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Size dependent phase transformation in atomized TiAl powders
Intermetallics 61 (2015) 72e79
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Size dependent phase transformation in atomized TiAl powders Dong-Ye Yang a, b, c, Shu Guo a, b, Hua-Xin Peng c, d, Fu-Yang Cao a, b, Na Liu e, Jian-Fei Sun a, b, * a
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, PR China c Advanced Composites Centre for Innovation and Science (ACCIS), Department of Aerospace Engineering, University of Bristol, Bristol BS8 1TR, UK d School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China e Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 20 November 2014 Received in revised form 11 February 2015 Accepted 25 February 2015 Available online
The paper focuses on the phase transformation of undercooled Tie48Al (at.%) droplets atomized by high pressure gas. The microstructural evolution was analyzed using the transient nucleation theory in combination with the microstructural examinations. The primary phase, final phase volume fraction and the hardness are related to the droplets size to provide fundamental understanding. The competitive formation of the primary phases a and b are strongly controlled by the droplets size, and a critical diameter of about 25 mm was identified corresponding to an undercooling of 102 K. The final phase volume fraction of g increased with the increase of powder size from 12% to 85%. The highest hardness with the value of 652 HV was obtained for powder of 50 mm in diameter. © 2015 Elsevier Ltd. All rights reserved.
Keywords: A. Intermetallics B. Phase transformation C. Rapid solidification D. Microstructure E. Phase stability, prediction
1. Introduction TiAl has been extensively investigated as a promising candidate material for high temperature applications in aviation, automotive and power generation industries, due to its low density, relatively high specific strength and excellent corrosion resistance, as well as good oxidation and creep resistance at high temperature [1e3]. However, the poor deformability at room temperature has severely limited its further development [4e6]. To overcome this deficiency, powder metallurgy is employed to provide fine homogenous microstructures and in some cases enhanced properties to improve the performance of TiAl based alloy [7,8]. In this context, gas atomization is a prominent technique applied with many benefits for pre-alloyed powders, due to its efficiency and high performance [9e11]. During gas atomization, a liquid stream is broken up into fine spherical droplets of different sizes by the impact of a highvelocity gas flow [12,13]. The TiAl alloy droplets undergo rapid * Corresponding author. Harbin Institute of Technology, School of Materials Science and Engineering, Harbin, Heilongjiang 150001, PR China. Tel./fax: þ86 451 86413904. E-mail address: [email protected] (J.-F. Sun). http://dx.doi.org/10.1016/j.intermet.2015.02.017 0966-9795/© 2015 Elsevier Ltd. All rights reserved.
solidification processing (RSP) with the change of phase selection process toward more ductile metastable phases. The phase evolution in a non-equilibrium solidification process depends critically on the interplay between undercooling and cooling rate as a function of droplets size. Further influences under investigation include primary phase competition and suppression of solid state transformations. The final phases state, volume fraction and hardness are the important factors affecting the sintering performance of pre-alloy powders. During rapid solidification processing, the competitive nucleation of high-temperature metastable phases is seriously separated from the equilibrium solidification state. The thermodynamic and kinetic conditions were considered to fit for the guidance of phase separation under non-equilibrium solidification. Increasing the undercooling, the primary phase changes from b (A2-type bcc) to a (A3-type hcp) and g (L10-type fcc) phases [14]. When the primary phase is b-Ti, Zhou [15] has observed morphological transition from b to a in a hypo-peritectic Ti47Al53 alloy sample undercooled by 20 K. The primary or secondary a is always surrounded by g segregate. The g acts the primary phase in the achievable undercooling range of Ti44Al56 alloy [16]. The phase transformation of a / g is determined by diffusion rate which is
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quite sensitive to the cooling rate [17]. During the atomization, the transformation a / g was suppressed due to the high cooling rate of droplets about 104e106 K s1, resulting in a decrease of a2 þ g lamella structure replaced by a / a2 ordering transformation. The objective of this paper is to analyze the solidification pathways of undercooled Tie48Al droplets based on the thermodynamic and kinetic theories. Relationships are established between the primary phase, volume fraction, hardness and the gas atomized Tie48Al droplets size to provide fundamental understanding of the subsequent powder metallurgy (PM) process. 2. Experimental procedure The powders of TiAl alloy (nominal composition: Tie48Al, at.%) were prepared by high pressure gas atomization in an argon spray
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chamber environment, and the TiAl alloy was melted by levitation melting technique in order to achieve composition homogeneity and purity (140 wt.ppm N, 600 wt.ppm O). These powders were firstly classified into 5 different size groups, ranging from 5 to 150 mm in diameter. The classification was carried out by screening and a sedimentation technique [18]. The microstructure of the powders was characterized using scanning electron microscopy (SEM), X-ray diffraction (XRD) and transmission electron microscopy (TEM). For SEM observations, the powders were mounted in epoxy, polished, etched with Kroll solution 4 vol. % HNO3 þ 2 vol. % HF þ 94 vol. % H2O and coated with conductive layer. Transmission electron microscopy (TEM) was used to determine the phases and the microstructure of the powders in addition to SEM and optical microscopy. The TEM specimens were prepared by
Fig. 1. Microstructures of atomized Tie48Al powder (less than 15 mm): (a) Featureless surface (SEM), (b) TEM image of the powder, (c) Corresponding SAED pattern of the matrix a2 phase, zone axis ½0001a2 , (d) Lamella structure, corresponding SAED pattern of the g phase, zone axis ½110g , (e) Corresponding SAED pattern of a phase, zone axis ½1120a , (f) Mixed a and a2 electron diffraction, from the ½0110a =½0001a2 zone axis.
