribution was characterized by Electron Microprobe (EPMA, 1720).
In order to select the appropriate reaction temperature, the Gibbs free energy ∆G needs to be calculated according to Equation (5): In the formula, ∆H is the enthalpy change, T is the temperature, and ∆S is the entropy change. Figure 2 shows the ∆G of Equations (1) and (2) with a function of temperature. The ∆G value was less than 0 when the temperature was above 500 • C, indicating that the reaction can proceed spontaneously since the ambient temperature was above 500 • C. Therefore, the minimum reaction temperature was set as 500 • C. The reaction of TiH 2 and Al included two reaction processes: Firstly, the endothermic reaction of TiH 2 decomposing into Ti and H 2 . Secondly, the exothermic reaction of Ti and Al forming TiAl 3 . When temperature was lower than 800 • C, the ∆G > 0 of TiH 2 , the decomposition reaction indicated that the heat released by the formation of TiAl 3 between Ti and Al was absorbed by the decomposition of TiH 2 .
In order to select the appropriate reaction temperature, the Gibbs free energy ΔG needs to be calculated according to Equation (5): In the formula, ΔH is the enthalpy change, T is the temperature, and ΔS is the entropy change. Figure 2 shows the ΔG of Equations (1) and (2) with a function of temperature. The ΔG value was less than 0 when the temperature was above 500 °C, indicating that the reaction can proceed spontaneously since the ambient temperature was above 500 °C. Therefore, the minimum reaction temperature was set as 500 °C. The reaction of TiH2 and Al included two reaction processes: Firstly, the endothermic reaction of TiH2 decomposing into Ti and H2. Secondly, the exothermic reaction of Ti and Al forming TiAl3. When temperature was lower than 800 °C, the ΔG > 0 of TiH2, the decomposition reaction indicated that the heat released by the formation of TiAl3 between Ti and Al was absorbed by the decomposition of TiH2. It has been reported that Ti and Al powders synthesize TiAl by self-propagation high-temperature synthesis under pressure-less and high temperature conditions [33]. The reaction product has a non-metallurgical combination of layered structure with segregation. This is because the combustion wave takes place in the form of periodic oscillations or spins during the combustion process. In such cases, it is difficult to obtain micronsized powder by mechanically crushing the product with a layered structure, which cannot be applied to the pressure-less sintering process. In order to utilize TiH2 decomposition to absorb the heat released by the reaction of Ti and Al and prevent the self-propagating reaction, the upper limit of the reaction temperature was selected to be 800 °C.
The reactions in Equations (3) and (4) describe the diffusion process between TiAl3 and Ti, during which Ti3Al, TiAl, TiAl2 were formed at the interface. When TiAl3 was still present, TiAl and TiAl2 diffuse competitively and grow up at the same time. When TiAl3 was consumed, Ti3Al and TiAl2 continued to decrease, and TiAl phase continued to increase. There were Ti3Al and TiAl phases in the final reaction product [34].
Based on the thermodynamic calculation results, the blended TiH2 and Al powders were designed to pretreat at 500, 600, 700, and 800 °C. After treatment, the phases of the powders are shown in Figure 3a. At 500 °C, TiH2 decomposed partially, and the main phases were TiH1.5 and Al. There was no obvious diffraction peak of intermetallic compound. At 600 °C, Al had been completely consumed, and the main phases were γ-TiAl, It has been reported that Ti and Al powders synthesize TiAl by self-propagation high-temperature synthesis under pressure-less and high temperature conditions [33]. The reaction product has a non-metallurgical combination of layered structure with segregation. This is because the combustion wave takes place in the form of periodic oscillations or spins during the combustion process. In such cases, it is difficult to obtain micron-sized powder by mechanically crushing the product with a layered structure, which cannot be applied to the pressure-less sintering process. In order to utilize TiH 2 decomposition to absorb the heat released by the reaction of Ti and Al and prevent the self-propagating reaction, the upper limit of the reaction temperature was selected to be 800 • C.
The reactions in Equations (3) and (4) describe the diffusion process between TiAl 3 and Ti, during which Ti 3 Al, TiAl, TiAl 2 were formed at the interface. When TiAl 3 was still present, TiAl and TiAl 2 diffuse competitively and grow up at the same time. When TiAl 3 was consumed, Ti 3 Al and TiAl 2 continued to decrease, and TiAl phase continued to increase. There were Ti 3 Al and TiAl phases in the final reaction product [34].
