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A novel powder metallurgy processing approach to prepare fine-grained Ti-rich TiAl-based alloys from pre-alloyed powders
Intermetallics 42 (2013) 146e155
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
A novel powder metallurgy processing approach to prepare fine-grained Ti-rich TiAl-based alloys from pre-alloyed powders Sanjay K. Vajpai a, *, Kei Ameyama b,1 a b
Department of Mechanical Engineering, Ritsumeikan University, 1-1-1, Noji-higashi, Kusatsu, Shiga 525-8577, Japan Faculty of Science and Engineering, Ritsumeikan University, 1-1-1, Noji-higashi, Kusatsu, Shiga 525-8577, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 9 April 2013 Received in revised form 21 June 2013 Accepted 25 June 2013 Available online 13 July 2013
A relatively simple yet attractive and efficient powder metallurgy processing method, based on mechanical milling and spark plasma sintering of PREP (Plasma Rotating Electrode Processed) pre-alloyed powders, has been proposed to prepare full density fine-grained Ti-rich TieAl alloys with controlled microstructure. Full density fine-grained Ti-rich TieAl alloy compacts, with grain size in the range of 2e4 mm, were successfully prepared from dendritic coarse PREP pre-alloyed powders, which are otherwise difficult to consolidate, by mechanical milling followed by spark plasma sintering at moderate consolidation conditions. It has been demonstrated that the microstructure, i.e. lamellar or equiaxed, depends on the extent of mechanical milling. The fine-grained TieAl compacts exhibited extremely good mechanical properties. An attempt has been made to correlate the effect of mechanical milling on the microstructure and properties of the sintered TieAl compacts. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl B. Mechanical properties at ambient temperature C. Mechanical alloying and milling D. Powder metallurgy, including consolidation E. Sintering
1. Introduction In recent decades, Titanium-rich TiAl-based alloys have been considered a promising candidate as a high temperature structural material in automobile and aerospace sectors, due to their better high temperature specific strength than alloy steels and Nibased super-alloys, together with good oxidation, corrosion, and creep resistance at elevated temperatures [1e9]. Moreover, these alloys have been also found suitable as nuclear structural material due to their excellent radiation resistance and low neutron activation, together with good high-temperature creep resistance [10e13]. However, these alloys suffer from the inherent room temperature brittleness, which not only limits their commercial applications but also poses severe restrictions on the processing due to limited room temperature formability. The limited commercial application of these alloys can be attributed primarily to the coarse grained microstructure of the components fabricated
* Corresponding author. Tel.: þ81 77 561 2749; fax: þ81 77 561 2665. E-mail addresses: [email protected] (S.K. Vajpai), [email protected] (K. Ameyama). 1 Tel.: þ81 77 561 2756; fax: þ81 77 561 2665. 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.06.006
by the conventional ingot metallurgy processing, wherein the coarse-grained microstructure further aggravates the problem of brittleness and formability. It would be worth mentioning that the mechanical properties of the TiAl alloys can be improved by not only refining grain size but also controlling the amount and distribution of different phases in the microstructure, such as a2Ti3Al and g-TiAl phases in the form of either duplex or lamellar structure [14e20]. Although the desired phases can be obtained by controlling the composition and thermo-mechanical treatments [20e24], obtaining a fine-grained microstructure with a controlled distribution of various phases is extremely difficult. Therefore, the fabrication of near-net shape products with controlled microstructure, which governs the mechanical properties, via conventional ingot metallurgy processing remains a challenge due to limited room temperature formability and requirements of complex thermo-mechanical treatments [5e7]. The mechanical properties of TiAl based alloys depend on several factors, including purity of alloy, grain size, phase constitution, and distribution of various phases in the finished product [1e9,14e24]. Therefore, a microstructural design must include all the above mentioned factors to achieve optimum mechanical properties. Moreover, it is equally important to design a processing strategy which is not only capable of delivering near-net shape component with desired microstructure but also technologically
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viable and cost effective [24]. Powder metallurgy (PM) processing is an attractive way to achieve all the above mentioned goals. PM processing makes it possible to prepare near-net shape products, which minimizes the requirements of post processing machining [25]. Moreover, it also provides an opportunity to prepare finegrained components together with better control over chemical and phase compositions. There have been some efforts to prepare TiAl based alloys via powder metallurgy route [26e42]. The PM routes based on the mechanical milling of elemental powders suffer from the problems of contamination due to severe milling conditions, low yield due to sticking of elemental powders to the milling media, difficulty in handling the milled powders during subsequent processing due to their high reactivity, and limitations in achieving chemical homogeneity in the sintered components [26e32]. Moreover, controlling the microstructure is still an issue due to the formation of several undesirable phases, which have extremely adverse effect on the mechanical properties. The aforementioned problems associated with the PM processing can be effectively overcome by using atomized pre-alloyed powders as starting material. The rapid solidification during atomization ensures better control over chemical composition and chemical homogeneity, together with fine-grained microstructure, in the pre-alloyed powders. These characteristics provide an opportunity to control the microstructure in the finished components only by proper adjustments of the conditions during subsequent consolidation stage. In recent years, a few efforts have been made to prepare fine-grained TiAl-based alloys by spark plasma sintering of inert gas atomized pre-alloyed powders [33e37]. The sintered nearfull density compacts, thus prepared, exhibited improved mechanical properties as compared to their counterparts prepared by conventional ingot metallurgy method. However, it was found that the final density, microstructure, and properties of these compacts depend on the starting powder characteristics and sintering conditions. These parameters include average powder particle size, powder size distribution, initial composition, microstructure of the powder, sintering temperature, sintering pressure, and total sintering time. Therefore, the PM processing based on spark plasma sintering of pre-alloyed powders can become a viable method to prepare TiAl-based alloys with desired microstructure and properties, provided the correlation amongst the above parameters can be understood more clearly. Plasma rotating electrode process (PREP) is an attractive rapid solidification process which is capable of producing high purity pre-alloyed powders on an industrial scale. However, only a few efforts were made to prepare Ti-rich TiAl based alloys using PREPed powder as starting material, owing to the issues related with the consolidation of the PREPed powder with controlled microstructure [38e42]. It was found that the PREPed TiAl powder mass primarily consists of two different types of powder particles, i.e. (i) fine martensitic powder particles with size less than 100 mm, and (ii) coarse dendritic powder particles with size over 100 mm. It was also demonstrated that the fine martensitic particles are easy to consolidate to near full density, whereas coarse dendritic powders are very difficult to consolidate. Generally, a combination of high pressure, high temperature, and long pressing time is required to achieve near full density. However, the requirements of these consolidation conditions restrict the control over the final microstructure in the sintered compacts. Furthermore, the PREPed powder mass contains coarse dendritic powders in the range of 60e80 mass percent, which severely restricts the commercial viability of using PREPed powders as starting material. Therefore, it is desired to develop a PM processing methodology based on PREP powders wherein near-net shape full density components can be prepared under moderate consolidation conditions together with better control over microstructure.