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the FIB-TEM foil preparation technique and a conventional plating technique [14]. For plating technique, the powders were firstly placed loosely on a copper sheet substrate. This copper sheet was then electroplated with nickel, using a low current density (4e5 mA cm2) to ensure good adherence of the nickel plating to the powders. The plating was about 650 mm thick. The copper sheet was removed, and the foils were then subjected to polish, dimple and ion mill. 3. Experimental results The microstructures and phases developed in the Tie48Al at.% powders change with the rapid solidification regime such as the undercooling and cooling rates. The regime is strongly dependent on the size of powders. Thus the results are presented for powder size ranges representing as fine (less than 15 mm), medium (40e70 mm) and large (100e150 mm). In the fine powders, a featureless surface morphology is observed by SEM ranging in size less than 15 mm as shown in Fig. 1(a). Examination of the powders cross-sections generally shows the presence of cellular and lamella structure morphologies by TEM in Fig. 1(b). The volume fraction of cellular microstructure in powders is about 80% with an average cellular spacing of 1e3 mm. Electron diffraction pattern shows the matrix phase of the cells to be a2, from the ½0001a2 zone axis in Fig. 1(c). Fig. 1(b) shows some region containing only a2 (marked as region A) while region B contains the usual a2 þ g lamella structure with a spacing of 0.1e0.25 mm, with higher magnification micrograph as shown in Fig. 1(d). However, the grey phase of g lamella is fine and limited precipitation (less than 0.03 mm). The segregate
structures obtained in the a2 grain boundaries consist of the lamella and feather-like phase, and the electron diffraction pattern from the ½110g zone axis shows the phase to be g in Fig. 1(d). Fig. 1(e) presents corresponding electron diffraction pattern composed of plates of a significantly different hexagonal phase, and this is disordered a phase formed from the liquid, from the ½1120a zone axis reported previously by McCullough [14]. Fig. 1(f) shows a mixed a and a2 electron diffraction pattern from the ½0110a =½0001a2 zone axis, and a direct evidence of suppression in the ordering a / a2. Thus the primary phase formed from the liquid is a phase. In addition, energy dispersive spectroscopy (EDS) analysis an average aluminum concentrations for the cells and segregate g phase areas of 47.9 at.% and 50.2 at.%, respectively, indicating segregation was suppressed owing to the result of rapid solidification. For powders of size from 40 to 70 mm, the surface morphology exhibits a dendrite appearance with a little short dendrite arms as shown in Fig. 2(a). The grains are rather small (about 15 mm) and the growth is almost equiaxed to show a clear hexagonal or polygon pattern. Examination of the powders cross-sections generally shows the presence of concentric liquid/solid interface solidification geometry due to the multi-nucleation events taking place in Fig. 2(b). The volume fraction of a2 þ g lamella structure increases to about 30% with an average spacing of 0.3 mm and the grey phase of g lamella is about 0.1 mm. From the TEM observations this can largely be attributed to the suppressed decomposition of a2 into the a2 þ g. The net-like structure shows the presence of fine g phase grains and a2 þ g lamella (less than 0.02 mm) with twins structures
Fig. 2. Microstructures of atomized Tie48Al powder (40e70 mm): SEM views of (a) Dendrite surface and (b) Equiaxed dendritic morphology, TEM views of (c) a2 þ g lamella structure with twins structure from net-like structure with corresponding SAED pattern showing ½1210a2 ==½110g , (d) Fine g phase grain.
D.-Y. Yang et al. / Intermetallics 61 (2015) 72e79
morphologies as indicated by TEM images in Figs. 2(c) and (d), respectively. Fig. 2(c) shows the selected area diffraction patterns of the net-like structure phase, which proved to be the phase a2 þ g with ½1210a2 ==½110g for the transformed region. In the large powders, the equiaxed dendrite surface morphology with development of long dendrite arms ranging in size from 100 to 150 mm is observed by SEM as shown in Fig. 3(a). The equiaxed dendrites are similar to those seen in conventional solidification in which the primary phase was b-Ti. Fig. 3(b) shows cross-sectional view of a large powder, where the dendrite arms are almost orthogonal to one another. Thus, it seems more likely to be the characteristic growth of the cubic b phase from the liquid. Electron diffraction (Fig. 3(d)) of the selected area in Fig.3(c) reveals the presence of B2 phase of the transformation of parent b, from the ½001B2 zone axis. Hence, in the large powders there is evidences for the primary phase forming from the liquid is b phase.
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The volume fraction of a2 þ g lamella and net-like structure increases to about 70% with sole a2 surrounded by g phase in Fig. 3(c). The spacing of a2 þ g lamella is about 0.3 mm and the grey phase of g lamella is about 0.15 mm as shown in Fig. 3(e). EDS analysis reveals the average composition for the dendrite and segregate regions to be 47.7 at.% Al and 51.1% Al, respectively. The average aluminum concentration is slightly higher in comparison to fine powders. Fig. 4 shows the X-ray diffraction analyses of powders of different sizes at room temperature (298 K). The spectrum contains a mixture of a2 and g peaks, but is dominated by the large peaks of ð201Þa2 and (111)g near 41 and 39 , respectively. With increasing powder size, the intensity of the a2 peaks reduced greatly with the loss of some previous peaks. Concurrently new g signals corresponding to (001), (110) and (222) appeared at 22 , 32 and 85 , respectively. In addition, the peaks of (202)g and (220)g near 65 are
Fig. 3. Microstructures of atomized Tie48Al powder (100e150 mm): SEM views of (a) Dendrite surface and (b) Orthogonal symmetry dendrites, TEM views of (c) sole a2 surrounded by g phase and B2 phase with corresponding SAED pattern (d), zone axis ½001B2 , (e) a2þg lamella structure.