Based on the thermodynamic calculation results, the blended TiH 2 and Al powders were designed to pretreat at 500, 600, 700, and 800 • C. After treatment, the phases of the powders are shown in Figure 3a. At 500 • C, TiH 2 decomposed partially, and the main phases were TiH 1.5 and Al. There was no obvious diffraction peak of intermetallic compound. At 600 • C, Al had been completely consumed, and the main phases were γ-TiAl, Ti 2 Al and a small amount of α2-Ti 3 Al phase. At 700 • C, the Ti 2 Al phase content decreased, while the γ-TiAl and α2-Ti 3 Al phase content increased. Upon increasing to 800 • C, the reaction progressed completely, and only γ-TiAl and α2-Ti 3 Al phases existed. enriched with Ti and leant with Al. Based on XRD analysis, the outer layer was γphase, while the core was Ti3Al phase. Moreover, multiple intermetallic layers mig formed in the PA powder, and diffusion between layers of various intermetallic pounds continued in the sintering process [35]. It was beneficial to achieve pressure sintering of the prepared PA powder since the diffusion between the intermetallic pound layers had little effect on the volume change. Besides, the homogenization o alloy would be completed in the sintering process. According to the Ti-Al phase diagram (shown in Figure 4) [4], the Ti-48Al alloy s to solidify at 1456 °C. With cooling (top→bottom analysis), the composition of Tialloy gradually narrows. The mechanism is the thermodynamic feasibility of the me In addition, Al powder was involved in the reaction as Equation (2). The process of forming TiAl 3 was an exothermic reaction with great volume expansion because of the Kirkendall effect, and it was difficult to achieve densification without pressure [19]. It can be inferred that the PA powder without Al was necessary for pressure-less sintering, and thus, the reaction temperature of 800 • C was suitable for PA powder preparation.
SEM image of the pretreated PA powder at 800 • C followed by crushing is shown in Figure 3b. It was found that the prepared PA powder was irregular. Figure 3c shows the cross-section and composition distribution of the PA powder. The PA powder was of "core-shell" structure, as shown in Figure 3c. According to the EDS results, the atomic ratio of aluminum and titanium in the outer layer was close to 1:1, while the core was enriched with Ti and leant with Al. Based on XRD analysis, the outer layer was γ-TiAl phase, while the core was Ti 3 Al phase. Moreover, multiple intermetallic layers might be formed in the PA powder, and diffusion between layers of various intermetallic compounds continued in the sintering process [35]. It was beneficial to achieve pressure-less sintering of the prepared PA powder since the diffusion between the intermetallic compound layers had little effect on the volume change. Besides, the homogenization of this alloy would be completed in the sintering process.
According to the Ti-Al phase diagram (shown in Figure 4) [4], the Ti-48Al alloy starts to solidify at 1456 • C. With cooling (top→bottom analysis), the composition of Ti-48Al alloy gradually narrows. The mechanism is the thermodynamic feasibility of the melting and solidification route [36,37]. The phase diagram confirms the stability of γ-TiAl at room temperature. Although the powder metallurgy route for producing Ti-Al alloy is not shown in the equilibrium phase diagram, the mixed powder of titanium and aluminum alloy inherits the bottom→top direction. and solidification route [36,37]. The phase diagram confirms the stability of γ-TiAl at room temperature. Although the powder metallurgy route for producing Ti-Al alloy is not shown in the equilibrium phase diagram, the mixed powder of titanium and aluminum alloy inherits the bottom→top direction.

γ-TiAl Alloy Preparation
In addition to strict process control, the oxygen content was controlled by adjusting the particle size of Al powder. Table 2 shows the O, N, H contents of sintered samples at 1350 °C prepared with different Al powder particle sizes. As the particle size of the Al powder increased, the O content decreased. The sample prepared using coarse Al powder of 120 μm as raw material had the lowest oxygen content, about 1710 ppm. The specific surface area of coarse Al powder decreased, resulting in the absorbed oxygen content decreasing. There was an aluminum oxide film on the surface of Al powder, and the oxygen atom would dissolve into the matrix during the powder preparation and sintering process, which was difficult to eliminate. In order to prepare γ-TiAl alloy with low oxygen content and high relative density, TiH2 powder and coarse Al powder (~120 μm) were selected as final raw materials. Besides, all experimental operations were carried out under argon atmosphere. The PA powder was prepared at 800 °C and then compacted using cold isostatic pressing. Green body was vacuum pressure-less sintered at different temperatures. Finally, the residual pores were eliminated by HIP without steel capsule.