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In the present work, a new powder metallurgy processing approach has been proposed to prepare full density compacts, with controlled microstructure, under moderate consolidation conditions using “difficult-to-consolidate” coarse dendritic PREPed TiAl pre-alloyed powders as starting material. The proposed PM route primarily consists of mechanical milling of coarse dendritic powders followed by spark plasma sintering. It is envisaged that the mechanical milling of the prealloyed powder would enhance the sintering kinetics, resulting in near full density compacts at relatively lower sintering temperature and pressure in short period of time. In order to understand the effect of mechanical milling on the density and microstructure of the sintered compacts, PREP powders were milled for different time period. The microstructural characteristics of the starting PREPed powder and mechanically milled powders are presented to understand the effect of milling on the microstructure and morphology of prealloyed powders. Subsequently, the microstructural characteristics of the sintered compacts prepared from mechanically milled powders are also presented to elaborate the effect of milling time on the densification and microstructural evolution in the sintered compacts. Finally, in order to assess the mechanical properties of the sintered compacts, the results of the hardness measurements and flexural tests are presented and discussed. 2. Experimental procedure Ti-rich TieAl pre-alloyed powders, with particle size in the range of 200e1400 mm, produced by helium-PREP were used in the present study. The chemical composition of the powders is shown in the Table 1. The pre-alloyed powder was mechanically milled (MM) in a planetary ball mill up to 108 ks (30 h). To minimize the contamination due to mechanical milling, the milling was carried out under inert gas (argon) atmosphere and low ball-to-powder ratio (2:1). Subsequently, the powders milled for 18 ks (5 h) and 108 ks (30 h) were sintered by spark plasma sintering at 1323 K for 3.6 ks (1 h) under the pressure of 50 MPa. The powder was filled in the graphite die and the pressure was applied during sintering. Henceforth, the sintered compacts prepared from initial PREP, MMed for 5 h, and MMed for 30 h powders will be referred as MM0h, MM5h, and MM30h, respectively. The average particle size and size distribution of the initial and milled powders were evaluated by laser particle size analyzer. X-ray diffraction and scanning electron microscopy (SEM) techniques were used for phase identification and microstructural characteristics of the starting powder, milled powders, and sintered compacts. In order to achieve the phase contrast in the micrographs, images were acquired using backscattered electron detector in the SEM. Four point bend test and micro-hardness tests were carried out to evaluate the mechanical behaviour of the sintered compacts. The micro-hardness tests were carried out under the load of 980.7 mN, and average hardness was calculated using 25 measurements. The four point bend test was carried out on rectangular specimens with approximate dimensions 1.5 2.5 15 (all dimensions in mm). The rectangular bend test specimens were polished using 0.1 mm alumina suspension to minimize the roughness on the surface. Finally, fractured surfaces of the specimens were observed under the SEM to understand the nature of fracture of the specimens.
Table 1 Chemical composition of the starting PREPed powder. Element
Al
O
Fe
Si
Ti
Amount (mass%)
33.67
0.387
0.02
0.07
Bal.
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3. Results and discussion 3.1. Microstructural characteristics of initial PREP powder Fig. 1 shows the morphology and particle size distribution of the PREPed TiAl powders. It can be seen that the powder particles are spherical in shape, and have wide and bimodal size distribution. The particle size varied in the range of 180e1400 mm, and the average particle size was approximately 500 mm. It can be noted that most of the particles are large in size. The bimodal particle size distribution is a common feature of PREPed powders. The typical surface morphology of the PREPed powders is shown in Fig. 2. It was noted that most of the particles consisted of dendritic structure. The PREP TiAl powders have been found to consist of either martensitic or dendritic microstructure depending on the cooling rate during processing [39e42]. The cooling rate also depends on the particle size, and it has been observed frequently that the small sized particles (<100 mm size) exhibit martensitic structure, whereas larger particles exhibit dendritic microstructure. Since the present PREP TiAl powders are very large in size, it is obvious that the cooling rate during flight would be “relatively” slower as compared to small sized martensitic particles, leading to the formation of dendritic microstructure.
Fig. 2. SEM micrographs of PREPed powders showing the dendritic surface morphology.
Fig. 3 shows the XRD pattern of the PREPed powders. It can be seen that hexagonal a2 (Ti3Al) is the predominant phase present in the particles together with a small amount of tetragonal g (TiAl) phase. The presence of a2 as predominant phase suggests that the hexagonal a phase is the first phase which was formed during solidification. The presence of ordered a2 hexagonal phase together with a precipitation of g phase in small amounts also suggests that
Fig. 1. (a) Morphology and (b) particle size distribution of PREPed TieAl powder.
Fig. 3. XRD pattern of the initial PREPed powder.
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the solidification rate remains sufficiently high to retain the hexagonal a phase in the solidified powder particles. However, the ordering of a phase is not suppressed even at such a rapid cooling rate, leading to the formation of the ordered variant of a, i.e. hexagonal a2. Moreover, the presence of ordered a2 phase in large amounts also suggests that the rapid solidification during preparation of PREP TiAl powder also leads to the retention of supersaturated a2 phase. The presence of dendritic microstructure, accompanied with the formation of hexagonal a2 phase, has been found to be a typical feature of the coarse particles of PREP TiAl powders [39e42]. 3.2. Microstructural characteristics of mechanically milled powders The PREPed TiAl powders were mechanically milled in a planetary ball mill for different period of time. Fig. 4 shows the average particle size and size distribution of the MMed powders together with initial powder. It can be clearly observed that the average particle size reduced very rapidly within initial 5 h of milling followed by a sluggish refinement in particle size (Fig. 4a). The particle size after initial 5 h of milling was found to be approximately 50 mm. Subsequently, the average particle size of approximately 10 mm was obtained after 15 h of milling. Furthermore, it can be noticed that further milling under the present conditions did not result in any further refinement in particle size even after prolong milling time, i.e. 30 h, and the average particle size remained approximately 10 mm. It would be worth mentioning that the mechanical milling resulted in not only a small average particle size
Fig. 4. Effect of mechanical milling on the (a) average particle size, and (b) size distribution of PREPed powders.
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but also reduced size distribution (Fig. 4b). Fig. 5 shows the effect of milling time on the morphology of the PREPed powder. It can be seen that after initial 5 h of milling, initial PREP powder particles fragmented in to smaller size irregular shaped particles, primarily via brittle mode of fracture (Fig. 5a and b). Although the initial milling period appears to be dominated by the brittle fracture of initial PREP powders, it would be worth mentioning that the fractured powder particles would also undergo to some degree of plastic deformation during this stage. Fig. 5c and d shows the morphology of the PREP powders milled for 15 h. A comparison of Fig. 5a and c clearly shows that the increased milling time led to the further fragmentation of the powder particles, resulting in the refinement of the particle size. Moreover, it can also be observed that the resulting powder particles are slightly flat and irregularshaped, suggesting that the particles appear to have undergone a significant plastic deformation. Fig. 5e and f shows the morphology of the PREP powders milled for 15 h. A comparison of Fig. 5c and e indicates that the subsequent milling after 15 h did not induced any significant fragmentation of the powder particles, as demonstrated by particle size analysis earlier and shown in Fig. 4. However, the flaky morphology of the milled powder particles, as shown in Fig. 5f, is a clear evidence of severe plastic deformation of the powders due to milling. Fig. 6 shows the XRD pattern of initial and MMed powders. It can be seen that there is neither any shifting in the positions of peaks nor any change in the relative intensity of g-peaks with respect to a2 peaks. These results indicate that the mechanical milling did not induce any phase change in the PREP TiAl powder. However, it can be noted that there was a significant broadening in the XRD peaks as a result of milling. Fig. 7 shows the variation in the width of the three most intense peaks with increasing milling time, wherein full width at half maxima (FWHM) of the peak has been considered as the peak width. It can be observed that the peak width increased rapidly up to approximately 5 h of milling. The subsequent milling led to a moderate increment in peak broadening between 5 and 15 h followed by a sluggish increment up to 30 h of milling. The important factors contributing to the XRD peak broadening during mechanical milling are (i) introduction of lattice strains caused by the accumulation of defects, such as dislocations and point defects, and (ii) crystallite size refinement. It is well established that large amounts of lattice defects are accumulated during the early stages of mechanical milling followed by achieving a saturation level, wherein the rate of generation and annihilation of these defects attain equilibrium state. The attainment of the equilibrium state is characterized by the formation of very fine size crystallites in the milled powders. Thus, the XRD peak broadening during early stages of milling is primarily related with the accumulation of large amounts of lattice defects, whereas the peak broadening after prolonged milling can be attributed to the formation of extremely small crystallites. Therefore, in the present case, it appears that the early stages of milling, up to 15 h, involves fracturing of powder particles together with an accumulation of large amounts of lattice defects due to plastic deformation of fragmented particles. However, the subsequent milling leads to the severe plastic deformation of fragmented particles, as evidenced by the flaky morphology of the 30 h milled powder, resulting in the powder particles with extremely fine-sized crystallites. It is envisaged that the mechanically milled irregular-shaped small-sized powder particles would exhibit better packing, in terms of interparticle contact area, as compared to the large-sized spherical PREP powders. Moreover, the accumulated defects and fine-grained structure in the milled powder particles would enhance the diffusion kinetics during subsequent sintering, resulting in the improved sinterability and chemical homogeneity. As a result, it would be possible to prepare near full density
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Fig. 5. SEM micrographs of PREPed powders milled for (a, b) 5h, (c, d) 15 h, and (e, f) 30 h.
