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Evolution of the microstructure and hardness of the Ti–Si alloys during high temperature heat-treatment
Journal of Alloys and Compounds 479 (2009) 246–251
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Evolution of the microstructure and hardness of the Ti–Si alloys during high temperature heat-treatment Yongzhong Zhan a,∗ , Xinjiang Zhang a , Jing Hu a , Qinghua Guo a , Yong Du b a b
Laboratory of Nonferrous Metal Materials and New Processing Technology, Ministry of Education, Guangxi University, Nanning, Guangxi 530004, PR China State Key Laboratory of Powder Metallurgy, Central South University, Changsha, Hunan 410083, PR China
a r t i c l e
i n f o
Article history: Received 9 December 2008 Received in revised form 5 January 2009 Accepted 7 January 2009 Available online 20 January 2009 Keywords: Metals and alloys Microstructure X-ray diffraction Scanning electron microscopy Optical microscopy
a b s t r a c t Samples of Ti–8Si, Ti–13.67Si and Ti–23S (at.%) alloys were prepared by vacuum ar melting technique and heat-treated for 2 h at 1373 K and 1473 K in a vacuum environment. The as-cast and annealed samples were examined by optical microscope (OM), X-ray diffraction (XRD), scanning electron microscope (SEM), energy dispersive X-ray analysis (EDX) and hardness test. During high temperature heat-treatment, the quasi-continuous network-shaped Ti5 Si3 + Ti phases in the as-cast Ti–8Si alloy finally converted to near-equiaxed Ti5 Si3 grains dispersed in the Ti matrix. For Ti–13.67Si alloy, the Ti5 Si3 grains finally also transformed to the near-equiaxed ones compared with the lamellar grains in the as-cast alloy. Moreover, compared with the as-cast alloy, Vickers hardness of Ti–8Si and Ti–13.67Si alloys was increased by the annealing. However, the microstructure of the Ti–23Si alloy could not change significantly. So, adversely Vickers hardness decreased after high temperature heat-treatment. The microstructural evolution of Ti–Si alloys is related with the diffusion of Si atoms in the Ti/Ti5 Si3 interface during high temperature heattreatment producing as much as possible homogeneous distribution of Ti5 Si3 particles in a continuous ␣-Ti matrix. The task was not achievable for Ti–23Si alloy because of the stability of the primary Ti5 Si3 phase particle during heat-treatment. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Titanium and titanium alloys has been intensively studied due to a lot of practical desirable properties, i.e. high strength, specific modulus and good corrosion resistance, etc. [1–3]. However, some major inherent problems like severe metal mould reaction, low castability and high-energy consumption of traditional casting titanium alloys have greatly impeded the development of these materials. As a new titanium casting alloy, Ti–Si eutectic alloy has some excellent casting properties, such as low melting point, narrow crystallization range and good fluidity, which is similar to the extensively applied Fe–C eutectic alloy and Al–Si eutectic alloy. Moreover, it consists of a ductile phase (␣-Ti, which is the matrix) and a brittle phase (Ti5 Si3 , acts as the reinforcement) [4–10]. This eutectic alloy has been considered to be the thirdgeneration eutectic alloy. Heat-treatments are usually adopted to increase the ductility and strength of titanium alloys by modifying their solid-state phase phase change behavior and microstructure. It can improve the mechanical properties of bulk materials that are related to the microstructure, processing history and heattreatment procedures. Some works conducted to the knowledge
∗ Corresponding author. Tel.: +86 771 3272311; fax: +86 771 3233530. E-mail address: [email protected] (Y. Zhan). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.01.017
of microstructure evolution and deformation mechanism of Ti alloys during hot deformation under different processing and heattreatment procedures [11–14]. These previous studies have shown that the thermal treatment of the Ti–6Al–4V alloy carried out at temperature above the  transus temperature leads to a ‘lamellar’ microstructure colonized plate-like ␣ at a low cooling rate, basket-weave at an intermediate cooling rate, Widmanstätten at a high cooling rate, and martensite after water-quenching. When processing is carried out in the (␣ + ) phase field, a much finer (␣ + ) structure can be obtained, because the primary ␣ phase limits the growth of the  phase [15,16]. In addition to processing temperature, other hot working parameters, such as strain and strain rate, also affect the microstructure, e.g. the volume fractions of the ␣ and  phases, the phase size and the ‘lamellar’ dimensions of the ␣ phase. The aspect ratio of ␣ lamellar phase is also a very important factor that influences the mechanical properties of the Ti–6Al–4V alloy [17–19]. These studies on microstructure and property of some Ti alloys during different heattreatment procedures are of great importance in exploring Ti–Si alloys. According to the currently accepted Ti–Si phase diagram (Fig. 1) [20], the reported stable phases in Ti-rich region are: -Ti, ␣-Ti, Ti3 Si, Ti5 Si3 and liquid (L). The compound Ti3 Si is formed from both eutectoid and peritectoid transformations (-Ti ↔ Ti3 Si + ␣Ti; -Ti + Ti5 Si3 ↔ Ti3 Si). Additionally, an eutectic reaction of (Ti–Si)
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247
2. Experimental procedure The stating materials used in this work were pure Ti (99.99 wt.%) and Si (99.99 wt.%). The alloy samples (all weighting 25 g) were arc-melted in a watercooled copper cast with a non-consumable tungsten electrode under pure argon atmosphere. Each sample was melted three times and turned around after melting for better homogeneity. For examining the effect of heat-treatment temperature on microstructure and mechanical properties, the alloy samples was heat-treated at 1373 K and 1473 K for 2 h at each temperature, and then cooled down to the room temperature with the furnace. The annealing temperature was determined according to Ti–Si binary phase diagram. Table 1 shows the nominal compositions of the alloys and the process of high temperature treatment. The as-cast and annealed samples were characterized via scanning electron microscope (SEM), optical microscopy (OM) and X-ray diffraction (XRD). The SEM images were obtained in a Hitachi S-3400N scanning electron microscope equipped with energy dispersive X-ray analysis (EDX). For the XRD experiment, the samples were performed on a Rigaku D/Max 2500 V diffractometer with Cu-K␣ radiation and graphite monochromator operated at 40 kV, 250 mA (20◦ ≤ 2 ≤ 60◦ ). The Materials Data software Jade 5.0 [21] and Powder Diffraction File (PDF release 2002) were used for phase analysis. Average hardness of the alloys were measured using model HV-30 Vickers microhardness tester with a test load of 3 kg and a load-dwell time of 30 s. More than eight indentations were performed on each samples to obtain the mean value. Fig. 1. Ti-rich region of the Ti–Si phase diagram.
