Microstructures and mechanical properties of ternary Ti–Si–Sn alloys

Microstructures and mechanical properties of ternary Ti–Si–Sn alloys

Journal Pre-proof Microstructures and mechanical properties of ternary Ti–Si–Sn alloys Chandra Sekhar Tiwary, Manas Paliwal, Sanjay Kashyap, Praful Pa...

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Journal Pre-proof Microstructures and mechanical properties of ternary Ti–Si–Sn alloys Chandra Sekhar Tiwary, Manas Paliwal, Sanjay Kashyap, Praful Pandey, Suman Sarkar, Ipshita Kundu, Shakti Bhaskar, In-Ho Jung, K. Chattopadhyay, Dipankar Banerjee PII:





MSA 138472

To appear in:

Materials Science & Engineering A

Received Date: 16 July 2019 Revised Date:

26 September 2019

Accepted Date: 27 September 2019

Please cite this article as: C.S. Tiwary, M. Paliwal, S. Kashyap, P. Pandey, S. Sarkar, I. Kundu, S. Bhaskar, I.-H. Jung, K. Chattopadhyay, D. Banerjee, Microstructures and mechanical properties of ternary Ti–Si–Sn alloys, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/ j.msea.2019.138472. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Microstructures and mechanical properties of ternary Ti-Si-Sn alloys Chandra Sekhar Tiwary1,2*, Manas Paliwal3, Sanjay Kashyap1, Praful Pandey1, Suman









Chattopadhyay1, Dipankar Banerjee1 1

Materials Engineering, Indian Institute of Science, Bangalore, India


Metallurgical and Materials Engineering, Indian Institute of technology,

Kharagpur, West Bengal, India 3









Gandhinagar, India 4

Department of Materials Science and Engineering, Seoul National University,

South Korea


Abstract Titanium based alloys are one of the important structural materials with applications ranging from aerospace to biomedical industries. Several ternary TiSi-Sn alloys were explored in this study to design the microstructure comprising of binary and ternary eutectics corresponding to Ti-Si, Ti-Sn and Ti-Si-Sn system. The microstructural evolution in these alloys was studied using a combination of characterization techniques and thermodynamic calculations. The predicted solidification path from thermodynamic calculations well supported the experimental microstructural data. In addition, mechanical property and microstructure







measurements and evaluating stress-strain curves under compression. Keywords: Ti-Alloys; Eutectic microstructure; Thermodynamic calculations; Solidification. *Corresponding authors: [email protected], [email protected], [email protected], [email protected]


1. INTRODUCTION The mechanical and chemical properties of Titanium alloy with low density makes it as one of the most attractive structural materials [1,2]. Ti alloys are widely used as aerospace materials [1-4]. In pursuit of development of high temperature-high strength Ti based alloys, researchers have alloyed it with several class of different elements [5-15]. The isomorphous bcc elements such as Nb, Mo, Zr and V can improve the mechanical properties through solid solution hardening [5,15]. On the other-hand, alloying elements such as Ni, Sn, Ga and Al improve the corrosion/oxidation resistance along with mechanical properties [9,15]. Solute Au is known to improve mechanical, corrosion and wear resistance of Ti alloy [16]. Ti alloys with eutectic or composite structure can significantly improve the high temperature mechanical properties and oxidation resistance. Although, carbide and boride reinforced Ti composite attracted attention due to their high temperature stability [17], the interface engineering of matrix and their reinforcement is rather challenging. The in-situ grown metal-intermetallic eutectic structure provides better interface and good control over morphology, hence mechanical properties can be improved [18-21]. In last few years, Al, Ni, Sn and Nb based Ti eutectic alloys exhibited remarkable improvement in strength and oxidation resistance [18-29]. Among Ti based eutectic alloys, Ti-Si eutectic alloys are the most promising due to their low density and high strength at elevated temperature [30,31]. However, room temperature ductility and environmental resistance of these alloys remain a challenge. In recent years, several ternary alloying additions have been used to improve these limitations. The chemistry of the eutectic alloys has been widely varied from low amount of alloying addition to formation of ductile ternary phases. 3