Fig. 6. XRD patterns of the MMed PREP TiAl powders.
Fig. 7. Variation in the XRD peak width with increasing milling time.
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Fig. 8. Typical micrographs showing the thickness cross-section of the as-polished sintered compacts prepared from (a) initial PREP powder, and (b) MMed PREP powder.
products with improved chemical homogeneity in relatively short processing time under the low pressure and low temperature processing conditions. In the present work, powders milled for 5 h (short-time milled) and 30 h (prolonged milled) were only processed further to understand the effect of two extreme milling conditions on the microstructure and properties of the resulting TiAl compacts.
followed by sintering under low pressure, low temperature, and relatively shorter sintering time. Fig. 9 shows the XRD results of the sintered compacts prepared from initial as well as mechanically milled PREP powders. It is interesting to note that the sintered compacts consist of primarily tetragonal g phase (TiAl) together with a small amount of hexagonal a2 (Ti3Al), which is altogether different from the as PREPed powder wherein hexagonal a2 (Ti3Al) was the primary phase present. These results suggest that the initial PREP powder consisted of supersaturated hexagonal a2 (Ti3Al) phase, as a result of rapid cooling. Moreover, the presence of tetragonal g phase (TiAl) as the major phase in the sintered compacts is not unexpected considering the composition of the initial powder. It appears that the sintering has led to the formation of large amounts of hexagonal g phase (TiAl) via precipitation from supersaturated hexagonal a2 (Ti3Al) phase, which is present in the rapidly solidified initial PREP TiAl powder. Fig. 10a and b shows the SEM micrographs depicting the microstructures of sintered compacts prepared from the initial PREP TiAl powder. It can be seen that the most of the powder particles in the sintered compacts consisted of diffused dendritic microstructure (Fig. 10a). Although some recrystallized equiaxed grains, g as well as a2/g lamellar microstructure were also found to be present in the compacts (Fig. 10b), only a few such areas were found to exist. Moreover, it can also be noticed that a2 phase is present either on the grain boundaries or as a layer in the lamellar structure. These observations are in conformity with the XRD results, showing that the initial PREP powders consisted of supersaturated hexagonal a2 (Ti3Al) phase and its dissociation during sintering leads to the formation of ordered Ti3Al and TiAl phases. Fig. 10c shows a typical SEM micrograph depicting the microstructure of sintered compacts prepared from 5 h milled PREP TieAl alloy powders. It can be clearly observed that the sintered compacts consist of primarily lamellar microstructure, more or less equiaxed lamella colonies consisting of alternate layers of a2/g phases, together with equiaxed fine size g grains. The presence of equiaxed grains/lamella colonies indicates that the microstructural evolution during sintering was dominated by the recrystallization and grain growth. It can also be noted that the a2 phase is distributed in the matrix as either submicron-sized particles on the boundaries of equiaxed g grains or extremely fine sized layers, having less than 0.2 mm thickness, in the lamellar structure. Therefore, it is evident that a short time milling led to not only the disappearance of initial coarse dendritic microstructure but also the formation of fine grained microstructure. It has been already shown that the short time milling, i.e. 5 h in the present case, induces large amounts of
3.3. Microstructural characteristics of sintered compacts The initial PREP and MMed powders were sintered by spark plasma sintering at 1323 K for 60 min under 50 MPa pressure. Fig. 8 shows the micrographs depicting the as-polished thickness crosssection of the sintered compacts. Fig. 8a shows that the sintered compacts of initial PREP powders consist of significant amount of open porosity, i.e. approximately 25 volume%. It can be noticed that the spherical PREP powder particles only joined together by forming a neck as a result of sintering and did not undergo any significant plastic deformation during sintering, resulting in large amounts of retained porosity. Fig. 8b shows the thickness crosssection of the sintered compacts prepared from 5 h milled powder. The absence of any appreciable amount of residual porosity shows that a short-time milled irregular shaped powder particles resulted in the full density sintered compacts. Similarly, 30 h milled powder also resulted in full density sintered compacts. Therefore, it is evident that full density TiAl compacts can be successfully prepared from large sized PREP TiAl powders by mechanical milling
151
Fig. 9. XRD patterns of the sintered TiAl compacts.
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Fig. 10. Typical SEM micrographs of sintered TiAl compacts prepared from (a, b) initial PREP powder, (c) 5 h MMed powder, and (d) 30 h MMed powder.
microstructural defects in the powders. Furthermore, it is well known that these microstructural defects enhance the diffusion kinetics, and the diffusivity of atoms of various elements in an alloy becomes very high even at very low temperatures. Therefore, it appears that the presence of these microstructural defects, such as dislocations and point defects, in the short time milled PREP powders enhanced the diffusion process significantly during sintering, leading to the disappearance of coarse dendritic microstructure and formation of refined microstructure with a particular phase distribution. Since short time milling does not ensures uniform accumulated plastic deformation in the milled powders, the microstructural evolution during sintering appears to be dependent on the local accumulated plastic deformation in the short time milled powders. As a result, both a2/g lamellar and (a2 þ g) equiaxed microstructures are found in the sintered compacts prepared from 5 h MMed powders. The average grain size of the sintered compacts was found to be approximately 3.82 1.11 mm. Fig. 10d shows a typical SEM micrograph exhibiting the microstructure of the sintered compacts prepared from 30 h milled PREP TieAl alloy powders. It can be observed that sintered compacts primarily consist of very fine size equiaxed grains of g phase (tetragonal ordered TiAl), having average grain size 2.34 1 mm. It can also be noticed that there is no evidence of existence of lamellar microstructure in significant amounts and the a2 phase is present primarily on the boundaries of the equiaxed g grains in the form of very small particles with size less than 2 mm. It would be worth mention that a significant amount of a2 phase particles have size in the submicron range. These microstructural features indicate that the severely deformed supersaturated PREP powders undergo recrystallization during sintering to form strain free small grains, together with the simultaneous precipitation of small sized a2 on the grain boundaries of these newly formed grains. It appears that these extremely small sized a2 phase present on the grain boundaries restricted further growth of the recrystallized g grains during sintering, resulting in the fine-grained microstructure consisting of equiaxed g grains. Recently, Guyon et al. [37] have also reported the formation of submicron/nano size recrystallized grains during
spark plasma sintering of milled TieAl-based prealloyed powders prepared by inert gas atomization. Similar to the present work, fine size a2 was found to be present on the grain boundaries of the finesize g grains wherein it was postulated that these fine size a2 particles/grains restrict any significant growth of the recrystallized g grains during subsequent heating. Table 2 shows the relative amounts (volume%) of a2 and g phases in the sintered compacts prepared from initial and MMed powders. The average amounts were estimated from the SEM micrographs which were acquired using backscattered electron detector. It can be noted that all the sintered compacts consist of comparable relative amounts of a2 and g phases irrespective of the condition of the starting powder, i.e. as PREPed or MMed for different periods of time. Therefore, these results suggest that the extent of mechanical milling only brings about the redistribution of the a2 and g phases in the sintered compacts prepared from PREP powders, and their relative amounts is primarily governed by the chemical composition of the initial PREP powder. The above results show that the coarse PREP pre-alloyed TieAl powders are difficult to consolidate to near-full density under the moderate sintering conditions such as short sintering time, relatively lower sintering temperature, and low pressure. Moreover, the initial dendritic microstructure also remains retained under these sintering conditions. On the other hand, it has been demonstrated that mechanical milling of these “difficult-to-consolidate” coarse PREP powders leads to not only near-full density compacts with fine-grained microstructure but also the disappearance of dendritic microstructure. Ameyama et al. [27] demonstrated that fine grained Table 2 Average amounts of a2 and g phases present in the sintered compacts. Material