3. Results and discussions L ↔ Ti5 Si3 + -Ti occurs at a temperature of 1603 K and 8.5 wt.% Si in Ti-rich region of the Ti–Si phase diagram. In order to improve the combined properties of Ti–Si alloys and step up their practical applications, the present work did research work to examine the evolution of the microstructure and hardness during high temperature processing of Ti–Si alloys, and then discussed their evolution mechanism from a microstructural viewpoint.
3.1. As-cast microstructure Fig. 2 shows the optical micrographs of the as-cast Ti–Si alloys. XRD data (Fig. 3) and EDX analysis have indicated the presence of Ti5 Si3 phases in all alloy samples. The micrograph of Ti–8Si hypoeutectic alloy (Fig. 2a) shows that the Ti5 Si3 + Ti lamellar particles distribute the domain boundaries of the mas-
Table 1 Compositions of the Ti–Si alloys and process of high temperature treatment in the present work. Alloys
Alloy type
Composition (at.% Si)
Holding temperature (K)
Holding time (h)
Ti–8Si Ti–13.67Si Ti–23Si
Hypoeutectic Eutectic Hypereutectic
8.0 13.67 23.0
1373 1373 1373
2 2 2
1473 1473 1473
Fig. 2. As-cast optical micrographs of Ti–Si alloy samples: (a) Ti–8Si; (b) Ti–13.67Si; (c) Ti–23Si.
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Fig. 3. XRD patterns of the as-cast Ti–Si alloy samples.
sive Ti proeutectic phase forming continuous or quasi-continuous network-shaped microstructure. In the micrograph of Ti–13.67Si eutectic alloy (Fig. 2b), a typical Ti + Ti5 Si3 eutectic microstructure is observed. A higher magnification should reveal peritectoid Ti and Ti5 Si3 plus Ti3 Si particles, eventually some finer eutectoid lamella. Fig. 2c shows the microstructural feature of Ti–23Si hypereutectic alloy. Some massive proeutectic Ti5 Si3 phase domain and typical Ti + Ti5 Si3 eutectic structure are observed in this alloy. Based on above analytical result, it is clear that the phase distribution of the as-cast Ti–13.67Si eutectic alloy is more homogeneous than that of the hypoeutectic and hypereutectic alloys. With the increase of Si content, the content of Ti5 Si3 phase gradually rises as indicated from the intensity of Ti5 Si3 of the XRD pattern in Fig. 3, which is the same result provided by the microstructure analysis. These results confirm the eutectic reaction of L ↔ Ti5 Si3 + -Ti at approximately 13.67 at.% Si as predicted in Fig. 1. The high temperature phase -Ti is observed in Ti–8Si
and Ti–13.67Si alloys from the XRD data, suggesting that noticeable amounts of the -Ti phase have been retained in the as-cast alloys under the rapid cooling rate of the arc furnace. Moreover, Fig. 2(a–c) shows that Ti–13.67 Si alloy exhibits smaller and more homogeneous grain-sized than Ti–8Si and Ti–23Si alloys. According to the SEM micrographs of this eutectic alloy (Fig. 4), it is observed that eutectic cells consist of some clubbed and lamellar Ti5 Si3 grains. The dendritic microstructure is possibly produced by reason of constitutional supercooling resulted from the rapid cooling process. It is clear that the long clubbed and lamellar phases are oriented differently in adjacent eutectic cells, which are produced by the peritectoid reaction -Ti + Ti5 Si3 (or Ti3 Si at lower temperature). From Fig. 1, the silicide in the equilibrium microstructure of Ti–Si should be Ti3 Si. However, the silicide is analyzed to be Ti5 Si3 by the result of XRD in Fig. 3. This is due to the fact that the formation of Ti3 Si would be a very slow solid-state phase change process at eutectoid temperature. Therefore, its formation was quite hindered (but not at all because the eutectoid ␣-Ti is present in the XRD pattern) in the present experimental condition, and the amount of peritectoid Ti3 Si formed should be too short to be detected by XRD. 3.2. Microstructural evolution at high temperature The optical micrographs of the samples after annealing at different temperature show the microstructure evolution of Ti–Si alloys during heat-treatment process. Fig. 5 shows the XRD result of Ti–8Si heat-treated hypoeutectic samples. Compared with the as-cast alloy, the phase component of this alloy did not altered during different high temperature processing, namely Ti5 Si3 , ␣-Ti and -Ti. However, due to quick atomic diffusion at high temperature, some continuous network-shape Ti5 Si3 phase arrays at Ti phase domain boundaries of the as-cast Ti–8Si alloy changed by means of coalescence mechanism during annealing at 1373 K, as is shown
Fig. 4. SEM micrographs of the as-cast Ti–13.67Si alloy: (a) low magnification, showing the different orientations of adjacent colonies; (b) low magnification, showing the clubbed and lamellar Ti5 Si3 grains; (c) high magnification, showing the dendritic microstructure.