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are very close, and it is limited to predict the primary phase under non-equilibrium conditions. Thus, further kinetic consideration of incubation time has to be performed. Previous studies have concluded that the nucleation mechanism of gas atomized droplets undercool below the alloy liquidus temperature is the heterogeneous nucleation [22,23]. Heterogeneous nucleation generally dictates the practical limit of metastability which is dependent on the transient time. Temperature vs transient time curves can be calculated and used to predict the selection of competing phases during rapid solidification. The nucleation behavior of competing phases in the undercooled TiAl droplets could be evaluated by the transient nucleation theory proposed by Shao and Tsakiropoulos [24]. The incubation time t can be expressed as:
t¼
Fig. 4. XRD spectra for atomized Tie48Al alloy powders of different sizes.
overlapped due to the decrease of lattice constant by rapid solidification. 4. Discussion 4.1. Thermodynamic analysis of phase formation In order to analyze the primary phase of the atomized TiAl droplets, the disordered solution phase, liquid (L), a (Ti) and b (Ti) and the intermetallic compound g (TiAl) phase are considered in the calculation of the Gibbs energy of the Tie48Al, at.% alloy. The Gibbs energy of the present phases has to be expressed as analytical functions of composition, temperature and pressure. The liquid (L), a (Ti) and b (Ti) phases could be accurately described as random mixtures of Ti and Al according to the quasi-subregular solution model, using the SGTE data [19]. The intermetallic compound g (TiAl) phase is considered to consist of two sublattices based on a sublattice model [16,20,21]. The Gibbs energy calculated using foregoing models is shown in Fig. 5. Fig. 5 shows that all the phases a, b and g would have sufficient thermodynamic conditions for candidate of the primary phase with an undercooling of about 70 K. The liquidus temperature is about 1805 K using this model. The nucleation temperatures for a and b
7:2Rf ðqÞ a4 Tr $ $ 1 cos q d2a xL;eff DSm DTr2
(1)
where q is the wetting angle, f(q) ¼ (23cosq þ cos3q)/4, xL,eff is the effective alloy concentration, Tr ¼ T/Tm, DTr ¼ 1Tr, Sm is the molar entropy of fusion, R is gas constant, da is the average atomic 1 1/3 diameter of the solid phase taken as da¼(Wm$N1 , Wm is the 0 $r ) average weight of a mole of solid phase atoms, N0 is Avogadro's number, r is the density of the solid phase, D is the diffusivity in the undercooled melt and a is the atomic jump distance taken as 0.5 nm [24]. Thermodynamic parameters of competing phase, adopted for calculation, were given in Table 1 [15,24,25]. Grant et al. [22] found that the f(q) (catalytic efficiency) decreased with the increase of droplet diameter. During atomization, the undercooling of initial nucleation increases with decreasing droplet diameter. Thus, assuming a proportional relationship between f(q) and undercooling with typical values of q ¼ 40 for undercooling DT ¼ ~70 K (150 mm diameter droplets) and q ¼ 60 for DT ¼ ~295 K (5 mm diameter droplets [26]), respectively. Predictions based on transient nucleation theory with the case of heterogeneous nucleation (q ¼ 40 ) were consistent with experiment results well by the work of G. S. Shao [24].
. f ðqÞ ¼ 2 3 cos q þ cos3 q 4 ¼ A þ BDT
(2)
where A ¼ 2 104 and B ¼ 8 106. The diffusion coefficient can be related to the liquid viscosity h by the StokeseEinstein expression [27]:
D ¼ kB T=3pah
(3)
with h ¼ 0.925 103exp(1.164 104/RT) in present study by Zhou et al. [28]. Fig. 6(a) shows the calculated results between the incubation time and undercooling of competing phases. Based on the consideration of transient nucleation theory [29], the longer the incubation time of the competing phase, the more suppressed the formation of the competing phase is. Thus the nucleation and growth would preferably take place at the competing phase with the shortest incubation time from the undercooled melt. Two
Table 1 Thermodynamical parameters of the atomized Tie48Al alloy.
Fig. 5. Calculated Gibbs energy of liquid (L), and different competing phases (a, b and g) as a function of temperature in the Tie48Al alloy.
Phase
r (kg m3)
Tm (K)
DSm (J mol1 K1)a
a b g
3850 3850 3850
1773 1804 1729
9.196 6.821 12.324
a
Data from Refs. [15,24,25].
D.-Y. Yang et al. / Intermetallics 61 (2015) 72e79
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Fig. 7. Relative volume fraction of g phase and the cooling rate of the TieAl powders.
{111}g parallel to the basal plane of the a2 [30,31]. With the decrease of undercooling, solidification starts with the formation of primary b in droplets (D > 25 mm). At low undercooling, solidification of phase b is firstly surrounded by a and subsequently by g. The residual high temperature phase b has transformed to B2 as the final phase. The volume fraction of a2 þ g lamella and net-like structure increases from about 20% to 70% with the increase of powder size. Meanwhile, the maximum spacing of a2 þ g is almost the same of about 0.2e0.3 mm. However, the g lamella increases greatly from about 0.03 to 0.15 mm. Thus, the solidification and solid state transformations for the powders D < 25 mm can be summarized as:
L/½a þ L/½aþgs /½a2 þgs /½a2 þ g þ gs /½ða2 þ gÞ þ gt þ gs þ gg Fig. 6. Relationship between (a) incubation times of competing phases and temperature; (b) undercooling and the size of Tie48Al alloy droplets.