The third alloying element gives addition degree of freedom to fine tune the microstructure and hence mechanical properties [32-40]. Recently, it is demonstrated that Sn addition in Ti-Fe based eutectic alloy can result in improvement of its mechanical properties [41-45]. In the current work, we have explored the addition of Sn to optimize microstructure and mechanical properties of Ti-Si eutectic alloy. That is, the main focus of the present study is Ti-Si-Sn system. Ti rich region of Ti-Si and Ti-Sn systems is explored recently [46,47]. The Ti-Si phase diagram [34] shows an intermediate phase Ti5Si3 which participates in the eutectic reaction, L

β-Ti (beta-Ti) +Ti5Si3. The Ti3Si phase forms through solid

state reaction between β-Ti (beta-Ti) and Ti5Si3. In the Ti-Sn phase diagram [33], an intermediate solution phase Ti3Sn also forms through the eutectic solidification, and subsequent peritectoid reaction, β-Ti (beta-Ti) + Ti3Sn

α-Ti (alpha-Ti), is

also observed in Ti rich side. Previously, isothermal sections of ternary Ti-Si-Sn were investigated at 200, 900 and 1200 oC using a combination of thermal analysis, X-ray diffraction and SEM-EDS [48-49]. A ternary solution phase (Tau phase) with average composition of 60.25–61.03 Ti, 15.01–21.77 Si (at. %) and balance Sn was reported to be stable till 1200 oC. In addition no Si solubility was reported in Ti-Sn binary phases and no Sn solubility was found in Ti-Si binary phases with the exception of Ti5Si3 which showed a limited Sn solubility [49]. However, the liquidus of Ti-Si-Sn system has not been determined so far, and therefore the solidification path and as cast microstructural change in the ternary system are not investigated yet. The Ti-Si and Ti-Sn systems contain eutectics between β-Ti (beta-Ti) and the intermetallics Ti5Si3 and Ti3Sn, respectively. Our objective is to evaluate the 4

solidification behavior of ternary Ti-Si-Sn alloys and explore the possibility of designing as cast microstructure containing binary intermetallic and solution phases. The results of this work can act as a road map to develop new Ti alloys with enhanced mechanical properties.

2. EXPERIMENTAL DETAILS The alloy composition studied in the present study are shown in Fig.1. The compositions were designed to evaluate interactions between the Ti solid solution (beta-Ti) and Ti3Sn and Ti5Si3 intermetallics. Various compositions were prepared using a vacuum arc remelting furnace. The gases species such as nitrogen, oxygen and hydrogen can be dissolved in liquid metal during melting in open furnaces, and they can be detrimental to the mechanical properties of majority of alloys. Therefore, it is ideal to melt the alloys under a vacuum or argon atmosphere. About 20 grams of the mixture of starting alloying elements were placed in the Cu crucible equipped with water cooling jacket. 200mA of DC current was applied to start an arc between the metallic mixture and a non-consumable tungsten electrode, and the melting was continued till the entire charge was molten. In the present study, each alloy was remelted five times in order to achieve homogeneity in the alloy. The microstructures of the as solidified samples were imaged with the help of Scanning Electron Microscope (SEM) under secondary electron (SE) and back scatter electron mode to reveal the topographic and compositional differences between the phases, respectively. The phase analysis was conducted using Electron Probe Micro Analysis (EPMA) equipped with wavelength-dispersive spectroscopy (WDS). JEOL JXA-8530F EPMA is equipped with 6 WDS detectors (3 for heavy 5