Average amount of a2 phase (vol.%)
Average amount of g phase (vol.%)
PREP þ SPS (1323 K) MMed (5 h) þ SPS (1323 K) MMed (30 h) þ SPS (1323 K)
26 23 20
74 77 80
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Fig. 11 shows the average micro-hardness values of the sintered specimens prepared from initial PREP, 5 h milled, and 30 h milled powders. It can be noticed that the average micro-hardness of the compacts prepared from short-time milled powders, i.e. 5 h milled, is only slightly higher than that of the compacts prepared from initial PREP powders. However, it is interesting to note that the average micro-hardness of the compacts prepared from 30 h milled powders is significantly higher than that of the compacts prepared from initial PREP and 5 h milled powders. Moreover, it can also be noticed that the dispersion of micro-hardness values from the mean value decreased with increasing milling time. It has been demonstrated earlier that the sintered compacts of initial PREP powder consist of diffused dendritic and lamellar microstructure, together with the precipitation of a2 particles of 2e5 mm size. On the other hand, short-time milling results in the formation of
lamellar microstructure due to the enhanced short-range diffusivity. It would be worth pointing out that both the microstructures are not significantly different in terms of distribution of a2 and g phases, resulting in comparable hardness values. It would also be worth mentioning that the distribution of a2 and g phases in the lamellar microstructure is relatively much finer as compared to that of diffused dendritic microstructure. This fact is reflected by the smaller dispersion of hardness values from the mean hardness value in the sintered compacts prepared from 5 h milled powder as compared to that of initial PREP powders. The significantly high average hardness of sintered compacts from 30 h milled powder can be attributed to the presence of fine-size equiaxed grains of g phase, which inherently exhibits higher brittleness as compared to the lamellar microstructure. Fig. 12 shows representative flexural strength vs. deflection curves of the specimens prepared from the MM5h and MM30h sintered compacts, having average fracture strength 978.41 85.31 MPa and 883.23 3.64 MPa, respectively. It can be noted that the sintered compacts of 5 h MMed powder exhibited slightly higher average fracture strength and deflection-to-fracture, together with lower relative hardness, than those of compacts of 30 h MMed powder. Since a combination of higher deflection-to-fracture and lower hardness values can be regarded as an indicator of better ductility in case of brittle materials, the above results show that the MM5h sintered compacts exhibit better combination of strength and ductility as compared to the MM30h compacts. It would be worth mention that the average fracture strength of the present TiAl compacts is considerably higher than coarse grained TiAl-based alloys and comparable to that of the fine-grained compacts prepared from inert gas atomized prealloyed powders [1,2,33,35,36]. The high fracture strength of the present compacts can be attributed to the refined microstructure. It can also be noted that the MM5h compacts have not only the higher average fracture strength but also larger scattering in the strength values as compared to MM30h compacts. However, the scattering in the fracture strength values is considerably small as compared to those reported in the literature for TiAl-based alloys. Fig. 13 shows the morphology of the fractured surfaces of the flexural test specimens prepared from the MM5h and MM30h sintered compacts. It can be clearly observed that all the sintered compacts exhibit brittle nature of fracture. However, it can be noted that the mode of fracture is different for both the compacts, i.e. the deformation mode of MM5h compacts is dominated by trans-
Fig. 11. Micro-hardness of the sintered compacts prepared from initial PREP and MMed powders.
Fig. 12. Representative Flexural strengthedeflection curves of sintered compacts prepared from mechanically milled powders.
TiAl-based alloys undergo superplastic deformation even at lower temperatures. In the present work, it has been demonstrated that the spark plasma sintering of milled powder resulted in the formation of fine-grained recrystallized microstructure. Therefore, the enhanced densification of the milled PREP pre-alloyed powders can be attributed to the superplastic deformation of the milled power particles even at lower sintering temperature. Similar observations regarding the enhanced densification of milled inert-gas-atomized pre-alloyed TiAl-based powder at lower temperature have also been made by Guyon et al. [37]. Furthermore, it is also demonstrated that the final microstructure of the sintered compacts also depends on the severity of plastic deformation induced in the powders by mechanical milling. Therefore, the present processing route provides an opportunity to prepare full density compacts with desired microstructure by adjusting the milling and sintering conditions. It would be worth mentioning that near-full density compacts can be prepared from coarse PREP powders by applying high temperature and high pressure for relatively longer sintering time, which is typically 4e5 h [39,41,42]. However, such processing conditions not only limit the bulk production of near-net shaped full density products but also result in the coarse-grained microstructure. Moreover, it is also extremely difficult to achieve the desired microstructure with controlled distribution of various phases in the microstructure when coarse PREP powders are used as starting material. 3.4. Mechanical properties of the sintered compacts
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Fig. 13. Fractured surfaces of the flexural test specimens prepared from the sintered compacts of (a, b) 5 h, and (c, d) 30 h MMed PREP powders.
granular mode of fracture (Fig. 13a and b) whereas it is primarily intergranular in case of MM30h compacts (Fig. 13c and d). Furthermore, it can also be observed that the fracture surface of the MM5h compacts exhibit a layered surface morphology (Fig. 13b). Such a layered fracture morphology indicates that the crack propagation during loading occurred through the lamellar colonies. It appears that the fine lamellar microstructure changes the mode of fracture by deflecting the crack propagation path at very short distances, leading to better fracture strength and ductility as compared to the equiaxed g microstructure with comparable grain size. As a result, the MM5h sintered compacts exhibit better mechanical properties as compared to the MM30h compacts. The relatively higher scattering of fracture strength values in the MM5h compacts as compared to that of MM30h compacts also appears to be related to the complex crack propagation during fracture due to its microstructural heterogeneities.
microstructure demonstrated a combination of higher fracture strength and ductility than the fine-grained g-based equiaxed structure. The better mechanical properties of “lamellar þ equiaxed” as compared to the equiaxed structure were attributed to the deflection of crack propagation path at short distances by the fine size lamellar structure.
4. Conclusions
References
Full density fine-grained Ti-rich TieAl alloy compacts were successfully prepared from dendritic coarse PREP pre-alloyed powders by mechanical milling followed by spark plasma sintering at moderate consolidation conditions, i.e. at 1323 K under 50 MPa. The enhanced densification of the milled powder was attributed to the superplastic deformation of milled powder particle during spark plasma sintering at lower temperatures due to the formation of recrystallized ultra-fine grains during sintering. It was also found that the microstructure of the sintered compacts depends on the extent of mechanical milling of PREPed powders. The short time milling, i.e. 5 h, resulted in fine-grained microstructure consisting of both a2/g lamellar and (a2 þ g) equiaxed grains, whereas extended milling times, i.e. 30 h, resulted in equiaxed g-based microstructure. Such a microstructural evolution was related with the uniformity of accumulated plastic deformation in the milled powders. The fine grained TieAl compacts exhibited high average fracture strength, over 900 MPa, irrespective of the nature of microstructure. The fine-grained “lamellar þ equiaxed”
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Acknowledgements This research was supported by the Japan Science and Technology Agency (JST) under Collaborative Research Based on Industrial Demand “Heterogeneous Structure Control: Towards Innovative Development of Metallic Structural Materials”, and by the Grant-inAid for Scientific Research on Innovative Area, “Bulk Nanostructured Metals”, through MEXT, Japan (contract No.22102004)). These supports are gratefully appreciated.