Y. Zhan et al. / Journal of Alloys and Compounds 479 (2009) 246–251
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Fig. 5. XRD patterns of Ti–8Si heat-treated sample.
Fig. 7. XRD patterns of the Ti–13.67Si heat-treated alloy.
in Fig. 6a. At the annealing temperature of 1473 K, the networkshape Ti5 Si3 phases of the as-cast alloy disappeared entirely by some initial re-solution of Ti5 Si3 into the -Ti matrix, followed by Ti5 Si3 grain-growth and re-precipitation of Ti5 Si3 from Ti matrix onto Ti5 Si3 grains. By this way, more homogeneous distribution of the phases in the final microstructure is provided. Spherodized or near-equiaxed Ti5Si3 phases (Fig. 6b) with Ti5 Si3 grain diameter in the range of 15–25 m after annealing at 1473 K resulted. For Ti–13.67Si eutectic alloy, XRD data of the annealed sample shows that the phase component of this alloy is Ti5 Si3 and ␣-Ti. Moreover, the high temperature phase -Ti that exists in the as-cast alloy disappears during high temperature processing, as is shown in Fig. 7. According to the optical micrographs of annealed sample (Fig. 8), the long clubbed and dendritic lamellar phase disappear after annealing at 1373 K and 1473 K for 2 h, compared with that of the as-cast microstructure. Moreover, it is obvious that after annealing the grains are much finer. From SEM images of the annealed sample, it can be seen that after isothermally treated at 1373 K for 2 h, most of the long clubbed and lamellar eutectic Ti5 Si3 grains have transformed to equiaxed shape as well as some short clubbed ones (Fig. 9a). However, when the isothermal holding temperature increases to 1473 K for 2 h, the Ti5 Si3 phases almost completely turned into symmetrically equiaxed or near-equiaxed grains (Fig. 9b) with diameter in the range of 5–30 m, confirming the coalescence mechanism, which provided a continuous matrix of ␣-Ti phase. According the XRD analysis of the annealed Ti–23Si sample (Fig. 10), the phase component is the same with that of the as-cast alloy, i.e. Ti5 Si3 and ␣-Ti. Compared with the as-cast microstructure, the abundant massive proeutectic Ti5 Si3 phases still exist in the Ti–23Si hypereutectic alloy after annealed at different high temperature, as is shown in Fig. 11. However, the Ti5 Si3 phase in the typical Ti + Ti5 Si3 eutectic structure of the as-cast alloy has been
spheroidized and coarsened after annealed at 1373 K for 2 h. This phase is much bulky after 1473 K heat-treatment. From the Ti–Si binary phase diagram, at 1373 K and 1473 K, it is clear that the abundant massive proeutectic Ti5 Si3 phases in the as-cast alloy may not in heat-treatment process. 3.3. Analysis of evolution mechanism The free energy change of a system taking place in a solid-state nucleation and growth of the new phase has been defined as [22]: F = −Vfv + s + fε − fd where Vfv is the chemistry free energy for producing V is the volume new phase, s is the surface energy, fε is the elastic strain energy consumed by new phase and fd is the free energy arisen from crystal defect. When the volume of Ti5 Si3 phases is phases is a fixed value, the lamellar Ti5 Si3 grain possesses higher surface than the spherical one and so spontaneously transforms into the later shape in order to reduce free energy of the system. The spheroidization transformation of the Ti5 Si3 grains requires atomic diffusion and, the diffusion of Si atoms inside Ti5 Si3 is accelerated at high temperature (1373 K and 1473 K). At the same time, a mass of defects such as cavity and dislocation contribute to the diffusion of Si atoms. On heating the as-cast alloys above 866 ◦ C, there will be resolution of Ti3 Si by means of the first peritectoid reaction; above 1170 ◦ C the second peritectoid reaction causes the re-solution of Ti5 Si3 . It is well known that the smaller the radius titanium silicide the bigger is its solubility into the titanium phase [22,23]. The lamellar Ti5 Si3 phases (Fig. 4) have large surface free energy and there is natural trend to coalescence in order to reduce surface free energy. This was the reason by which high temperature (1373 K or 1473 K for 2 h) heat-treatment of Ti–8Si and Ti–13.67Si alloys conducted to near-equiaxed Ti5 Si3 grains. In addition, it is found that
Fig. 6. Optical micrographs of Ti–8Si hypoeutectic alloy sample after high temperature processing at (a) 1373 K and (b) 1473 K.