elements and 3 for light elements) and an energy dispersive X-ray spectrometer (EDS). This combination can simultaneously analyze 3 elements WDS + 16 elements EDS and collect image signals from backscatter and secondary electron detectors as well. An accelerating voltage of 15 kV and beam current of 10nA with counting time of 30 s on peaks and 10 s on the background was used for the Ti, Si and Sn analysis. The raw data were finally converted into weight fraction data by employing ZAF correction using pure Ti, Si and Sn. The structures of the phases in the samples were identified by X-Pert PRO X-ray diffraction (XRD) equipment. The X-Pert PRO has a fixed tube Cu target, secondary graphite monochromator and a flat plate sample holder. It also has X-Celerator that reduces data acquisition time by approximately 1/4th. The mechanical properties of alloys were determined using micro-indentation (CSM Instruments SA). The uniaxial compression test (Instron UTM) was performed to determine the yield strength and fracture strain. SEM analysis was used to examine the fracture surface. The thermodynamic and solidification calculations for the Ti-Si-Sn system were performed with FactSage 7.2 software [50]. Thermodynamic database of the Ti-SiSn system was prepared by combining the optimized thermodynamic parameters (solution phases and compounds) of the three constituent binary systems: Ti-Si, TiSn and Si-Sn. Ternary Tau phase was also incorporated in the system by fixing its Gibbs energy to reproduce the experimental isothermal phase diagrams of the Ti-SiSn system at 200, 900 and 1200 oC [48,49]. For the thermodynamic optimization, Modified Quasichemical Model [51-52], and Compound Energy Formulism [53] were employed to describe the Gibbs energy of liquid and solid solution phases, respectively. For proper interpolation of the Gibbs energy of ternary liquid solution, 6

Toop interpolation technique [54] with Ti as asymmetric component in the Ti-Si-Sn liquid solution was adopted; As both binary liquid Ti-Si and Ti-Sn solutions have strong negative enthalpy of mixing while binary Si-Sn liquid shows positive enthalpy of mixing, Toop interpolation technique is the proper choice for the description of the Gibbs energy of the ternary liquid solution. In order to confirm the accuracy of the thermodynamic database, DSC experiments were also performed. Two ternary parameters were used for the liquid phase to explain the solidified microstructure in Ti-Si-Sn alloys. The details of thermodynamic optimizations will be published elsewhere [55]. 3. RESULTS and DISCUSSION The liquidus projection for the Ti-Si-Sn system is calculated in Fig. 1 using the new thermodynamic database. The primary crystallization regions along with the isotherms are indicated. The beta-Ti phase represents Ti rich solid solution with mutual solubility of Sn and Si. Ti3Sn and Ti5Si3 phases are the binary phases which are of main interests in this work. A small primary crystallization phase region of ternary Tau phase is also calculated. In Ti region, liquid miscibility gap is calculated which originates from the metastable liquid miscibility gap in the Si-Sn system. The microstructure of Ti-15 at. % Si binary alloy is shown in Fig. 2. The SEM images in BSE mode at two different magnifications are shown in Fig. 2(a) and (b). A fine rod shaped eutectic structure with sub-micron meter spacing can be observed in high magnification images. Along with fine eutectics, the alloy consists of small fraction of primary Ti phase. The phase identification is performed using TEM as 7

shown in Fig. 2(c)-(g). The bright field TEM image with diffraction pattern (see in inset) shown in Fig. 2(c) reveals sub-micron eutectic spacing and confirms the eutectic structure to consist of beta-Ti and Ti5Si3 phase. The experimental as cast microstructures of ternary alloys and Scheil cooling calculation results are presented in Figs. 3 to 6. The as solidified microstructure of alloy A (Fig.3(a)) consists of plate like colonies of both coarse and fine regions together with globules of a phase in light contrast. The compositions of various regions in the micrograph were determined by EPMA and summarized in Table 2. The globular phase in grey contrast is beta-Ti solid solution with composition 87.9 % Ti, 10.9 % Sn, 1.2 % Si (at. %). The darker phase associated with the coarser regions of the eutectic structure corresponds to Ti5Si3 phase with composition 67.4 % Ti, 2.6 % Sn, 30.0 % Si (at. %). The lighter plate underlying this structure has composition of 77.5% Ti, 12.7 % Sn, 9.8 % Si (at.%). In order to understand the development of the microstructural features, Scheil cooling calculations were performed for the alloys and the results are shown in Fig. 3. The calculations reveal that beta-Ti and Ti3Sn phase co-precipitate from liquid at 1458.26 oC. Following this, Ti5Si3 phase forms at 1436.26 o C, owing to a ternary peritectic reaction. This is followed by co-precipitation of beta-Ti and Ti5Si3 phases till the final eutectic temperature on the Ti-Si binary is reached. The calculations support the experimental observations for alloy A as the globular phase in the SEM image that corresponds to the primary beta-Ti phase. The lighter region in the SEM image corresponds to the region containing binary eutectic of beta-Ti and the intermetallic Ti3Sn phase while the coarser and darker eutectic region corresponds to beta-Ti and Ti5Si3 phase. 8