Sintering temperature was one of the important factors that affect the density. Relative densities of Ti-48Al alloy sintered at different sintering temperatures (1300-1450 °C) are shown in Figure 5. With the sintering temperature increasing, the relative density showed an upward trend, and the increasing amplitude gradually decreased. Compared to the T1400 sample, the relative density of T1450 had no obvious change. Therefore, when the sintering temperature was higher than 1400 °C, the temperature had less influence on the relative density.

γ-TiAl Alloy Preparation
In addition to strict process control, the oxygen content was controlled by adjusting the particle size of Al powder. Table 2 shows the O, N, H contents of sintered samples at 1350 • C prepared with different Al powder particle sizes. As the particle size of the Al powder increased, the O content decreased. The sample prepared using coarse Al powder of 120 µm as raw material had the lowest oxygen content, about 1710 ppm. The specific surface area of coarse Al powder decreased, resulting in the absorbed oxygen content decreasing. There was an aluminum oxide film on the surface of Al powder, and the oxygen atom would dissolve into the matrix during the powder preparation and sintering process, which was difficult to eliminate. In order to prepare γ-TiAl alloy with low oxygen content and high relative density, TiH 2 powder and coarse Al powder (~120 µm) were selected as final raw materials. Besides, all experimental operations were carried out under argon atmosphere. The PA powder was prepared at 800 • C and then compacted using cold isostatic pressing. Green body was vacuum pressure-less sintered at different temperatures. Finally, the residual pores were eliminated by HIP without steel capsule.
Sintering temperature was one of the important factors that affect the density. Relative densities of Ti-48Al alloy sintered at different sintering temperatures (1300-1450 • C) are shown in Figure 5. With the sintering temperature increasing, the relative density showed an upward trend, and the increasing amplitude gradually decreased. Compared to the T1400 sample, the relative density of T1450 had no obvious change. Therefore, when the sintering temperature was higher than 1400 • C, the temperature had less influence on the relative density.
x FOR PEER REVIEW In general, the sample prepared by HIP is into a capsule welded by steel or titanium p cessing is required to remove the outer capsul and high cost. For this reason, HIP without a ca porosity. The pores in the sintered samples inc the porosity was less than 5%, the proportion Therefore, the sintered T1350 and T1400 sampl selected for non-capsule HIP. The densities of Figure 6 marked as the "five-pointed stars". T reached to 99.8%.

Phase and Microstructure
The microstructures of the samples, sinte Figure 6. There were a large number of pores o In general, the sample prepared by HIP is developed to pack powder or green body into a capsule welded by steel or titanium plates [11,17]. Subsequent machining processing is required to remove the outer capsule, which is complicated with long process and high cost. For this reason, HIP without a capsule was employed to reduce the residual porosity. The pores in the sintered samples included open pores and closed pores. When the porosity was less than 5%, the proportion of open pores would be greatly reduced. Therefore, the sintered T1350 and T1400 samples with relative density of above 95% were selected for non-capsule HIP. The densities of the HIP processed samples are shown in Figure 6 marked as the "five-pointed stars". The relative density of H1350 and H1400 reached to 99.8%. Table 3. Chemical composition of different regions marked in Figure 6 obtained by EDS analysis.   Table 3. Table 3. Chemical composition of different regions marked in Figure 6  Accordingly, the element distribution of H1350 sample was analyzed (Figure the γ phase and α2 phase, the distribution of Ti and Al was relatively uniform. It co proven that the problem of multilayer intermetallic compounds in the prepared PA der (as seen in Figure 3c) had been solved.  Table 3.

Phase and Microstructure
The microstructures of the samples, sintered at different conditions, are shown in Figure 6. There were a large number of pores on the prior particle boundaries in the sample sintered at 1300 • C, which was in good agreement with the relative density result. As the sintering temperature increased, the pores gradually decreased. At 1450 • C, only a few pores remained in the sample.