crossing points, Tba, and Tag (corresponding to the undercooling of 102 K and 335 K, respectively) were marked out in the plot. At the two crossing points, both competing phases have the same incubation time. At a temperature lower than Tba, b phase has the shortest incubation time. Therefore, b would be the primary phase with an undercooling less than 102 K. The undercooling between 102 K and 335 K corresponds to the region of primary phase a formation. The relationship between the undercooling and the droplet size has been calculated in our previous study. The result of calculation is given shown in Fig. 6(b). Based on analyzing of Figs. 6(a) and (b), the primary phase is obviously dependent on the droplets size D with the critical diameter of about 25 mm. When D < 25 mm, the primary phase would be phase a, when D > 25 mm, the primary phase would be phase b. It reveals that the prediction fits the experiment results well. In addition, the rapid solidification of the primary phase would release enough fraction of latent heat into the liquid to cause recalescence of the droplet, because the Biot number of droplet is about 0.01. Thus, recalescence would ensue rapidly to depress the range of primary phase a region. There would be a transformation from primary phase a to b for the droplets size less than 25 mm, and resulting in a (a þ b) primary phase region. During cooling, the primary a phase forms in droplets (D < 25 mm) and is subsequently surrounded by segregate g-TiAl which is indeed observed in the final microstructure in all cases. Further cooling results in the ordering of a to a2 with fine antiphase boundaries (APBs). From these, plates of g grow with the
The solidification and solid state transformations for the powders D > 25 mm can be summarized as:
L/½b þ L/½b þ a þ L/½b þ a þ gs /½a þ B2 þ gs /½a2 þ B2 þ gs /½ða2 þ gÞ þ gt þ gc þB2 þ gs where gs refers to g segregate, gg to g grain formed at the net-like structure, and gt to the twin-related g variants formed during the decomposition of a2, and gc to cellular g.
Fig. 8. Relative microhardness of the different size of powders (each data point corresponds to one powder).
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Fig. 9. TEM microstructures of atomized Tie48Al powder: of (a) less than 15 mm, (b) 100e150 mm.
4.2. Volume fraction of final phases Relative volume fraction of a2 and g phase was achieved with a quantitative analysis method [32]. The proportion of a2 and g phase diffraction peak area was measured in Fig. 4. Then the phase fraction was estimated by considering the Reference Intensity Ratio (RIR), as shown in Fig. 7. The solid phase cooling rate of different powders has been calculated in our previous study, and the results are shown in Fig. 7. The phase transformation is controlled by diffusion in the powders, and appears quite sensitive to the cooling rate. The average cooling rate of powder decreased with the increasing of powder size. That means the bigger powder has more diffusion time to finish the phase transformation. The volume fraction of g phase increased with the increasing of powder size from about 12% to 85%. 4.3. Mechanical property of different powders In order to gain some insight into the change in properties, microhardness measurements were carried out on the as polished sections of different powders electroplated with nickel. The results are shown in Fig. 8. The hardness increased with the increasing of powder size until the diameter reached about 50 mm (652 HV). As the diameter increased further, the hardness decreased. This variation in hardness will be correlated with the phase compositions and morphologies present in the powder. Generally, the hardness values of a2 phase are higher than g phase with rapid solidification [10]. Thus, in powders D > 50 mm the hardness declined with the obviously increasing of g phase fractions from 20% to 85%. In the fine powders ranging in size less than 15 mm, rapid solidification defectdislocations are commonly observed by TEM as shown in Fig. 9(a). With the reduction of cooling rate, the defects are obviously decreased as shown in Fig. 9(b). Thus, the non-equilibrium characteristic is weakened with the reduction of solidified defect. These would also make the hardness fall down. In powder D < 50 mm, the transformation of main phase morphologies from a2 cells to a2 þ g lamella in net-like structure might lead to the increase of hardness. The final phase B2 in larger powders would also improve the hardness of the material. 5. Conclusions In the present study, microstructure evolution of undercooled Tie48Al (at.%) droplets was analyzed with the transient nucleation theory in combination with the microstructural observation. It has
been demonstrated that the phase transformation and final fraction can be directly related to the undercooling and cooling rate determined by the droplet size. The main findings are: 1. Microstructure observations of the suppressed primary phase are in good agreement with the analysis using transient nucleation theory. The competitive primary phases a and b are dependent on the droplet size with the critical diameter of about 25 mm corresponding to an undercooling of 102 K. 2. With the decrease of undercooling, the volume fraction of a2 þ g lamella and net-like structure increases from 20% to 70%. The width of g lamella increases greatly from about 0.03 mm to 0.15 mm in constant max spacing of a2 þ g lamella structure about 0.3 mm. Segregate g-TiAl was observed in the final microstructure in all case. 3. The volume fraction of g phase is easily controlled by the cooling rate, and increased with increasing powder size from 12% to 85%. 4. The hardness increased to the increasing of powder size until 50 mm in diameter (652 HV). As the diameter increased further, the hardness decreased. Variation in hardness is correlated with the phase compositions and morphologies present in the powder. Acknowledgments The authors gratefully acknowledge the financial supports from the National Natural Science Foundation of China (Grant No. 51301157 and 51434007) and the National High Technology Research and Development Program of China 863 Program (Grant No. 2013AA031103). When the work was carried out, Dongye Yang was a visiting student under the supervision of Professor Hua-Xin Peng in ACCIS at the University of Bristol and was financially supported by the China Scholarship Council. References [1] Kim YW. Ordered intermetallic alloys, part III: gamma titanium aluminides. JOM 1994;46:30e9. [2] Yang C, Jiang H, Hu D, Huang A, Dixon M. Effect of boron concentration on phase transformation texture in as-solidified Ti44Al8NbxB. Scr Mater 2012;67: 85e8. [3] Kulkarni K, Sun Y, Sachdev A, Lavernia E. Field-activated sintering of blended elemental g-TiAl powder compacts: porosity analysis and growth kinetics of Al3Ti. Scr Mater 2013;68:841e4. [4] Wang Y, Ding H, Zhang H, Chen R, Guo J, Fu H. Microstructures and fracture toughness of Ti(43e48) Al2Cr2Nb prepared by electromagnetic cold crucible directional solidification. Mater Des 2014;64:153e9.