Alloys B and C exhibit similar microstructure. Beta-Ti solid solution (see the composition in Table 2) is formed as primary crystalline phase in both alloys. As shown in Fig. 4 (a) and (b), the amount of this primary phase (light contrast) is significantly larger than the other alloys. Eutectic structure is also formed between the primary phases, and their overall compositions are presented in Table 2. The composition analysis of each phase in the eutectic region was also performed. Based on this result, the phase in dark contrast in the eutectic region is determined to be Ti5Si3 phase. The other eutectic component having the similar contrast as that of the primary nodules is most likely beta-Ti phase. The volume fraction of primary crystalline beta-Ti phase in alloy C is higher than that of alloy B. The results of Scheil cooling calculations for alloys B and C are well matched with the experimental microstructure. In general, the ternary alloy typically can have ternary eutectic structure with three kinds of crystalline phases. However, alloys B and C show only beta-Ti and Ti5Si3 phase in the eutectic structure. This happens because both beta-Ti and Ti5Si3 phases form solid solutions. As can be seen in Fig. 6(a), the liquidus temperature along the univariant line between beta-Ti and Ti5Si3 phases decreases toward binary Ti-Si side. That is, the composition of liquid phase (solidification path) in alloys B and C changes from ternary Ti-Si-Sn to binary Ti-Si with the progress of the solidification. Alloys D and E have similar microstructural features, as can be seen in Fig. 4 (c) and (d). Alloy D contains small amount of residual eutectic phase. The composition of the residual eutectic phase is given in Table 2. The composition of the phase in grey contrast corresponds to that of the beta-Ti while the composition of the rod shaped phase in light contrast corresponds to that of the Ti3Sn phase. The 9

morphology of Ti3Sn suggests that it has precipitated in Widmanstatten form in both the alloys during solidification. The Scheil cooling calculations for alloys D and E show the same solidification path. In the solidification of alloy D, for example, the primary beta-Ti phase is calculated to form at 1574.08 oC, followed by co-precipitation of beta-Ti and Ti3Sn phases starting at 1549.08 oC. Similar to alloys A-C, Ti5Si3 phase is formed due to the ternary peritectic reaction in alloys D and E, followed by co-precipitation of beta-Ti and Ti5Si3 phase. According to the calculation, the volume fraction of residual eutectic phase (beta-Ti + Ti5Si3 ) is very less in alloy E compared to alloy D, which is

consistent with the as cast

microstructures. This difference in the volume of residual eutectic phase is simply due to longer binary eutectic solidification path of alloy E than that of alloy D, as can be seen in Fig. 6 (a). Alloy G has distinct microstructure appearance (Fig. 5(a)) from the previous compositions. The light contrast phase has a composition corresponding to that of the Ti3Sn phase. The 6-fold symmetry of the dendritic structure is consistent with the hexagonal structure of this phase. The regions between the dendrites of Ti3Sn phase have two distinct appearances. The majority of the area is occupied by a eutectic structure whose composition is similar to that of the eutectic regions of alloys A-E. This area appears as a dark lamellae coupled with a phase in grey contrast (corresponding to beta-Ti), as seen in the magnified inset of the figure. A eutectic type structure of Ti3Sn and Ti5Si3 is also observed in other regions where Ti5Si3 appears as dark grey faceted phase. The complex microstructural evolution in this alloy could be understood with the aid of Scheil cooling calculation. As seen in the Scheil calculations (Fig. 5), Ti3Sn is the primary phase to precipitate from 10