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Fig. 8. Optical micrographs of Ti–8Si alloy sample after high temperature processing at (a) 1373 K and (b) 1473 K.
Fig. 9. SEM micrographs of Ti–13.67Si eutectic alloy sample during high temperature processing: (a) 1373 K and (b) 1473 K.
spherical new phase. Moreover, the diffusion coefficient is temperature dependent according to Arrhenius equation and so its value at 1473 K is bigger than that at 1373 K, so the growth of spherical new phase at 1473 K is accelerated. That is the reason for the formation of much more bulky grains in the samples that annealed at 1473 K. However, there is contribution from the peritectoid reactions, i.e. at 1373 K, -Ti (p) + Ti3 Si = -Ti (r), and at 1473 K, -Ti (p) + Ti5Si = -Ti (r), where p: poor in Si and r: rich in Si content. 3.4. Hardness change during heat-treatment
Fig. 10. XRD patterns of the Ti–23Si heat-treated alloy.
the thermal treatment of Ti–Si alloys at 1373 K was not enough to provide structure separation and spheroidization of Ti5 Si3 grains. Apparently, the incomplete structure separation and spheroidization would possibly affect the final coarsening. From the microstructures of Ti–8Si and Ti–13.67Si alloys shown in Figs. 6 and 8, the grains of samples annealed at 1473 K are coarser than that of the annealed sample at 1373 K. At 1373 K or 1473 K, the diffusion coefficient (D) is the main factor for the growth of
In order to evaluate the effect of high temperature processing on the mechanical properties of Ti–Si alloys, Vickers hardness measurements were carried out on all samples, as is presented in Table 2. More than eight indentations were performed on each sample in its three different states. The indent sizes are in the range of 100–125 m, larger than the grain-sizes in the microstructures analyzed (as-cast and heat-treated) to obtain the mean value. Compared with the as-cast alloys, the Vickers hardness of Ti–Si alloys had been obviously altered after high temperature heat-treatment. For the Ti–8Si alloy, hardness increments of 141 HV and 130 HV were observed by annealing at 1373 K and 1473 K, respectively. Moreover, the hardness value of Ti–13.67Si alloy was increased from 379 HV for the as-cast state to 410 HV and 435 HV after annealing at 1373 K
Table 2 Average Vickers hardness of as-cast and high temperature treated Ti–Si alloy samples. Alloys Hardness (HV)
Ti–8Si
Ti–13.67Si
Ti–23Si
As-cast
1373 K
1473 K
As-cast
1373 K
1473 K
As-cast
1373 K
1473 K
376
517
506
379
410
435
458
387
365
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Fig. 11. Optical micrographs of Ti–23Si alloy sample during high temperature processing: (a) 1373 K and (b) 1473 K.
and 1473 K respectively. However, the hardness of Ti–23Si alloy was decreased by high temperature heat-treatment. The as-cast hardness of Ti–Si alloy was altered due to the variation of grain-size and constituent particles. In addition, the near-equiaxed Ti5 Si3 particles that distribute homogeneously in the continuous Ti matrix are propitious to the improvement of both hardness and strength. The above mechanisms may contribute to the increase of Vickers hardness of the hypoeutectic and eutectic Ti–Si alloys after high temperature heat-treatment and long diffusion time.
Acknowledgements
4. Conclusion
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The microstructural evolution of Ti–Si alloys that took place during high temperature heat-treatment of the as-cast alloys is described in the following. The quasi-continuous network-shape Ti5 Si3 + Ti intergranular phases of the as-cast Ti–8Si alloy transformed to near-equiaxed grains of Ti5 Si3 by annealing at high temperature, with grain-sizes in the range from 15 m to 25 m. In turn, the lamellar microstructure of Ti–13.67Si as-cast alloy transformed to Ti5 Si3 grains (5–30 m in size) near-equiaxed and homogeneously dispersed in the ␣-Ti matrix. Moreover, the Vickers hardness of Ti–8Si and Ti–13.67Si alloys was increased by high temperature heat-treatment. Unlikely, the microstructures of Ti–23Si alloy had not been altered significantly, and the hardness was decreased during high temperature heat-treatment. The microstructural evolution of Ti–Si alloys is related to the diffusion of Si atoms in the Ti/Ti5 Si3 particles interface when the alloys are isothermally treated for 2 h at 1373 K and 1473 K promoting coalescence of Ti5 Si3 particles. As much was possible for the kinetics conditions provided to the heat-treated alloys, there have been produced dispersion of the Ti5 Si3 particles in the Ti matrix, but this target was very much hindered by the inertial behavior of the proeutectic Ti5 Si3 particles of the Ti–23Si alloy during heat-treatment step.