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Size dependent phase transformation in atomized TiAl powders Dong-Ye Yang a, b, c, Shu Guo a, b, Hua-Xin Peng c, d, Fu-Yang Cao a, b, Na Liu e, Jian-Fei Sun a, b, * a
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, PR China c Advanced Composites Centre for Innovation and Science (ACCIS), Department of Aerospace Engineering, University of Bristol, Bristol BS8 1TR, UK d School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China e Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 20 November 2014 Received in revised form 11 February 2015 Accepted 25 February 2015 Available online
The paper focuses on the phase transformation of undercooled Tie48Al (at.%) droplets atomized by high pressure gas. The microstructural evolution was analyzed using the transient nucleation theory in combination with the microstructural examinations. The primary phase, final phase volume fraction and the hardness are related to the droplets size to provide fundamental understanding. The competitive formation of the primary phases a and b are strongly controlled by the droplets size, and a critical diameter of about 25 mm was identified corresponding to an undercooling of 102 K. The final phase volume fraction of g increased with the increase of powder size from 12% to 85%. The highest hardness with the value of 652 HV was obtained for powder of 50 mm in diameter. © 2015 Elsevier Ltd. All rights reserved.
Keywords: A. Intermetallics B. Phase transformation C. Rapid solidification D. Microstructure E. Phase stability, prediction
1. Introduction TiAl has been extensively investigated as a promising candidate material for high temperature applications in aviation, automotive and power generation industries, due to its low density, relatively high specific strength and excellent corrosion resistance, as well as good oxidation and creep resistance at high temperature [1e3]. However, the poor deformability at room temperature has severely limited its further development [4e6]. To overcome this deficiency, powder metallurgy is employed to provide fine homogenous microstructures and in some cases enhanced properties to improve the performance of TiAl based alloy [7,8]. In this context, gas atomization is a prominent technique applied with many benefits for pre-alloyed powders, due to its efficiency and high performance [9e11]. During gas atomization, a liquid stream is broken up into fine spherical droplets of different sizes by the impact of a highvelocity gas flow [12,13]. The TiAl alloy droplets undergo rapid * Corresponding author. Harbin Institute of Technology, School of Materials Science and Engineering, Harbin, Heilongjiang 150001, PR China. Tel./fax: þ86 451 86413904. E-mail address: [email protected] (J.-F. Sun). http://dx.doi.org/10.1016/j.intermet.2015.02.017 0966-9795/© 2015 Elsevier Ltd. All rights reserved.
solidification processing (RSP) with the change of phase selection process toward more ductile metastable phases. The phase evolution in a non-equilibrium solidification process depends critically on the interplay between undercooling and cooling rate as a function of droplets size. Further influences under investigation include primary phase competition and suppression of solid state transformations. The final phases state, volume fraction and hardness are the important factors affecting the sintering performance of pre-alloy powders. During rapid solidification processing, the competitive nucleation of high-temperature metastable phases is seriously separated from the equilibrium solidification state. The thermodynamic and kinetic conditions were considered to fit for the guidance of phase separation under non-equilibrium solidification. Increasing the undercooling, the primary phase changes from b (A2-type bcc) to a (A3-type hcp) and g (L10-type fcc) phases [14]. When the primary phase is b-Ti, Zhou [15] has observed morphological transition from b to a in a hypo-peritectic Ti47Al53 alloy sample undercooled by 20 K. The primary or secondary a is always surrounded by g segregate. The g acts the primary phase in the achievable undercooling range of Ti44Al56 alloy [16]. The phase transformation of a / g is determined by diffusion rate which is
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quite sensitive to the cooling rate [17]. During the atomization, the transformation a / g was suppressed due to the high cooling rate of droplets about 104e106 K s1, resulting in a decrease of a2 þ g lamella structure replaced by a / a2 ordering transformation. The objective of this paper is to analyze the solidification pathways of undercooled Tie48Al droplets based on the thermodynamic and kinetic theories. Relationships are established between the primary phase, volume fraction, hardness and the gas atomized Tie48Al droplets size to provide fundamental understanding of the subsequent powder metallurgy (PM) process. 2. Experimental procedure The powders of TiAl alloy (nominal composition: Tie48Al, at.%) were prepared by high pressure gas atomization in an argon spray
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chamber environment, and the TiAl alloy was melted by levitation melting technique in order to achieve composition homogeneity and purity (140 wt.ppm N, 600 wt.ppm O). These powders were firstly classified into 5 different size groups, ranging from 5 to 150 mm in diameter. The classification was carried out by screening and a sedimentation technique [18]. The microstructure of the powders was characterized using scanning electron microscopy (SEM), X-ray diffraction (XRD) and transmission electron microscopy (TEM). For SEM observations, the powders were mounted in epoxy, polished, etched with Kroll solution 4 vol. % HNO3 þ 2 vol. % HF þ 94 vol. % H2O and coated with conductive layer. Transmission electron microscopy (TEM) was used to determine the phases and the microstructure of the powders in addition to SEM and optical microscopy. The TEM specimens were prepared by
Fig. 1. Microstructures of atomized Tie48Al powder (less than 15 mm): (a) Featureless surface (SEM), (b) TEM image of the powder, (c) Corresponding SAED pattern of the matrix a2 phase, zone axis ½0001a2 , (d) Lamella structure, corresponding SAED pattern of the g phase, zone axis ½110g , (e) Corresponding SAED pattern of a phase, zone axis ½1120a , (f) Mixed a and a2 electron diffraction, from the ½0110a =½0001a2 zone axis.