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[29] Keskinen J, Pogany A, Rubin J, Ruuskanen P. Carbide and hydride formation during mechanical alloying of titanium and aluminum with hexane. Mater Sci Eng A 1995;196:205e11. [30] Liu Y, Huang B, He Y, Zhou K. Processing TiAl-based alloy by elemental powder metallurgy. J Mater Sci Technol 2000;16:605e10. [31] Sun ZM, Hashimoto H. Fabrication of TiAl alloys by MA-PDS process and the mechanical properties. Intermetallics 2003;11:825e34. [32] Bhattacharya P, Bellon P, Averback RS, Hales SJ. Nanocrystalline TiAl powders synthesized by high energy ball milling: effects of milling parameters on yield and contamination. J Alloys Compd 2004;368:187e96. [33] Couret A, Molenat G, Galy J, Thomas M. Microstructure and mechanical properties of TiAl alloys consolidated by spark plasma sintering. Intermetallics 2008;16:1134e41. [34] Wang YH, Lin JP, He YH, Wang YL, Chen GL. Fabrication and SPS microstructures of Ti-45Al-8.5Nb-(W, B, Y) alloying powders. Intermetallics 2008;16: 215e24. [35] Lu X, He XB, Zhang B, Zhang L, Qu XH, Guo ZX. Microstructure and mechanical properties of a spark plasma sintered Ti-45Al-8.5Nb-0.2B-0.1Y alloy. Intermetallics 2009;17:840e6. [36] Jabbar H, Monchoux J, Houdellier F, Dolle M, Schimansky F, Pyczak F, et al. Microstructure and mechanical properties of high niobium containing TiAl alloys elaborated by spark plasma sintering. Intermetallics 2010;18:2312e21. [37] Guyon J, Hazotte A, Monchoux JP, Bouzy E. Effect of powder state on spark plasma sintering of TiAl alloys. Intermetallics 2013;34:94e100. [38] McCullough C, Valencia JJ, Levi CG, Mehrabian R. Microstructural analysis of rapidly solidified Ti-Al-X powders. Mater Sci Eng A 1990;124:83e101. [39] Tokizane M, Fukami T, Inaba T. Structure and mechanical properties of the hot pressed compact of Ti-rich TiAl powder produced by the plasma rotating electrode process. ISIJ Int 1991;31:1088e92. [40] Nishida M, Tateyama T, Tomoshige R, Morita K, Chiba A. Electron microscopy studies of Ti-47 at.% Al powder produced by plasma rotating electrode process. Scripta Met Mater 1992;27:335e40. [41] Wajata T, Isonishi K, Ameyama K, Tokizane M. Superplasticity in TiAl two phase material made from PREPed powder. ISIJ Int 1993;33: 884e8. [42] Nishida M, Morizono Y, Kai T, Sugimoto J, Chiba A, Kumagae R. Microstructure dependence of compactibility of rapidly solidified Ti-rich TiAl alloy powders produced by plasma rotating electrode process. Mater Trans JIM 1997;38: 334e43.
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
A novel powder metallurgy processing approach to prepare fine-grained Ti-rich TiAl-based alloys from pre-alloyed powders Sanjay K. Vajpai a, *, Kei Ameyama b,1 a b
Department of Mechanical Engineering, Ritsumeikan University, 1-1-1, Noji-higashi, Kusatsu, Shiga 525-8577, Japan Faculty of Science and Engineering, Ritsumeikan University, 1-1-1, Noji-higashi, Kusatsu, Shiga 525-8577, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 9 April 2013 Received in revised form 21 June 2013 Accepted 25 June 2013 Available online 13 July 2013
A relatively simple yet attractive and efficient powder metallurgy processing method, based on mechanical milling and spark plasma sintering of PREP (Plasma Rotating Electrode Processed) pre-alloyed powders, has been proposed to prepare full density fine-grained Ti-rich TieAl alloys with controlled microstructure. Full density fine-grained Ti-rich TieAl alloy compacts, with grain size in the range of 2e4 mm, were successfully prepared from dendritic coarse PREP pre-alloyed powders, which are otherwise difficult to consolidate, by mechanical milling followed by spark plasma sintering at moderate consolidation conditions. It has been demonstrated that the microstructure, i.e. lamellar or equiaxed, depends on the extent of mechanical milling. The fine-grained TieAl compacts exhibited extremely good mechanical properties. An attempt has been made to correlate the effect of mechanical milling on the microstructure and properties of the sintered TieAl compacts. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl B. Mechanical properties at ambient temperature C. Mechanical alloying and milling D. Powder metallurgy, including consolidation E. Sintering
1. Introduction In recent decades, Titanium-rich TiAl-based alloys have been considered a promising candidate as a high temperature structural material in automobile and aerospace sectors, due to their better high temperature specific strength than alloy steels and Nibased super-alloys, together with good oxidation, corrosion, and creep resistance at elevated temperatures [1e9]. Moreover, these alloys have been also found suitable as nuclear structural material due to their excellent radiation resistance and low neutron activation, together with good high-temperature creep resistance [10e13]. However, these alloys suffer from the inherent room temperature brittleness, which not only limits their commercial applications but also poses severe restrictions on the processing due to limited room temperature formability. The limited commercial application of these alloys can be attributed primarily to the coarse grained microstructure of the components fabricated
* Corresponding author. Tel.: þ81 77 561 2749; fax: þ81 77 561 2665. E-mail addresses: [email protected] (S.K. Vajpai), [email protected] (K. Ameyama). 1 Tel.: þ81 77 561 2756; fax: þ81 77 561 2665. 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.06.006
by the conventional ingot metallurgy processing, wherein the coarse-grained microstructure further aggravates the problem of brittleness and formability. It would be worth mentioning that the mechanical properties of the TiAl alloys can be improved by not only refining grain size but also controlling the amount and distribution of different phases in the microstructure, such as a2Ti3Al and g-TiAl phases in the form of either duplex or lamellar structure [14e20]. Although the desired phases can be obtained by controlling the composition and thermo-mechanical treatments [20e24], obtaining a fine-grained microstructure with a controlled distribution of various phases is extremely difficult. Therefore, the fabrication of near-net shape products with controlled microstructure, which governs the mechanical properties, via conventional ingot metallurgy processing remains a challenge due to limited room temperature formability and requirements of complex thermo-mechanical treatments [5e7]. The mechanical properties of TiAl based alloys depend on several factors, including purity of alloy, grain size, phase constitution, and distribution of various phases in the finished product [1e9,14e24]. Therefore, a microstructural design must include all the above mentioned factors to achieve optimum mechanical properties. Moreover, it is equally important to design a processing strategy which is not only capable of delivering near-net shape component with desired microstructure but also technologically
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viable and cost effective [24]. Powder metallurgy (PM) processing is an attractive way to achieve all the above mentioned goals. PM processing makes it possible to prepare near-net shape products, which minimizes the requirements of post processing machining [25]. Moreover, it also provides an opportunity to prepare finegrained components together with better control over chemical and phase compositions. There have been some efforts to prepare TiAl based alloys via powder metallurgy route [26e42]. The PM routes based on the mechanical milling of elemental powders suffer from the problems of contamination due to severe milling conditions, low yield due to sticking of elemental powders to the milling media, difficulty in handling the milled powders during subsequent processing due to their high reactivity, and limitations in achieving chemical homogeneity in the sintered components [26e32]. Moreover, controlling the microstructure is still an issue due to the formation of several undesirable phases, which have extremely adverse effect on the mechanical properties. The aforementioned problems associated with the PM processing can be effectively overcome by using atomized pre-alloyed powders as starting material. The rapid solidification during atomization ensures better control over chemical composition and chemical homogeneity, together with fine-grained microstructure, in the pre-alloyed powders. These characteristics provide an opportunity to control the microstructure in the finished components only by proper adjustments of the conditions during subsequent consolidation stage. In recent years, a few efforts have been made to prepare fine-grained TiAl-based alloys by spark plasma sintering of inert gas atomized pre-alloyed powders [33e37]. The sintered nearfull density compacts, thus prepared, exhibited improved mechanical properties as compared to their counterparts prepared by conventional ingot metallurgy method. However, it was found that the final density, microstructure, and properties of these compacts depend on the starting powder characteristics and sintering conditions. These parameters include average powder particle size, powder size distribution, initial composition, microstructure of the powder, sintering temperature, sintering pressure, and total sintering time. Therefore, the PM processing based on spark plasma sintering of pre-alloyed powders can become a viable method to prepare TiAl-based alloys with desired microstructure and properties, provided the correlation amongst the above parameters can be understood more clearly. Plasma rotating electrode process (PREP) is an attractive rapid solidification process which is capable of producing high purity pre-alloyed powders on an industrial scale. However, only a few efforts were made to prepare Ti-rich TiAl based alloys using PREPed powder as starting material, owing to the issues related with the consolidation of the PREPed powder with controlled microstructure [38e42]. It was found that the PREPed TiAl powder mass primarily consists of two different types of powder particles, i.e. (i) fine martensitic powder particles with size less than 100 mm, and (ii) coarse dendritic powder particles with size over 100 mm. It was also demonstrated that the fine martensitic particles are easy to consolidate to near full density, whereas coarse dendritic powders are very difficult to consolidate. Generally, a combination of high pressure, high temperature, and long pressing time is required to achieve near full density. However, the requirements of these consolidation conditions restrict the control over the final microstructure in the sintered compacts. Furthermore, the PREPed powder mass contains coarse dendritic powders in the range of 60e80 mass percent, which severely restricts the commercial viability of using PREPed powders as starting material. Therefore, it is desired to develop a PM processing methodology based on PREP powders wherein near-net shape full density components can be prepared under moderate consolidation conditions together with better control over microstructure.