The authors wish to express thanks to the financial support of the National Natural Science Foundation of China (50761003, 50831007), the Key Project of China Ministry of Education (207085) and the opening project of the State Key Laboratory of Powder Metallurgy. References
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Evolution of the microstructure and hardness of the Ti–Si alloys during high temperature heat-treatment Yongzhong Zhan a,∗ , Xinjiang Zhang a , Jing Hu a , Qinghua Guo a , Yong Du b a b
Laboratory of Nonferrous Metal Materials and New Processing Technology, Ministry of Education, Guangxi University, Nanning, Guangxi 530004, PR China State Key Laboratory of Powder Metallurgy, Central South University, Changsha, Hunan 410083, PR China
a r t i c l e
i n f o
Article history: Received 9 December 2008 Received in revised form 5 January 2009 Accepted 7 January 2009 Available online 20 January 2009 Keywords: Metals and alloys Microstructure X-ray diffraction Scanning electron microscopy Optical microscopy
a b s t r a c t Samples of Ti–8Si, Ti–13.67Si and Ti–23S (at.%) alloys were prepared by vacuum ar melting technique and heat-treated for 2 h at 1373 K and 1473 K in a vacuum environment. The as-cast and annealed samples were examined by optical microscope (OM), X-ray diffraction (XRD), scanning electron microscope (SEM), energy dispersive X-ray analysis (EDX) and hardness test. During high temperature heat-treatment, the quasi-continuous network-shaped Ti5 Si3 + Ti phases in the as-cast Ti–8Si alloy finally converted to near-equiaxed Ti5 Si3 grains dispersed in the Ti matrix. For Ti–13.67Si alloy, the Ti5 Si3 grains finally also transformed to the near-equiaxed ones compared with the lamellar grains in the as-cast alloy. Moreover, compared with the as-cast alloy, Vickers hardness of Ti–8Si and Ti–13.67Si alloys was increased by the annealing. However, the microstructure of the Ti–23Si alloy could not change significantly. So, adversely Vickers hardness decreased after high temperature heat-treatment. The microstructural evolution of Ti–Si alloys is related with the diffusion of Si atoms in the Ti/Ti5 Si3 interface during high temperature heattreatment producing as much as possible homogeneous distribution of Ti5 Si3 particles in a continuous ␣-Ti matrix. The task was not achievable for Ti–23Si alloy because of the stability of the primary Ti5 Si3 phase particle during heat-treatment. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Titanium and titanium alloys has been intensively studied due to a lot of practical desirable properties, i.e. high strength, specific modulus and good corrosion resistance, etc. [1–3]. However, some major inherent problems like severe metal mould reaction, low castability and high-energy consumption of traditional casting titanium alloys have greatly impeded the development of these materials. As a new titanium casting alloy, Ti–Si eutectic alloy has some excellent casting properties, such as low melting point, narrow crystallization range and good fluidity, which is similar to the extensively applied Fe–C eutectic alloy and Al–Si eutectic alloy. Moreover, it consists of a ductile phase (␣-Ti, which is the matrix) and a brittle phase (Ti5 Si3 , acts as the reinforcement) [4–10]. This eutectic alloy has been considered to be the thirdgeneration eutectic alloy. Heat-treatments are usually adopted to increase the ductility and strength of titanium alloys by modifying their solid-state phase phase change behavior and microstructure. It can improve the mechanical properties of bulk materials that are related to the microstructure, processing history and heattreatment procedures. Some works conducted to the knowledge
∗ Corresponding author. Tel.: +86 771 3272311; fax: +86 771 3233530. E-mail address: [email protected] (Y. Zhan). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.01.017
of microstructure evolution and deformation mechanism of Ti alloys during hot deformation under different processing and heattreatment procedures [11–14]. These previous studies have shown that the thermal treatment of the Ti–6Al–4V alloy carried out at temperature above the  transus temperature leads to a ‘lamellar’ microstructure colonized plate-like ␣ at a low cooling rate, basket-weave at an intermediate cooling rate, Widmanstätten at a high cooling rate, and martensite after water-quenching. When processing is carried out in the (␣ + ) phase field, a much finer (␣ + ) structure can be obtained, because the primary ␣ phase limits the growth of the  phase [15,16]. In addition to processing temperature, other hot working parameters, such as strain and strain rate, also affect the microstructure, e.g. the volume fractions of the ␣ and  phases, the phase size and the ‘lamellar’ dimensions of the ␣ phase. The aspect ratio of ␣ lamellar phase is also a very important factor that influences the mechanical properties of the Ti–6Al–4V alloy [17–19]. These studies on microstructure and property of some Ti alloys during different heattreatment procedures are of great importance in exploring Ti–Si alloys. According to the currently accepted Ti–Si phase diagram (Fig. 1) [20], the reported stable phases in Ti-rich region are: -Ti, ␣-Ti, Ti3 Si, Ti5 Si3 and liquid (L). The compound Ti3 Si is formed from both eutectoid and peritectoid transformations (-Ti ↔ Ti3 Si + ␣Ti; -Ti + Ti5 Si3 ↔ Ti3 Si). Additionally, an eutectic reaction of (Ti–Si)
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2. Experimental procedure The stating materials used in this work were pure Ti (99.99 wt.%) and Si (99.99 wt.%). The alloy samples (all weighting 25 g) were arc-melted in a watercooled copper cast with a non-consumable tungsten electrode under pure argon atmosphere. Each sample was melted three times and turned around after melting for better homogeneity. For examining the effect of heat-treatment temperature on microstructure and mechanical properties, the alloy samples was heat-treated at 1373 K and 1473 K for 2 h at each temperature, and then cooled down to the room temperature with the furnace. The annealing temperature was determined according to Ti–Si binary phase diagram. Table 1 shows the nominal compositions of the alloys and the process of high temperature treatment. The as-cast and annealed samples were characterized via scanning electron microscope (SEM), optical microscopy (OM) and X-ray diffraction (XRD). The SEM images were obtained in a Hitachi S-3400N scanning electron microscope equipped with energy dispersive X-ray analysis (EDX). For the XRD experiment, the samples were performed on a Rigaku D/Max 2500 V diffractometer with Cu-K␣ radiation and graphite monochromator operated at 40 kV, 250 mA (20◦ ≤ 2 ≤ 60◦ ). The Materials Data software Jade 5.0 [21] and Powder Diffraction File (PDF release 2002) were used for phase analysis. Average hardness of the alloys were measured using model HV-30 Vickers microhardness tester with a test load of 3 kg and a load-dwell time of 30 s. More than eight indentations were performed on each samples to obtain the mean value. Fig. 1. Ti-rich region of the Ti–Si phase diagram.