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the FIB-TEM foil preparation technique and a conventional plating technique [14]. For plating technique, the powders were firstly placed loosely on a copper sheet substrate. This copper sheet was then electroplated with nickel, using a low current density (4e5 mA cm2) to ensure good adherence of the nickel plating to the powders. The plating was about 650 mm thick. The copper sheet was removed, and the foils were then subjected to polish, dimple and ion mill. 3. Experimental results The microstructures and phases developed in the Tie48Al at.% powders change with the rapid solidification regime such as the undercooling and cooling rates. The regime is strongly dependent on the size of powders. Thus the results are presented for powder size ranges representing as fine (less than 15 mm), medium (40e70 mm) and large (100e150 mm). In the fine powders, a featureless surface morphology is observed by SEM ranging in size less than 15 mm as shown in Fig. 1(a). Examination of the powders cross-sections generally shows the presence of cellular and lamella structure morphologies by TEM in Fig. 1(b). The volume fraction of cellular microstructure in powders is about 80% with an average cellular spacing of 1e3 mm. Electron diffraction pattern shows the matrix phase of the cells to be a2, from the ½0001a2 zone axis in Fig. 1(c). Fig. 1(b) shows some region containing only a2 (marked as region A) while region B contains the usual a2 þ g lamella structure with a spacing of 0.1e0.25 mm, with higher magnification micrograph as shown in Fig. 1(d). However, the grey phase of g lamella is fine and limited precipitation (less than 0.03 mm). The segregate
structures obtained in the a2 grain boundaries consist of the lamella and feather-like phase, and the electron diffraction pattern from the ½110g zone axis shows the phase to be g in Fig. 1(d). Fig. 1(e) presents corresponding electron diffraction pattern composed of plates of a significantly different hexagonal phase, and this is disordered a phase formed from the liquid, from the ½1120a zone axis reported previously by McCullough [14]. Fig. 1(f) shows a mixed a and a2 electron diffraction pattern from the ½0110a =½0001a2 zone axis, and a direct evidence of suppression in the ordering a / a2. Thus the primary phase formed from the liquid is a phase. In addition, energy dispersive spectroscopy (EDS) analysis an average aluminum concentrations for the cells and segregate g phase areas of 47.9 at.% and 50.2 at.%, respectively, indicating segregation was suppressed owing to the result of rapid solidification. For powders of size from 40 to 70 mm, the surface morphology exhibits a dendrite appearance with a little short dendrite arms as shown in Fig. 2(a). The grains are rather small (about 15 mm) and the growth is almost equiaxed to show a clear hexagonal or polygon pattern. Examination of the powders cross-sections generally shows the presence of concentric liquid/solid interface solidification geometry due to the multi-nucleation events taking place in Fig. 2(b). The volume fraction of a2 þ g lamella structure increases to about 30% with an average spacing of 0.3 mm and the grey phase of g lamella is about 0.1 mm. From the TEM observations this can largely be attributed to the suppressed decomposition of a2 into the a2 þ g. The net-like structure shows the presence of fine g phase grains and a2 þ g lamella (less than 0.02 mm) with twins structures
Fig. 2. Microstructures of atomized Tie48Al powder (40e70 mm): SEM views of (a) Dendrite surface and (b) Equiaxed dendritic morphology, TEM views of (c) a2 þ g lamella structure with twins structure from net-like structure with corresponding SAED pattern showing ½1210a2 ==½110g , (d) Fine g phase grain.
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morphologies as indicated by TEM images in Figs. 2(c) and (d), respectively. Fig. 2(c) shows the selected area diffraction patterns of the net-like structure phase, which proved to be the phase a2 þ g with ½1210a2 ==½110g for the transformed region. In the large powders, the equiaxed dendrite surface morphology with development of long dendrite arms ranging in size from 100 to 150 mm is observed by SEM as shown in Fig. 3(a). The equiaxed dendrites are similar to those seen in conventional solidification in which the primary phase was b-Ti. Fig. 3(b) shows cross-sectional view of a large powder, where the dendrite arms are almost orthogonal to one another. Thus, it seems more likely to be the characteristic growth of the cubic b phase from the liquid. Electron diffraction (Fig. 3(d)) of the selected area in Fig.3(c) reveals the presence of B2 phase of the transformation of parent b, from the ½001B2 zone axis. Hence, in the large powders there is evidences for the primary phase forming from the liquid is b phase.
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The volume fraction of a2 þ g lamella and net-like structure increases to about 70% with sole a2 surrounded by g phase in Fig. 3(c). The spacing of a2 þ g lamella is about 0.3 mm and the grey phase of g lamella is about 0.15 mm as shown in Fig. 3(e). EDS analysis reveals the average composition for the dendrite and segregate regions to be 47.7 at.% Al and 51.1% Al, respectively. The average aluminum concentration is slightly higher in comparison to fine powders. Fig. 4 shows the X-ray diffraction analyses of powders of different sizes at room temperature (298 K). The spectrum contains a mixture of a2 and g peaks, but is dominated by the large peaks of ð201Þa2 and (111)g near 41 and 39 , respectively. With increasing powder size, the intensity of the a2 peaks reduced greatly with the loss of some previous peaks. Concurrently new g signals corresponding to (001), (110) and (222) appeared at 22 , 32 and 85 , respectively. In addition, the peaks of (202)g and (220)g near 65 are
Fig. 3. Microstructures of atomized Tie48Al powder (100e150 mm): SEM views of (a) Dendrite surface and (b) Orthogonal symmetry dendrites, TEM views of (c) sole a2 surrounded by g phase and B2 phase with corresponding SAED pattern (d), zone axis ½001B2 , (e) a2þg lamella structure.