liquid. This is in agreement with the Ti3Sn dendrites observed in the as-cast microstructure. The solidification proceeds with the co-precipitation of Ti5Si3 and Ti3Sn phase which is observed as eutectic type structure in as-cast microstructure. After the ternary peritectic reaction, beta-Ti and Ti5Si3 phase co-precipitate during solidification, which appears as a dark lamellae combined with a phase in grey contrast in the as-cast microstructure. The microstructures of alloys (H, I and J) contain Ti5Si3 phase as the primary crystalline phase. This phase has a distinctly faceted appearance. In alloy H (Fig 5(b)), the Ti5Si3 phase in dark contrast is enveloped with Ti3Sn phase in light contrast while the residual region is in grey contrast with a composition corresponding to the beta-Ti phase. However, in alloys I and J (Fig. 5(c) and (d)) the primary phase is enveloped by a phase in grey contrast corresponding to the composition that is rich in both Si and Sn, and quite different in composition from the beta-Ti phase. This phase is confirmed to be the ternary Tau phase, which was reported in previous studies for sub-solidus isothermal sections [48,49]. The volume fraction of the primary Ti5Si3 phase in alloy J is much less compared to that in alloy I, due to lower Si content in alloy J. The Scheil cooling calculations for alloys H to J are shown in Fig. 5(b) to (d). The calculations indicate that Ti5Si3 is the primary precipitation phase in these alloys which clearly appear as dark grey faceted phase in the microstructures. In these alloys, the dark grey Ti5Si3 is enveloped by a light grey contrast Ti3Sn phase according to EPMA-WDS analysis and the cooling calculations. In case of alloy H, the Scheil cooling calculations indicate the formation of beta-Ti phase at the end of solidification which is also experimentally observed in the eutectic areas as a phase 11

with grey contrast. The calculations for alloys I and J suggest the formation of ternary Tau phase due to ternary peritectic reaction which is also supported by experimental evidence. In both these alloys the solidification ends with the formation of Ti2Sn phase. The Scheil cooling paths for alloys A to J are superimposed to the liquidus projection of Ti-Si-Sn system in Fig. 6. As seen in Fig. 6 (a) the initial alloy composition of alloy A lies on the univariant line between beta-Ti and Ti3Sn phase. Alloys B and C lie in the beta-Ti primary phase region and precipitate relatively large amount of beta-Ti phase during solidification. As seen in Fig. 6 (a) the alloys A-E undergo a ternary peritectic reaction during solidification. In these alloys, solidification ends at the binary Ti5Si3 composition where the last remaining liquid transforms to the binary eutectic microstructure of beta-Ti and Ti5Si3 phase. As indicated in Fig. 6 (b) alloy G composition lies in the Ti3Sn primary phase region whereas alloys H to J are located in the primary crystallization region of Ti5Si3 phase. As evident from Fig. 6 (b) the solidification for alloy G proceeds with the co-precipitation of Ti3Sn and Ti5Si3 phases followed by a ternary peritecitic reaction where beta-Ti, is formed. Alloy H follows a similar path with the solidification terminating at the Ti5Si3 composition in the Ti-Si binary. However, in case of alloy H the volume fraction of primary Ti5Si3 phase is more compared to alloy G. Alloys I and J solidify with Ti5Si3 as the primary phase followed by co-precipitation of Ti5Si3 and Ti3Sn phases and the solidification proceeds with the formation of ternary Tau phase due to a ternary peritectic reaction. As discussed above, all these solidification paths are well consistent with present experimental results. The hardness of the present solidified alloys was measured using microhardness 12