3. Results and discussions L ↔ Ti5 Si3 + -Ti occurs at a temperature of 1603 K and 8.5 wt.% Si in Ti-rich region of the Ti–Si phase diagram. In order to improve the combined properties of Ti–Si alloys and step up their practical applications, the present work did research work to examine the evolution of the microstructure and hardness during high temperature processing of Ti–Si alloys, and then discussed their evolution mechanism from a microstructural viewpoint.
3.1. As-cast microstructure Fig. 2 shows the optical micrographs of the as-cast Ti–Si alloys. XRD data (Fig. 3) and EDX analysis have indicated the presence of Ti5 Si3 phases in all alloy samples. The micrograph of Ti–8Si hypoeutectic alloy (Fig. 2a) shows that the Ti5 Si3 + Ti lamellar particles distribute the domain boundaries of the mas-
Table 1 Compositions of the Ti–Si alloys and process of high temperature treatment in the present work. Alloys
Alloy type
Composition (at.% Si)
Holding temperature (K)
Holding time (h)
Ti–8Si Ti–13.67Si Ti–23Si
Hypoeutectic Eutectic Hypereutectic
8.0 13.67 23.0
1373 1373 1373
2 2 2
1473 1473 1473
Fig. 2. As-cast optical micrographs of Ti–Si alloy samples: (a) Ti–8Si; (b) Ti–13.67Si; (c) Ti–23Si.
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Fig. 3. XRD patterns of the as-cast Ti–Si alloy samples.
sive Ti proeutectic phase forming continuous or quasi-continuous network-shaped microstructure. In the micrograph of Ti–13.67Si eutectic alloy (Fig. 2b), a typical Ti + Ti5 Si3 eutectic microstructure is observed. A higher magnification should reveal peritectoid Ti and Ti5 Si3 plus Ti3 Si particles, eventually some finer eutectoid lamella. Fig. 2c shows the microstructural feature of Ti–23Si hypereutectic alloy. Some massive proeutectic Ti5 Si3 phase domain and typical Ti + Ti5 Si3 eutectic structure are observed in this alloy. Based on above analytical result, it is clear that the phase distribution of the as-cast Ti–13.67Si eutectic alloy is more homogeneous than that of the hypoeutectic and hypereutectic alloys. With the increase of Si content, the content of Ti5 Si3 phase gradually rises as indicated from the intensity of Ti5 Si3 of the XRD pattern in Fig. 3, which is the same result provided by the microstructure analysis. These results confirm the eutectic reaction of L ↔ Ti5 Si3 + -Ti at approximately 13.67 at.% Si as predicted in Fig. 1. The high temperature phase -Ti is observed in Ti–8Si
and Ti–13.67Si alloys from the XRD data, suggesting that noticeable amounts of the -Ti phase have been retained in the as-cast alloys under the rapid cooling rate of the arc furnace. Moreover, Fig. 2(a–c) shows that Ti–13.67 Si alloy exhibits smaller and more homogeneous grain-sized than Ti–8Si and Ti–23Si alloys. According to the SEM micrographs of this eutectic alloy (Fig. 4), it is observed that eutectic cells consist of some clubbed and lamellar Ti5 Si3 grains. The dendritic microstructure is possibly produced by reason of constitutional supercooling resulted from the rapid cooling process. It is clear that the long clubbed and lamellar phases are oriented differently in adjacent eutectic cells, which are produced by the peritectoid reaction -Ti + Ti5 Si3 (or Ti3 Si at lower temperature). From Fig. 1, the silicide in the equilibrium microstructure of Ti–Si should be Ti3 Si. However, the silicide is analyzed to be Ti5 Si3 by the result of XRD in Fig. 3. This is due to the fact that the formation of Ti3 Si would be a very slow solid-state phase change process at eutectoid temperature. Therefore, its formation was quite hindered (but not at all because the eutectoid ␣-Ti is present in the XRD pattern) in the present experimental condition, and the amount of peritectoid Ti3 Si formed should be too short to be detected by XRD. 3.2. Microstructural evolution at high temperature The optical micrographs of the samples after annealing at different temperature show the microstructure evolution of Ti–Si alloys during heat-treatment process. Fig. 5 shows the XRD result of Ti–8Si heat-treated hypoeutectic samples. Compared with the as-cast alloy, the phase component of this alloy did not altered during different high temperature processing, namely Ti5 Si3 , ␣-Ti and -Ti. However, due to quick atomic diffusion at high temperature, some continuous network-shape Ti5 Si3 phase arrays at Ti phase domain boundaries of the as-cast Ti–8Si alloy changed by means of coalescence mechanism during annealing at 1373 K, as is shown
Fig. 4. SEM micrographs of the as-cast Ti–13.67Si alloy: (a) low magnification, showing the different orientations of adjacent colonies; (b) low magnification, showing the clubbed and lamellar Ti5 Si3 grains; (c) high magnification, showing the dendritic microstructure.