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are very close, and it is limited to predict the primary phase under non-equilibrium conditions. Thus, further kinetic consideration of incubation time has to be performed. Previous studies have concluded that the nucleation mechanism of gas atomized droplets undercool below the alloy liquidus temperature is the heterogeneous nucleation [22,23]. Heterogeneous nucleation generally dictates the practical limit of metastability which is dependent on the transient time. Temperature vs transient time curves can be calculated and used to predict the selection of competing phases during rapid solidification. The nucleation behavior of competing phases in the undercooled TiAl droplets could be evaluated by the transient nucleation theory proposed by Shao and Tsakiropoulos [24]. The incubation time t can be expressed as:
t¼
Fig. 4. XRD spectra for atomized Tie48Al alloy powders of different sizes.
overlapped due to the decrease of lattice constant by rapid solidification. 4. Discussion 4.1. Thermodynamic analysis of phase formation In order to analyze the primary phase of the atomized TiAl droplets, the disordered solution phase, liquid (L), a (Ti) and b (Ti) and the intermetallic compound g (TiAl) phase are considered in the calculation of the Gibbs energy of the Tie48Al, at.% alloy. The Gibbs energy of the present phases has to be expressed as analytical functions of composition, temperature and pressure. The liquid (L), a (Ti) and b (Ti) phases could be accurately described as random mixtures of Ti and Al according to the quasi-subregular solution model, using the SGTE data [19]. The intermetallic compound g (TiAl) phase is considered to consist of two sublattices based on a sublattice model [16,20,21]. The Gibbs energy calculated using foregoing models is shown in Fig. 5. Fig. 5 shows that all the phases a, b and g would have sufficient thermodynamic conditions for candidate of the primary phase with an undercooling of about 70 K. The liquidus temperature is about 1805 K using this model. The nucleation temperatures for a and b
7:2Rf ðqÞ a4 Tr $ $ 1 cos q d2a xL;eff DSm DTr2
(1)
where q is the wetting angle, f(q) ¼ (23cosq þ cos3q)/4, xL,eff is the effective alloy concentration, Tr ¼ T/Tm, DTr ¼ 1Tr, Sm is the molar entropy of fusion, R is gas constant, da is the average atomic 1 1/3 diameter of the solid phase taken as da¼(Wm$N1 , Wm is the 0 $r ) average weight of a mole of solid phase atoms, N0 is Avogadro's number, r is the density of the solid phase, D is the diffusivity in the undercooled melt and a is the atomic jump distance taken as 0.5 nm [24]. Thermodynamic parameters of competing phase, adopted for calculation, were given in Table 1 [15,24,25]. Grant et al. [22] found that the f(q) (catalytic efficiency) decreased with the increase of droplet diameter. During atomization, the undercooling of initial nucleation increases with decreasing droplet diameter. Thus, assuming a proportional relationship between f(q) and undercooling with typical values of q ¼ 40 for undercooling DT ¼ ~70 K (150 mm diameter droplets) and q ¼ 60 for DT ¼ ~295 K (5 mm diameter droplets [26]), respectively. Predictions based on transient nucleation theory with the case of heterogeneous nucleation (q ¼ 40 ) were consistent with experiment results well by the work of G. S. Shao [24].
. f ðqÞ ¼ 2 3 cos q þ cos3 q 4 ¼ A þ BDT
(2)
where A ¼ 2 104 and B ¼ 8 106. The diffusion coefficient can be related to the liquid viscosity h by the StokeseEinstein expression [27]:
D ¼ kB T=3pah
(3)
with h ¼ 0.925 103exp(1.164 104/RT) in present study by Zhou et al. [28]. Fig. 6(a) shows the calculated results between the incubation time and undercooling of competing phases. Based on the consideration of transient nucleation theory [29], the longer the incubation time of the competing phase, the more suppressed the formation of the competing phase is. Thus the nucleation and growth would preferably take place at the competing phase with the shortest incubation time from the undercooled melt. Two
Table 1 Thermodynamical parameters of the atomized Tie48Al alloy.
Fig. 5. Calculated Gibbs energy of liquid (L), and different competing phases (a, b and g) as a function of temperature in the Tie48Al alloy.
Phase
r (kg m3)
Tm (K)
DSm (J mol1 K1)a
a b g
3850 3850 3850
1773 1804 1729
9.196 6.821 12.324
a
Data from Refs. [15,24,25].
D.-Y. Yang et al. / Intermetallics 61 (2015) 72e79
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Fig. 7. Relative volume fraction of g phase and the cooling rate of the TieAl powders.
{111}g parallel to the basal plane of the a2 [30,31]. With the decrease of undercooling, solidification starts with the formation of primary b in droplets (D > 25 mm). At low undercooling, solidification of phase b is firstly surrounded by a and subsequently by g. The residual high temperature phase b has transformed to B2 as the final phase. The volume fraction of a2 þ g lamella and net-like structure increases from about 20% to 70% with the increase of powder size. Meanwhile, the maximum spacing of a2 þ g is almost the same of about 0.2e0.3 mm. However, the g lamella increases greatly from about 0.03 to 0.15 mm. Thus, the solidification and solid state transformations for the powders D < 25 mm can be summarized as:
L/½a þ L/½aþgs /½a2 þgs /½a2 þ g þ gs /½ða2 þ gÞ þ gt þ gs þ gg Fig. 6. Relationship between (a) incubation times of competing phases and temperature; (b) undercooling and the size of Tie48Al alloy droplets.