Combined the EDS results shown in Table 3 with the XRD results, it can be known that the gray area in Figure 6 was the γ-TiAl phase and the bright area was the α2-Ti 3 Al phase. The microstructure of the T1350 sample was near-γ. Near-γ was composed of equiaxed γ particles and a small number of α2 particles distributed in the γ grain boundaries [35][36][37]. The microstructure of T1400 sample was duplex, as shown in Figure 6c. The duplex structure was composed of equiaxed γ grains and γ/α2 layers, in which the α2 phase was precipitated from the disordered α phase existing at high temperature. The grain size of duplex was generally small because the α phase and γ phase were pinned to each other at high temperatures, resulting in slow grain growth. Moreover, the dispersed α2 phase in duplex would increase the strength with the sacrifice of plasticity [2,38]. The microstructure of the T1450 sample was nearly lamellar, as shown in Figure 6d. Nearly lamellar was composed of γ/α2 layers and a small number of equiaxed γ grains distributed between the layer clusters. As the γ phase was reduced, the pinning effect on the α phase Metals 2021, 11, 635 9 of 13 was weakened, resulting in larger lamellae. This kind of microstructures had high fracture toughness and creep resistance, but also showed low ductility at room temperature [2].
After HIP, the corresponding microstructures with high relatively density were observed. Comparing Figure 6b,c,e,f, it was found that HIP did not change the type of microstructure, and the pores reduced significantly. Although dense microstructure was observed in the HIP samples, grain growth was also inevitable under this HIP condition. Compared to Figure 6b,c, coarse microstructure was obviously found in Figure 6e,f. Accordingly, the element distribution of H1350 sample was analyzed (Figure 7). In the γ phase and α2 phase, the distribution of Ti and Al was relatively uniform. It could be proven that the problem of multilayer intermetallic compounds in the prepared PA powder (as seen in Figure 3c) had been solved.

Mechanical Properties
The relationship between Vickers hardness and sintering temperatu Figure 8. The hardness value of as-sintered compact increased with the inc tering temperature, but decreased after 1400 °C. The reason for the hardn was that the relative density increased and the porosity decreased with th increasing at 1300-1400 °C. Upon a further increase in the temperature t hardness value decreased due to the coarse grains and the transformed duplex to sub-lamellar structure. High sintering temperature promoted unn growth, which was harmful to mechanical properties. Therefore, 1350 °C were selected as appropriate sintering temperatures for PM Ti-48Al allo with as-sintered samples, the HIP samples had lower hardness, as shown Under the action of high temperature (1200 °C) and high pressure (140 M significantly grew and the structure became coarser.

Mechanical Properties
The relationship between Vickers hardness and sintering temperature is shown in Figure 8. The hardness value of as-sintered compact increased with the increasing of sintering temperature, but decreased after 1400 • C. The reason for the hardness increasing was that the relative density increased and the porosity decreased with the temperature increasing at 1300-1400 • C. Upon a further increase in the temperature to 1450 • C, the hardness value decreased due to the coarse grains and the transformed structure from duplex to sub-lamellar structure. High sintering temperature promoted unnecessary grain growth, which was harmful to mechanical properties. Therefore, 1350 • C and 1400 • C were selected as appropriate sintering temperatures for PM Ti-48Al alloys. Compared with as-sintered samples, the HIP samples had lower hardness, as shown in Figure 8a. Under the action of high temperature (1200 • C) and high pressure (140 MPa), the grains significantly grew and the structure became coarser. duplex to sub-lamellar structure. High sintering temperature promoted unnecessa growth, which was harmful to mechanical properties. Therefore, 1350 °C and were selected as appropriate sintering temperatures for PM Ti-48Al alloys. Co with as-sintered samples, the HIP samples had lower hardness, as shown in Fi Under the action of high temperature (1200 °C) and high pressure (140 MPa), th significantly grew and the structure became coarser.  Figure 7b presents the compression properties of these two HIP samples. Besides, Table 4 compares the compressive strengths of the HIP Ti-48Al alloys and some alloys reported in related literatures [39][40][41]. The compressive strength of H1300 sample was 2241 MPa, while that of the H1400 sample was 1931 MPa. Compared with the as-cast Ti-48Al-2Nb-2Cr alloy [40], the compressive strength of the prepared Ti-48Al alloy increased by 130% although there were no two alloying elements (Nb and Cr), which proved that the prepared PM Ti-Al alloys, using the new kind of PA powder, had higher properties. Furthermore, TiAl alloys prepared by traditional casting method would form coarse columnar crystals during the cooling process, because of the poor thermal conductivity of TiAl alloys. With the powder metallurgy process of γ-TiAl alloy, a very fine and chemically uniform microstructure could be obtained, which was beneficial for mechanical properties. In addition, compared with the Ti-50Al alloy prepared by hot pressing [39], the prepared Ti-48Al alloy had a 72% increase in the compressive strength.    