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Fig. 5. XRD patterns of Ti–8Si heat-treated sample.
Fig. 7. XRD patterns of the Ti–13.67Si heat-treated alloy.
in Fig. 6a. At the annealing temperature of 1473 K, the networkshape Ti5 Si3 phases of the as-cast alloy disappeared entirely by some initial re-solution of Ti5 Si3 into the -Ti matrix, followed by Ti5 Si3 grain-growth and re-precipitation of Ti5 Si3 from Ti matrix onto Ti5 Si3 grains. By this way, more homogeneous distribution of the phases in the final microstructure is provided. Spherodized or near-equiaxed Ti5Si3 phases (Fig. 6b) with Ti5 Si3 grain diameter in the range of 15–25 m after annealing at 1473 K resulted. For Ti–13.67Si eutectic alloy, XRD data of the annealed sample shows that the phase component of this alloy is Ti5 Si3 and ␣-Ti. Moreover, the high temperature phase -Ti that exists in the as-cast alloy disappears during high temperature processing, as is shown in Fig. 7. According to the optical micrographs of annealed sample (Fig. 8), the long clubbed and dendritic lamellar phase disappear after annealing at 1373 K and 1473 K for 2 h, compared with that of the as-cast microstructure. Moreover, it is obvious that after annealing the grains are much finer. From SEM images of the annealed sample, it can be seen that after isothermally treated at 1373 K for 2 h, most of the long clubbed and lamellar eutectic Ti5 Si3 grains have transformed to equiaxed shape as well as some short clubbed ones (Fig. 9a). However, when the isothermal holding temperature increases to 1473 K for 2 h, the Ti5 Si3 phases almost completely turned into symmetrically equiaxed or near-equiaxed grains (Fig. 9b) with diameter in the range of 5–30 m, confirming the coalescence mechanism, which provided a continuous matrix of ␣-Ti phase. According the XRD analysis of the annealed Ti–23Si sample (Fig. 10), the phase component is the same with that of the as-cast alloy, i.e. Ti5 Si3 and ␣-Ti. Compared with the as-cast microstructure, the abundant massive proeutectic Ti5 Si3 phases still exist in the Ti–23Si hypereutectic alloy after annealed at different high temperature, as is shown in Fig. 11. However, the Ti5 Si3 phase in the typical Ti + Ti5 Si3 eutectic structure of the as-cast alloy has been
spheroidized and coarsened after annealed at 1373 K for 2 h. This phase is much bulky after 1473 K heat-treatment. From the Ti–Si binary phase diagram, at 1373 K and 1473 K, it is clear that the abundant massive proeutectic Ti5 Si3 phases in the as-cast alloy may not in heat-treatment process. 3.3. Analysis of evolution mechanism The free energy change of a system taking place in a solid-state nucleation and growth of the new phase has been defined as [22]: F = −Vfv + s + fε − fd where Vfv is the chemistry free energy for producing V is the volume new phase, s is the surface energy, fε is the elastic strain energy consumed by new phase and fd is the free energy arisen from crystal defect. When the volume of Ti5 Si3 phases is phases is a fixed value, the lamellar Ti5 Si3 grain possesses higher surface than the spherical one and so spontaneously transforms into the later shape in order to reduce free energy of the system. The spheroidization transformation of the Ti5 Si3 grains requires atomic diffusion and, the diffusion of Si atoms inside Ti5 Si3 is accelerated at high temperature (1373 K and 1473 K). At the same time, a mass of defects such as cavity and dislocation contribute to the diffusion of Si atoms. On heating the as-cast alloys above 866 ◦ C, there will be resolution of Ti3 Si by means of the first peritectoid reaction; above 1170 ◦ C the second peritectoid reaction causes the re-solution of Ti5 Si3 . It is well known that the smaller the radius titanium silicide the bigger is its solubility into the titanium phase [22,23]. The lamellar Ti5 Si3 phases (Fig. 4) have large surface free energy and there is natural trend to coalescence in order to reduce surface free energy. This was the reason by which high temperature (1373 K or 1473 K for 2 h) heat-treatment of Ti–8Si and Ti–13.67Si alloys conducted to near-equiaxed Ti5 Si3 grains. In addition, it is found that
Fig. 6. Optical micrographs of Ti–8Si hypoeutectic alloy sample after high temperature processing at (a) 1373 K and (b) 1473 K.