indentation technique. The bar diagram of hardness value and the BSE SEM image of the indent for different alloys are shown in Fig.7 (a)-(j). The binary Ti-Si alloy shows lowest hardness of 450±15Hv. The ternary alloys show higher hardness compared to binary alloy, whereas the alloy A exhibits highest hardness of more than 700±22Hv. The hardness of the other alloys lay in the range of 570Hv to 650Hv. The SEM image of the indents are shown in Fig. 7 (b)-(j). The morphology of indents is correlated with deformation behaviour. In order to elucidate the relationship between mechanical property and microstructure, binary Ti-Si and ternary A, D and G alloys were chosen for the case studies. The stress-strain curves for the alloys deformed in compression mode are presented in Fig 8 (a). The stress-strain curve for binary alloy is similar to that observed for a brittle material, i.e. absence of sharp yielding point. The stress gradually increases with strain (linear portion of the stress-strain curve) till it reaches a maximum value and then drops down suddenly. It was observed that the samples failed layer by layer as shown in the fracture surface in Fig. 8(b) and (c). Yielding was observed in alloys A and D. The yield stress of 965MPa and 480MPa were revealed for alloys A and D, respectively. Interestingly, alloy G shows brittle fracture with a sharp fracture stress of 550MPa. The SEM images of the subsurface after compression of alloys A and D, revealing the deformation band and the crack initiation are as shown in Fig 8 (d)-(i). In alloy A (Fig.8 (d)) dimple like fracture sub-surface is seen, suggesting ductility in the alloy, which is consistent with the stress strain curve. The high magnification SEM image in BSE mode (Fig. 8(e)) shows cracks initiates in the Ti5Si3 phase. The deformation of alloy D is shown in Fig. 8(f)-(g). It is seen that the deformation 13

bands are mostly present in the eutectic region. The cracks propagate straight through this phase and clearly avoid the tin (Sn) rich phase. However, the cracks are present only in the Ti3Sn phase which is responsible for crack initiation. In alloy G (Fig. 8(h)-8(i)) subsurface after fracture indicates brittle failure. The dark phases correspond to Ti5Si3 which acts as crack initiation phase. The dendritic structure (Ti5Si3) appears to fail in a brittle manner. The crack passes through these dendrites. The surrounding eutectic regions indicate some amount of crack resistance in the alloy. In alloy G, there are no slip lines visible. This means that the alloys have failed before entering into the plastic deformation domain. In alloy G, Ti5Si3 phase appears to be the most brittle, which is consistent with stress-strain curve. Among all these alloys, alloy A has the maximum strength and plasticity. In alloy A large amount of deformation bands are seen both in Ti3Sn and beta-Ti which means that the alloy has undergone considerable plastic deformation before failure and this is in accordance with the stress-strain plot obtained from compression test. Although the slip lines are present in both the phases the cracks appear only in the Si rich phase indicating that Ti5Si3 phase fails first under compression. The eutectic region appears unaffected. In all the cases, Ti5Si3 phase favors crack initiation and propagation. The SEM image of indent (shown in Fig. 7(b)) shows cracks in the Ti5Si3 phase which further confirms the brittle behaviour. In alloy D the Si content is small and correspondingly the amount of eutectic is also low. Therefore, out of the two main constituents (Ti3Sn and beta-Ti), Ti3Sn is the crack initiation site because it is more brittle compared to beta-Ti.


4. CONCLUSION Several ternary alloys were selected in the ternary Ti-Si-Sn system to obtain a solidified microstructure comprising of eutectics originating from the binary Ti-Si and Ti-Sn system. Microstructural characterization of the solidified alloys was performed using SEM, EPMA, XRD and TEM. In addition, thermodynamic and solidification calculations were performed to assist the analysis of experimental microstructural data. Moreover, mechanical properties were determined for these alloys using hardness and stress-strain curve measurements. The solidification microstructures of these alloys changed drastically with the formation of different primary and secondary phases depending on the starting composition. Depending on the alloy composition, beta-Ti, Ti5Si3 and Ti3Sn were either the primary solidification phases or the part of the secondary microstructure that form during the later stages in the solidification. Scheil cooling calculations reasonably explained the evolution of solidified microstructures in all the ternary alloys. Ternary alloys reveal drastic improvements in hardness value compared to binary Ti-Si alloy. The compression test shows an improvement of plasticity in eutectic structure due to the microstructure modification by addition of Sn to Ti-Si alloy. The alloy A exhibits best strength and plasticity under compression. The present study of experiments and thermodynamic calculations can help in designing high strength Ti based alloys.