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In the present work, a new powder metallurgy processing approach has been proposed to prepare full density compacts, with controlled microstructure, under moderate consolidation conditions using “difficult-to-consolidate” coarse dendritic PREPed TiAl pre-alloyed powders as starting material. The proposed PM route primarily consists of mechanical milling of coarse dendritic powders followed by spark plasma sintering. It is envisaged that the mechanical milling of the prealloyed powder would enhance the sintering kinetics, resulting in near full density compacts at relatively lower sintering temperature and pressure in short period of time. In order to understand the effect of mechanical milling on the density and microstructure of the sintered compacts, PREP powders were milled for different time period. The microstructural characteristics of the starting PREPed powder and mechanically milled powders are presented to understand the effect of milling on the microstructure and morphology of prealloyed powders. Subsequently, the microstructural characteristics of the sintered compacts prepared from mechanically milled powders are also presented to elaborate the effect of milling time on the densification and microstructural evolution in the sintered compacts. Finally, in order to assess the mechanical properties of the sintered compacts, the results of the hardness measurements and flexural tests are presented and discussed. 2. Experimental procedure Ti-rich TieAl pre-alloyed powders, with particle size in the range of 200e1400 mm, produced by helium-PREP were used in the present study. The chemical composition of the powders is shown in the Table 1. The pre-alloyed powder was mechanically milled (MM) in a planetary ball mill up to 108 ks (30 h). To minimize the contamination due to mechanical milling, the milling was carried out under inert gas (argon) atmosphere and low ball-to-powder ratio (2:1). Subsequently, the powders milled for 18 ks (5 h) and 108 ks (30 h) were sintered by spark plasma sintering at 1323 K for 3.6 ks (1 h) under the pressure of 50 MPa. The powder was filled in the graphite die and the pressure was applied during sintering. Henceforth, the sintered compacts prepared from initial PREP, MMed for 5 h, and MMed for 30 h powders will be referred as MM0h, MM5h, and MM30h, respectively. The average particle size and size distribution of the initial and milled powders were evaluated by laser particle size analyzer. X-ray diffraction and scanning electron microscopy (SEM) techniques were used for phase identification and microstructural characteristics of the starting powder, milled powders, and sintered compacts. In order to achieve the phase contrast in the micrographs, images were acquired using backscattered electron detector in the SEM. Four point bend test and micro-hardness tests were carried out to evaluate the mechanical behaviour of the sintered compacts. The micro-hardness tests were carried out under the load of 980.7 mN, and average hardness was calculated using 25 measurements. The four point bend test was carried out on rectangular specimens with approximate dimensions 1.5 2.5 15 (all dimensions in mm). The rectangular bend test specimens were polished using 0.1 mm alumina suspension to minimize the roughness on the surface. Finally, fractured surfaces of the specimens were observed under the SEM to understand the nature of fracture of the specimens.
Table 1 Chemical composition of the starting PREPed powder. Element
Al
O
Fe
Si
Ti
Amount (mass%)
33.67
0.387
0.02
0.07
Bal.
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3. Results and discussion 3.1. Microstructural characteristics of initial PREP powder Fig. 1 shows the morphology and particle size distribution of the PREPed TiAl powders. It can be seen that the powder particles are spherical in shape, and have wide and bimodal size distribution. The particle size varied in the range of 180e1400 mm, and the average particle size was approximately 500 mm. It can be noted that most of the particles are large in size. The bimodal particle size distribution is a common feature of PREPed powders. The typical surface morphology of the PREPed powders is shown in Fig. 2. It was noted that most of the particles consisted of dendritic structure. The PREP TiAl powders have been found to consist of either martensitic or dendritic microstructure depending on the cooling rate during processing [39e42]. The cooling rate also depends on the particle size, and it has been observed frequently that the small sized particles (<100 mm size) exhibit martensitic structure, whereas larger particles exhibit dendritic microstructure. Since the present PREP TiAl powders are very large in size, it is obvious that the cooling rate during flight would be “relatively” slower as compared to small sized martensitic particles, leading to the formation of dendritic microstructure.
Fig. 2. SEM micrographs of PREPed powders showing the dendritic surface morphology.
Fig. 3 shows the XRD pattern of the PREPed powders. It can be seen that hexagonal a2 (Ti3Al) is the predominant phase present in the particles together with a small amount of tetragonal g (TiAl) phase. The presence of a2 as predominant phase suggests that the hexagonal a phase is the first phase which was formed during solidification. The presence of ordered a2 hexagonal phase together with a precipitation of g phase in small amounts also suggests that
Fig. 1. (a) Morphology and (b) particle size distribution of PREPed TieAl powder.
Fig. 3. XRD pattern of the initial PREPed powder.
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the solidification rate remains sufficiently high to retain the hexagonal a phase in the solidified powder particles. However, the ordering of a phase is not suppressed even at such a rapid cooling rate, leading to the formation of the ordered variant of a, i.e. hexagonal a2. Moreover, the presence of ordered a2 phase in large amounts also suggests that the rapid solidification during preparation of PREP TiAl powder also leads to the retention of supersaturated a2 phase. The presence of dendritic microstructure, accompanied with the formation of hexagonal a2 phase, has been found to be a typical feature of the coarse particles of PREP TiAl powders [39e42]. 3.2. Microstructural characteristics of mechanically milled powders The PREPed TiAl powders were mechanically milled in a planetary ball mill for different period of time. Fig. 4 shows the average particle size and size distribution of the MMed powders together with initial powder. It can be clearly observed that the average particle size reduced very rapidly within initial 5 h of milling followed by a sluggish refinement in particle size (Fig. 4a). The particle size after initial 5 h of milling was found to be approximately 50 mm. Subsequently, the average particle size of approximately 10 mm was obtained after 15 h of milling. Furthermore, it can be noticed that further milling under the present conditions did not result in any further refinement in particle size even after prolong milling time, i.e. 30 h, and the average particle size remained approximately 10 mm. It would be worth mentioning that the mechanical milling resulted in not only a small average particle size
Fig. 4. Effect of mechanical milling on the (a) average particle size, and (b) size distribution of PREPed powders.
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but also reduced size distribution (Fig. 4b). Fig. 5 shows the effect of milling time on the morphology of the PREPed powder. It can be seen that after initial 5 h of milling, initial PREP powder particles fragmented in to smaller size irregular shaped particles, primarily via brittle mode of fracture (Fig. 5a and b). Although the initial milling period appears to be dominated by the brittle fracture of initial PREP powders, it would be worth mentioning that the fractured powder particles would also undergo to some degree of plastic deformation during this stage. Fig. 5c and d shows the morphology of the PREP powders milled for 15 h. A comparison of Fig. 5a and c clearly shows that the increased milling time led to the further fragmentation of the powder particles, resulting in the refinement of the particle size. Moreover, it can also be observed that the resulting powder particles are slightly flat and irregularshaped, suggesting that the particles appear to have undergone a significant plastic deformation. Fig. 5e and f shows the morphology of the PREP powders milled for 15 h. A comparison of Fig. 5c and e indicates that the subsequent milling after 15 h did not induced any significant fragmentation of the powder particles, as demonstrated by particle size analysis earlier and shown in Fig. 4. However, the flaky morphology of the milled powder particles, as shown in Fig. 5f, is a clear evidence of severe plastic deformation of the powders due to milling. Fig. 6 shows the XRD pattern of initial and MMed powders. It can be seen that there is neither any shifting in the positions of peaks nor any change in the relative intensity of g-peaks with respect to a2 peaks. These results indicate that the mechanical milling did not induce any phase change in the PREP TiAl powder. However, it can be noted that there was a significant broadening in the XRD peaks as a result of milling. Fig. 7 shows the variation in the width of the three most intense peaks with increasing milling time, wherein full width at half maxima (FWHM) of the peak has been considered as the peak width. It can be observed that the peak width increased rapidly up to approximately 5 h of milling. The subsequent milling led to a moderate increment in peak broadening between 5 and 15 h followed by a sluggish increment up to 30 h of milling. The important factors contributing to the XRD peak broadening during mechanical milling are (i) introduction of lattice strains caused by the accumulation of defects, such as dislocations and point defects, and (ii) crystallite size refinement. It is well established that large amounts of lattice defects are accumulated during the early stages of mechanical milling followed by achieving a saturation level, wherein the rate of generation and annihilation of these defects attain equilibrium state. The attainment of the equilibrium state is characterized by the formation of very fine size crystallites in the milled powders. Thus, the XRD peak broadening during early stages of milling is primarily related with the accumulation of large amounts of lattice defects, whereas the peak broadening after prolonged milling can be attributed to the formation of extremely small crystallites. Therefore, in the present case, it appears that the early stages of milling, up to 15 h, involves fracturing of powder particles together with an accumulation of large amounts of lattice defects due to plastic deformation of fragmented particles. However, the subsequent milling leads to the severe plastic deformation of fragmented particles, as evidenced by the flaky morphology of the 30 h milled powder, resulting in the powder particles with extremely fine-sized crystallites. It is envisaged that the mechanically milled irregular-shaped small-sized powder particles would exhibit better packing, in terms of interparticle contact area, as compared to the large-sized spherical PREP powders. Moreover, the accumulated defects and fine-grained structure in the milled powder particles would enhance the diffusion kinetics during subsequent sintering, resulting in the improved sinterability and chemical homogeneity. As a result, it would be possible to prepare near full density
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Fig. 5. SEM micrographs of PREPed powders milled for (a, b) 5h, (c, d) 15 h, and (e, f) 30 h.