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Fig. 8. Optical micrographs of Ti–8Si alloy sample after high temperature processing at (a) 1373 K and (b) 1473 K.
Fig. 9. SEM micrographs of Ti–13.67Si eutectic alloy sample during high temperature processing: (a) 1373 K and (b) 1473 K.
spherical new phase. Moreover, the diffusion coefficient is temperature dependent according to Arrhenius equation and so its value at 1473 K is bigger than that at 1373 K, so the growth of spherical new phase at 1473 K is accelerated. That is the reason for the formation of much more bulky grains in the samples that annealed at 1473 K. However, there is contribution from the peritectoid reactions, i.e. at 1373 K, -Ti (p) + Ti3 Si = -Ti (r), and at 1473 K, -Ti (p) + Ti5Si = -Ti (r), where p: poor in Si and r: rich in Si content. 3.4. Hardness change during heat-treatment
Fig. 10. XRD patterns of the Ti–23Si heat-treated alloy.
the thermal treatment of Ti–Si alloys at 1373 K was not enough to provide structure separation and spheroidization of Ti5 Si3 grains. Apparently, the incomplete structure separation and spheroidization would possibly affect the final coarsening. From the microstructures of Ti–8Si and Ti–13.67Si alloys shown in Figs. 6 and 8, the grains of samples annealed at 1473 K are coarser than that of the annealed sample at 1373 K. At 1373 K or 1473 K, the diffusion coefficient (D) is the main factor for the growth of
In order to evaluate the effect of high temperature processing on the mechanical properties of Ti–Si alloys, Vickers hardness measurements were carried out on all samples, as is presented in Table 2. More than eight indentations were performed on each sample in its three different states. The indent sizes are in the range of 100–125 m, larger than the grain-sizes in the microstructures analyzed (as-cast and heat-treated) to obtain the mean value. Compared with the as-cast alloys, the Vickers hardness of Ti–Si alloys had been obviously altered after high temperature heat-treatment. For the Ti–8Si alloy, hardness increments of 141 HV and 130 HV were observed by annealing at 1373 K and 1473 K, respectively. Moreover, the hardness value of Ti–13.67Si alloy was increased from 379 HV for the as-cast state to 410 HV and 435 HV after annealing at 1373 K
Table 2 Average Vickers hardness of as-cast and high temperature treated Ti–Si alloy samples. Alloys Hardness (HV)
Ti–8Si
Ti–13.67Si
Ti–23Si
As-cast
1373 K
1473 K
As-cast
1373 K
1473 K
As-cast
1373 K
1473 K
376
517
506
379
410
435
458
387
365
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Fig. 11. Optical micrographs of Ti–23Si alloy sample during high temperature processing: (a) 1373 K and (b) 1473 K.
and 1473 K respectively. However, the hardness of Ti–23Si alloy was decreased by high temperature heat-treatment. The as-cast hardness of Ti–Si alloy was altered due to the variation of grain-size and constituent particles. In addition, the near-equiaxed Ti5 Si3 particles that distribute homogeneously in the continuous Ti matrix are propitious to the improvement of both hardness and strength. The above mechanisms may contribute to the increase of Vickers hardness of the hypoeutectic and eutectic Ti–Si alloys after high temperature heat-treatment and long diffusion time.
Acknowledgements
4. Conclusion
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The microstructural evolution of Ti–Si alloys that took place during high temperature heat-treatment of the as-cast alloys is described in the following. The quasi-continuous network-shape Ti5 Si3 + Ti intergranular phases of the as-cast Ti–8Si alloy transformed to near-equiaxed grains of Ti5 Si3 by annealing at high temperature, with grain-sizes in the range from 15 m to 25 m. In turn, the lamellar microstructure of Ti–13.67Si as-cast alloy transformed to Ti5 Si3 grains (5–30 m in size) near-equiaxed and homogeneously dispersed in the ␣-Ti matrix. Moreover, the Vickers hardness of Ti–8Si and Ti–13.67Si alloys was increased by high temperature heat-treatment. Unlikely, the microstructures of Ti–23Si alloy had not been altered significantly, and the hardness was decreased during high temperature heat-treatment. The microstructural evolution of Ti–Si alloys is related to the diffusion of Si atoms in the Ti/Ti5 Si3 particles interface when the alloys are isothermally treated for 2 h at 1373 K and 1473 K promoting coalescence of Ti5 Si3 particles. As much was possible for the kinetics conditions provided to the heat-treated alloys, there have been produced dispersion of the Ti5 Si3 particles in the Ti matrix, but this target was very much hindered by the inertial behavior of the proeutectic Ti5 Si3 particles of the Ti–23Si alloy during heat-treatment step.
The authors wish to express thanks to the financial support of the National Natural Science Foundation of China (50761003, 50831007), the Key Project of China Ministry of Education (207085) and the opening project of the State Key Laboratory of Powder Metallurgy. References