Table 4 compares the compressive strengths of the HIP Ti-48Al alloys and some alloys reported in related literatures [39][40][41]. The compressive strength of H1300 sample was 2241 MPa, while that of the H1400 sample was 1931 MPa. Compared with the as-cast Ti-48Al-2Nb-2Cr alloy [40], the compressive strength of the prepared Ti-48Al alloy increased by 130% although there were no two alloying elements (Nb and Cr), which proved that the prepared PM Ti-Al alloys, using the new kind of PA powder, had higher properties. Furthermore, TiAl alloys prepared by traditional casting method would form coarse columnar crystals during the cooling process, because of the poor thermal conductivity of TiAl alloys. With the powder metallurgy process of γ-TiAl alloy, a very fine and chemically uniform microstructure could be obtained, which was beneficial for mechanical properties. In addition, compared with the Ti-50Al alloy prepared by hot pressing [39], the prepared Ti-48Al alloy had a 72% increase in the compressive strength. The HIP samples were subjected to a tensile test at room temperature. The tensile strength of the H1350 sample was 333 MPa, while that of the H 1400 sample was 350 MPa. Correspondingly, the tensile fracture morphology was observed, as shown in Figure 9. These two alloys had good densification state with fewer micro-cracks. The fracture was a brittle fracture. The main fracture modes in Figure 9a were consisted with a transcrystalline fracture and intergranular fracture. Besides, the fracture modes in Figure 8b included transcrystalline, intercrystalline, translamellar and interlamellar fractures. The fracture modes of these two alloys were different because the structures were different, as shown in Figure 6e,f. Compared with the H1350 sample with near-γ microstructure, the H1400 sample had duplex microstructure with near-γ and lamellar structure, and thus the fracture appeared as translayer fracture and interlayer fracture. In addition, the cleavage facets were smaller in Figure 9a. This was because the grains became coarser with the sintering temperature increasing. a brittle fracture. The main fracture modes in Figure 9a were consisted with a transcrystalline fracture and intergranular fracture. Besides, the fracture modes in Figure 8b included transcrystalline, intercrystalline, translamellar and interlamellar fractures. The fracture modes of these two alloys were different because the structures were different, as shown in Figures 6e,f. Compared with the H1350 sample with near-γ microstructure, the H1400 sample had duplex microstructure with near-γ and lamellar structure, and thus the fracture appeared as translayer fracture and interlayer fracture. In addition, the cleavage facets were smaller in Figure 9a. This was because the grains became coarser with the sintering temperature increasing.

Conclusions
A new kind of non-spherical PA TiAl powder was prepared by the reaction of TiH2 with Al at high temperature, and then Ti-48Al alloy with high relative density and low oxygen content was prepared by cold isostatic pressing and vacuum pressure-less sintering. These conclusions were as follows: (1) Using the blended TiH2 and Al powders as raw materials, a non-spherical TiAl pre-alloy powders were prepared by reacting at high temperature followed by crushing. There were only γ-TiAl and α2-Ti3Al phases in the PA powder pretreated at 800 °C.
(2) The irregular pre-alloyed powder had good sinterability. After pressure-less sintering at 1350 °C and 1400 °C temperatures, the relative densities were 95.2% and 97.3%, respectively. Followed by non-capsule HIP, both relative densities increased to nearly 100%.
(3) The particle size of the raw Al powder significantly affected the oxygen content

Conclusions
A new kind of non-spherical PA TiAl powder was prepared by the reaction of TiH 2 with Al at high temperature, and then Ti-48Al alloy with high relative density and low oxygen content was prepared by cold isostatic pressing and vacuum pressure-less sintering. These conclusions were as follows: (1) Using the blended TiH 2 and Al powders as raw materials, a non-spherical TiAl pre-alloy powders were prepared by reacting at high temperature followed by crushing. There were only γ-TiAl and α2-Ti 3 Al phases in the PA powder pretreated at 800 • C.
(2) The irregular pre-alloyed powder had good sinterability. After pressure-less sintering at 1350 • C and 1400 • C temperatures, the relative densities were 95.2% and 97.3%, respectively. Followed by non-capsule HIP, both relative densities increased to nearly 100%.
(3) The particle size of the raw Al powder significantly affected the oxygen content of the sample. The larger the Al particle size, the smaller the O content. The sample prepared with coarse Al powder of 120 µm had a low oxygen content of 1710 ppm.
(4) The as-sintered samples had a near-γ structure at 1350 • C a duplex structure at 1400 • C, and a nearly lamellar structure at 1450 • C. Besides, the sample sintered at 1350 • C followed by HIP had comprehensive properties with compressive strength of 2241 MPa and tensile strength of 333 MPa.