Fig. 6. XRD patterns of the MMed PREP TiAl powders.
Fig. 7. Variation in the XRD peak width with increasing milling time.
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Fig. 8. Typical micrographs showing the thickness cross-section of the as-polished sintered compacts prepared from (a) initial PREP powder, and (b) MMed PREP powder.
products with improved chemical homogeneity in relatively short processing time under the low pressure and low temperature processing conditions. In the present work, powders milled for 5 h (short-time milled) and 30 h (prolonged milled) were only processed further to understand the effect of two extreme milling conditions on the microstructure and properties of the resulting TiAl compacts.
followed by sintering under low pressure, low temperature, and relatively shorter sintering time. Fig. 9 shows the XRD results of the sintered compacts prepared from initial as well as mechanically milled PREP powders. It is interesting to note that the sintered compacts consist of primarily tetragonal g phase (TiAl) together with a small amount of hexagonal a2 (Ti3Al), which is altogether different from the as PREPed powder wherein hexagonal a2 (Ti3Al) was the primary phase present. These results suggest that the initial PREP powder consisted of supersaturated hexagonal a2 (Ti3Al) phase, as a result of rapid cooling. Moreover, the presence of tetragonal g phase (TiAl) as the major phase in the sintered compacts is not unexpected considering the composition of the initial powder. It appears that the sintering has led to the formation of large amounts of hexagonal g phase (TiAl) via precipitation from supersaturated hexagonal a2 (Ti3Al) phase, which is present in the rapidly solidified initial PREP TiAl powder. Fig. 10a and b shows the SEM micrographs depicting the microstructures of sintered compacts prepared from the initial PREP TiAl powder. It can be seen that the most of the powder particles in the sintered compacts consisted of diffused dendritic microstructure (Fig. 10a). Although some recrystallized equiaxed grains, g as well as a2/g lamellar microstructure were also found to be present in the compacts (Fig. 10b), only a few such areas were found to exist. Moreover, it can also be noticed that a2 phase is present either on the grain boundaries or as a layer in the lamellar structure. These observations are in conformity with the XRD results, showing that the initial PREP powders consisted of supersaturated hexagonal a2 (Ti3Al) phase and its dissociation during sintering leads to the formation of ordered Ti3Al and TiAl phases. Fig. 10c shows a typical SEM micrograph depicting the microstructure of sintered compacts prepared from 5 h milled PREP TieAl alloy powders. It can be clearly observed that the sintered compacts consist of primarily lamellar microstructure, more or less equiaxed lamella colonies consisting of alternate layers of a2/g phases, together with equiaxed fine size g grains. The presence of equiaxed grains/lamella colonies indicates that the microstructural evolution during sintering was dominated by the recrystallization and grain growth. It can also be noted that the a2 phase is distributed in the matrix as either submicron-sized particles on the boundaries of equiaxed g grains or extremely fine sized layers, having less than 0.2 mm thickness, in the lamellar structure. Therefore, it is evident that a short time milling led to not only the disappearance of initial coarse dendritic microstructure but also the formation of fine grained microstructure. It has been already shown that the short time milling, i.e. 5 h in the present case, induces large amounts of
3.3. Microstructural characteristics of sintered compacts The initial PREP and MMed powders were sintered by spark plasma sintering at 1323 K for 60 min under 50 MPa pressure. Fig. 8 shows the micrographs depicting the as-polished thickness crosssection of the sintered compacts. Fig. 8a shows that the sintered compacts of initial PREP powders consist of significant amount of open porosity, i.e. approximately 25 volume%. It can be noticed that the spherical PREP powder particles only joined together by forming a neck as a result of sintering and did not undergo any significant plastic deformation during sintering, resulting in large amounts of retained porosity. Fig. 8b shows the thickness crosssection of the sintered compacts prepared from 5 h milled powder. The absence of any appreciable amount of residual porosity shows that a short-time milled irregular shaped powder particles resulted in the full density sintered compacts. Similarly, 30 h milled powder also resulted in full density sintered compacts. Therefore, it is evident that full density TiAl compacts can be successfully prepared from large sized PREP TiAl powders by mechanical milling
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Fig. 9. XRD patterns of the sintered TiAl compacts.
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Fig. 10. Typical SEM micrographs of sintered TiAl compacts prepared from (a, b) initial PREP powder, (c) 5 h MMed powder, and (d) 30 h MMed powder.