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Fig. 1: Liquidus projection for the ternary Ti-Si-Sn system showing the initial alloy compositions used in the present study.


Fig.2: (a)-(b) SEM image of binary alloy at two different magnification (10micron for a, 2micron for b scale bar). (c) Bright field TEM image of the eutectic, inset shows diffraction pattern of the two lamellae. (d) Dark field TEM image of the primary phase with diffraction pattern. (e) HAADF TEM image of the eutectic. Inset shows composition profile. (f)-(g) EDS composition map of Si and Ti.


Fig.3: (a)-(b) BSE-SEM images at two different magnifications of alloy A (the scale bars in image a and b are 10 and 1 µm, respectively), (c) Scheil cooling calculation of alloy A, (d) Dark field TEM image of the eutectic, (e)-(f) Diffraction pattern of the two lamellae, (g) HAADF image of the eutectic, and (h)-(i) EDS mapping and composition line profile of the eutectic.


Fig.4: (i)-(ii) SEM image in BSE mode at two different magnification (10 µm and 2 µm) for (a) Alloy B (b) Alloy C (c) Alloy D (d) Alloy E. (iii) Scheil cooling calculations for the respective alloys.


Fig.5: (i)-(ii) SEM image in BSE mode at two different magnification (10 µm and 2 µm) for (a) Alloy G (b) Alloy H (c) Alloy I (d) Alloy J. (iii) Scheil cooling calculations for the respective alloys.




(b) Fig.6: Scheil cooling solidification path of ternary Ti-Si-Sn alloy. (a) alloys A-E (b) G-J


Fig.7. (a) Bar diagram of hardness values of different alloys. (b)-(j) The SEM image showing the indent in BSE mode for the binary alloy, and ternary alloys A, B, D, E, G, H, I and J (scale bar 10 µm).


Ti3Sn Ti

Ti5Si3 Ti3Sn


Fig.8. (a) Engineering stress-strain plot of binary alloy and ternary alloys A, D and G. SEM image in BSE mode of subsurface of alloys at two different magnification (10 µm and 5 µm). (b) – (c) binary alloy, (d) - (e) alloy A, (f) – (g) alloy D, (h) – (i) alloy G


Table:1 Composition of the alloys ( mole % )

Alloy A B C D E G H I J

Ti 79.7 85.3 82.7 82.1 82.2 74.3 71.3 65.9 68

Si 10.8 10.3 7.9 3.7 0.5 10.1 16.4 19.4 12.5

Sn 9.5 4.4 9.4 14.2 17.3 15.6 12.3 14.7 19.5


Table. 2. Constituent phases identified in the solidified microstructures of ternary Ti-Si-Sn alloys, along with EPMA composition data. Alloy A









Fine, eutectic, grey contrast




Fine, eutectic, light contrast




Grey nodules, Ti solid solution




Coarse eutectic, light contrast




Coarse eutectic, dark contrast, Ti5Si3




Eutectic region




Primary phase, light contrast, Ti solid solution







Eutectic region




Primary phase, light contrast, Ti solid solution




Dark contrast around primary phase, Ti5Si3




Eutectic region







Matrix, grey contrast, Ti solid solution





Matrix, grey contrast, Ti solid solution




Plates, light contrast, Ti3Sn





Primary phase, light contrast, Ti3Sn




Dark contrast, Ti5Si3




Grey contrast, Ti Solid solution



















Dark contrast, Ti5Si3




Light contrast, Ti3Sn




Grey contrast, Tau phase




Dark contrast, Ti5Si3






Dark contrast around primary phase, Ti5Si3

Plates, light contrast, Ti3Sn

Area with eutectic appearance, Ti3Sn+Ti5Si3 Eutectic Light contrast, Ti5Si3 Primary phase, dark contrast, Ti5Si3 Grey contrast, Ti solid solution




Light contrast, Ti3Sn




Grey contrast, Tau phase