crossing points, Tba, and Tag (corresponding to the undercooling of 102 K and 335 K, respectively) were marked out in the plot. At the two crossing points, both competing phases have the same incubation time. At a temperature lower than Tba, b phase has the shortest incubation time. Therefore, b would be the primary phase with an undercooling less than 102 K. The undercooling between 102 K and 335 K corresponds to the region of primary phase a formation. The relationship between the undercooling and the droplet size has been calculated in our previous study. The result of calculation is given shown in Fig. 6(b). Based on analyzing of Figs. 6(a) and (b), the primary phase is obviously dependent on the droplets size D with the critical diameter of about 25 mm. When D < 25 mm, the primary phase would be phase a, when D > 25 mm, the primary phase would be phase b. It reveals that the prediction fits the experiment results well. In addition, the rapid solidification of the primary phase would release enough fraction of latent heat into the liquid to cause recalescence of the droplet, because the Biot number of droplet is about 0.01. Thus, recalescence would ensue rapidly to depress the range of primary phase a region. There would be a transformation from primary phase a to b for the droplets size less than 25 mm, and resulting in a (a þ b) primary phase region. During cooling, the primary a phase forms in droplets (D < 25 mm) and is subsequently surrounded by segregate g-TiAl which is indeed observed in the final microstructure in all cases. Further cooling results in the ordering of a to a2 with fine antiphase boundaries (APBs). From these, plates of g grow with the
The solidification and solid state transformations for the powders D > 25 mm can be summarized as:
L/½b þ L/½b þ a þ L/½b þ a þ gs /½a þ B2 þ gs /½a2 þ B2 þ gs /½ða2 þ gÞ þ gt þ gc þB2 þ gs where gs refers to g segregate, gg to g grain formed at the net-like structure, and gt to the twin-related g variants formed during the decomposition of a2, and gc to cellular g.
Fig. 8. Relative microhardness of the different size of powders (each data point corresponds to one powder).
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Fig. 9. TEM microstructures of atomized Tie48Al powder: of (a) less than 15 mm, (b) 100e150 mm.
4.2. Volume fraction of final phases Relative volume fraction of a2 and g phase was achieved with a quantitative analysis method [32]. The proportion of a2 and g phase diffraction peak area was measured in Fig. 4. Then the phase fraction was estimated by considering the Reference Intensity Ratio (RIR), as shown in Fig. 7. The solid phase cooling rate of different powders has been calculated in our previous study, and the results are shown in Fig. 7. The phase transformation is controlled by diffusion in the powders, and appears quite sensitive to the cooling rate. The average cooling rate of powder decreased with the increasing of powder size. That means the bigger powder has more diffusion time to finish the phase transformation. The volume fraction of g phase increased with the increasing of powder size from about 12% to 85%. 4.3. Mechanical property of different powders In order to gain some insight into the change in properties, microhardness measurements were carried out on the as polished sections of different powders electroplated with nickel. The results are shown in Fig. 8. The hardness increased with the increasing of powder size until the diameter reached about 50 mm (652 HV). As the diameter increased further, the hardness decreased. This variation in hardness will be correlated with the phase compositions and morphologies present in the powder. Generally, the hardness values of a2 phase are higher than g phase with rapid solidification [10]. Thus, in powders D > 50 mm the hardness declined with the obviously increasing of g phase fractions from 20% to 85%. In the fine powders ranging in size less than 15 mm, rapid solidification defectdislocations are commonly observed by TEM as shown in Fig. 9(a). With the reduction of cooling rate, the defects are obviously decreased as shown in Fig. 9(b). Thus, the non-equilibrium characteristic is weakened with the reduction of solidified defect. These would also make the hardness fall down. In powder D < 50 mm, the transformation of main phase morphologies from a2 cells to a2 þ g lamella in net-like structure might lead to the increase of hardness. The final phase B2 in larger powders would also improve the hardness of the material. 5. Conclusions In the present study, microstructure evolution of undercooled Tie48Al (at.%) droplets was analyzed with the transient nucleation theory in combination with the microstructural observation. It has
been demonstrated that the phase transformation and final fraction can be directly related to the undercooling and cooling rate determined by the droplet size. The main findings are: 1. Microstructure observations of the suppressed primary phase are in good agreement with the analysis using transient nucleation theory. The competitive primary phases a and b are dependent on the droplet size with the critical diameter of about 25 mm corresponding to an undercooling of 102 K. 2. With the decrease of undercooling, the volume fraction of a2 þ g lamella and net-like structure increases from 20% to 70%. The width of g lamella increases greatly from about 0.03 mm to 0.15 mm in constant max spacing of a2 þ g lamella structure about 0.3 mm. Segregate g-TiAl was observed in the final microstructure in all case. 3. The volume fraction of g phase is easily controlled by the cooling rate, and increased with increasing powder size from 12% to 85%. 4. The hardness increased to the increasing of powder size until 50 mm in diameter (652 HV). As the diameter increased further, the hardness decreased. Variation in hardness is correlated with the phase compositions and morphologies present in the powder. Acknowledgments The authors gratefully acknowledge the financial supports from the National Natural Science Foundation of China (Grant No. 51301157 and 51434007) and the National High Technology Research and Development Program of China 863 Program (Grant No. 2013AA031103). When the work was carried out, Dongye Yang was a visiting student under the supervision of Professor Hua-Xin Peng in ACCIS at the University of Bristol and was financially supported by the China Scholarship Council. References [1] Kim YW. Ordered intermetallic alloys, part III: gamma titanium aluminides. JOM 1994;46:30e9. [2] Yang C, Jiang H, Hu D, Huang A, Dixon M. Effect of boron concentration on phase transformation texture in as-solidified Ti44Al8NbxB. Scr Mater 2012;67: 85e8. [3] Kulkarni K, Sun Y, Sachdev A, Lavernia E. Field-activated sintering of blended elemental g-TiAl powder compacts: porosity analysis and growth kinetics of Al3Ti. Scr Mater 2013;68:841e4. [4] Wang Y, Ding H, Zhang H, Chen R, Guo J, Fu H. Microstructures and fracture toughness of Ti(43e48) Al2Cr2Nb prepared by electromagnetic cold crucible directional solidification. Mater Des 2014;64:153e9.
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