microstructural defects in the powders. Furthermore, it is well known that these microstructural defects enhance the diffusion kinetics, and the diffusivity of atoms of various elements in an alloy becomes very high even at very low temperatures. Therefore, it appears that the presence of these microstructural defects, such as dislocations and point defects, in the short time milled PREP powders enhanced the diffusion process significantly during sintering, leading to the disappearance of coarse dendritic microstructure and formation of refined microstructure with a particular phase distribution. Since short time milling does not ensures uniform accumulated plastic deformation in the milled powders, the microstructural evolution during sintering appears to be dependent on the local accumulated plastic deformation in the short time milled powders. As a result, both a2/g lamellar and (a2 þ g) equiaxed microstructures are found in the sintered compacts prepared from 5 h MMed powders. The average grain size of the sintered compacts was found to be approximately 3.82 1.11 mm. Fig. 10d shows a typical SEM micrograph exhibiting the microstructure of the sintered compacts prepared from 30 h milled PREP TieAl alloy powders. It can be observed that sintered compacts primarily consist of very fine size equiaxed grains of g phase (tetragonal ordered TiAl), having average grain size 2.34 1 mm. It can also be noticed that there is no evidence of existence of lamellar microstructure in significant amounts and the a2 phase is present primarily on the boundaries of the equiaxed g grains in the form of very small particles with size less than 2 mm. It would be worth mention that a significant amount of a2 phase particles have size in the submicron range. These microstructural features indicate that the severely deformed supersaturated PREP powders undergo recrystallization during sintering to form strain free small grains, together with the simultaneous precipitation of small sized a2 on the grain boundaries of these newly formed grains. It appears that these extremely small sized a2 phase present on the grain boundaries restricted further growth of the recrystallized g grains during sintering, resulting in the fine-grained microstructure consisting of equiaxed g grains. Recently, Guyon et al. [37] have also reported the formation of submicron/nano size recrystallized grains during
spark plasma sintering of milled TieAl-based prealloyed powders prepared by inert gas atomization. Similar to the present work, fine size a2 was found to be present on the grain boundaries of the finesize g grains wherein it was postulated that these fine size a2 particles/grains restrict any significant growth of the recrystallized g grains during subsequent heating. Table 2 shows the relative amounts (volume%) of a2 and g phases in the sintered compacts prepared from initial and MMed powders. The average amounts were estimated from the SEM micrographs which were acquired using backscattered electron detector. It can be noted that all the sintered compacts consist of comparable relative amounts of a2 and g phases irrespective of the condition of the starting powder, i.e. as PREPed or MMed for different periods of time. Therefore, these results suggest that the extent of mechanical milling only brings about the redistribution of the a2 and g phases in the sintered compacts prepared from PREP powders, and their relative amounts is primarily governed by the chemical composition of the initial PREP powder. The above results show that the coarse PREP pre-alloyed TieAl powders are difficult to consolidate to near-full density under the moderate sintering conditions such as short sintering time, relatively lower sintering temperature, and low pressure. Moreover, the initial dendritic microstructure also remains retained under these sintering conditions. On the other hand, it has been demonstrated that mechanical milling of these “difficult-to-consolidate” coarse PREP powders leads to not only near-full density compacts with fine-grained microstructure but also the disappearance of dendritic microstructure. Ameyama et al. [27] demonstrated that fine grained Table 2 Average amounts of a2 and g phases present in the sintered compacts. Material
Average amount of a2 phase (vol.%)
Average amount of g phase (vol.%)
PREP þ SPS (1323 K) MMed (5 h) þ SPS (1323 K) MMed (30 h) þ SPS (1323 K)
26 23 20
74 77 80
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Fig. 11 shows the average micro-hardness values of the sintered specimens prepared from initial PREP, 5 h milled, and 30 h milled powders. It can be noticed that the average micro-hardness of the compacts prepared from short-time milled powders, i.e. 5 h milled, is only slightly higher than that of the compacts prepared from initial PREP powders. However, it is interesting to note that the average micro-hardness of the compacts prepared from 30 h milled powders is significantly higher than that of the compacts prepared from initial PREP and 5 h milled powders. Moreover, it can also be noticed that the dispersion of micro-hardness values from the mean value decreased with increasing milling time. It has been demonstrated earlier that the sintered compacts of initial PREP powder consist of diffused dendritic and lamellar microstructure, together with the precipitation of a2 particles of 2e5 mm size. On the other hand, short-time milling results in the formation of
lamellar microstructure due to the enhanced short-range diffusivity. It would be worth pointing out that both the microstructures are not significantly different in terms of distribution of a2 and g phases, resulting in comparable hardness values. It would also be worth mentioning that the distribution of a2 and g phases in the lamellar microstructure is relatively much finer as compared to that of diffused dendritic microstructure. This fact is reflected by the smaller dispersion of hardness values from the mean hardness value in the sintered compacts prepared from 5 h milled powder as compared to that of initial PREP powders. The significantly high average hardness of sintered compacts from 30 h milled powder can be attributed to the presence of fine-size equiaxed grains of g phase, which inherently exhibits higher brittleness as compared to the lamellar microstructure. Fig. 12 shows representative flexural strength vs. deflection curves of the specimens prepared from the MM5h and MM30h sintered compacts, having average fracture strength 978.41 85.31 MPa and 883.23 3.64 MPa, respectively. It can be noted that the sintered compacts of 5 h MMed powder exhibited slightly higher average fracture strength and deflection-to-fracture, together with lower relative hardness, than those of compacts of 30 h MMed powder. Since a combination of higher deflection-to-fracture and lower hardness values can be regarded as an indicator of better ductility in case of brittle materials, the above results show that the MM5h sintered compacts exhibit better combination of strength and ductility as compared to the MM30h compacts. It would be worth mention that the average fracture strength of the present TiAl compacts is considerably higher than coarse grained TiAl-based alloys and comparable to that of the fine-grained compacts prepared from inert gas atomized prealloyed powders [1,2,33,35,36]. The high fracture strength of the present compacts can be attributed to the refined microstructure. It can also be noted that the MM5h compacts have not only the higher average fracture strength but also larger scattering in the strength values as compared to MM30h compacts. However, the scattering in the fracture strength values is considerably small as compared to those reported in the literature for TiAl-based alloys. Fig. 13 shows the morphology of the fractured surfaces of the flexural test specimens prepared from the MM5h and MM30h sintered compacts. It can be clearly observed that all the sintered compacts exhibit brittle nature of fracture. However, it can be noted that the mode of fracture is different for both the compacts, i.e. the deformation mode of MM5h compacts is dominated by trans-
Fig. 11. Micro-hardness of the sintered compacts prepared from initial PREP and MMed powders.
Fig. 12. Representative Flexural strengthedeflection curves of sintered compacts prepared from mechanically milled powders.
TiAl-based alloys undergo superplastic deformation even at lower temperatures. In the present work, it has been demonstrated that the spark plasma sintering of milled powder resulted in the formation of fine-grained recrystallized microstructure. Therefore, the enhanced densification of the milled PREP pre-alloyed powders can be attributed to the superplastic deformation of the milled power particles even at lower sintering temperature. Similar observations regarding the enhanced densification of milled inert-gas-atomized pre-alloyed TiAl-based powder at lower temperature have also been made by Guyon et al. [37]. Furthermore, it is also demonstrated that the final microstructure of the sintered compacts also depends on the severity of plastic deformation induced in the powders by mechanical milling. Therefore, the present processing route provides an opportunity to prepare full density compacts with desired microstructure by adjusting the milling and sintering conditions. It would be worth mentioning that near-full density compacts can be prepared from coarse PREP powders by applying high temperature and high pressure for relatively longer sintering time, which is typically 4e5 h [39,41,42]. However, such processing conditions not only limit the bulk production of near-net shaped full density products but also result in the coarse-grained microstructure. Moreover, it is also extremely difficult to achieve the desired microstructure with controlled distribution of various phases in the microstructure when coarse PREP powders are used as starting material. 3.4. Mechanical properties of the sintered compacts
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Fig. 13. Fractured surfaces of the flexural test specimens prepared from the sintered compacts of (a, b) 5 h, and (c, d) 30 h MMed PREP powders.
granular mode of fracture (Fig. 13a and b) whereas it is primarily intergranular in case of MM30h compacts (Fig. 13c and d). Furthermore, it can also be observed that the fracture surface of the MM5h compacts exhibit a layered surface morphology (Fig. 13b). Such a layered fracture morphology indicates that the crack propagation during loading occurred through the lamellar colonies. It appears that the fine lamellar microstructure changes the mode of fracture by deflecting the crack propagation path at very short distances, leading to better fracture strength and ductility as compared to the equiaxed g microstructure with comparable grain size. As a result, the MM5h sintered compacts exhibit better mechanical properties as compared to the MM30h compacts. The relatively higher scattering of fracture strength values in the MM5h compacts as compared to that of MM30h compacts also appears to be related to the complex crack propagation during fracture due to its microstructural heterogeneities.
microstructure demonstrated a combination of higher fracture strength and ductility than the fine-grained g-based equiaxed structure. The better mechanical properties of “lamellar þ equiaxed” as compared to the equiaxed structure were attributed to the deflection of crack propagation path at short distances by the fine size lamellar structure.
4. Conclusions
References
Full density fine-grained Ti-rich TieAl alloy compacts were successfully prepared from dendritic coarse PREP pre-alloyed powders by mechanical milling followed by spark plasma sintering at moderate consolidation conditions, i.e. at 1323 K under 50 MPa. The enhanced densification of the milled powder was attributed to the superplastic deformation of milled powder particle during spark plasma sintering at lower temperatures due to the formation of recrystallized ultra-fine grains during sintering. It was also found that the microstructure of the sintered compacts depends on the extent of mechanical milling of PREPed powders. The short time milling, i.e. 5 h, resulted in fine-grained microstructure consisting of both a2/g lamellar and (a2 þ g) equiaxed grains, whereas extended milling times, i.e. 30 h, resulted in equiaxed g-based microstructure. Such a microstructural evolution was related with the uniformity of accumulated plastic deformation in the milled powders. The fine grained TieAl compacts exhibited high average fracture strength, over 900 MPa, irrespective of the nature of microstructure. The fine-grained “lamellar þ equiaxed”
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Acknowledgements This research was supported by the Japan Science and Technology Agency (JST) under Collaborative Research Based on Industrial Demand “Heterogeneous Structure Control: Towards Innovative Development of Metallic Structural Materials”, and by the Grant-inAid for Scientific Research on Innovative Area, “Bulk Nanostructured Metals”, through MEXT, Japan (contract No.22102004)). These supports are gratefully appreciated.
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