Mechanics of Materials
Effect of TaC Powder Addition on the Microstructure and Mechanical Properties of Ti–Nb–Mn Alloy via Vacuum Sintering Process
2023 Volume 64 Issue 12 Pages 2714-2721
Details
2023 Volume 64 Issue 12 Pages 2714-2721
This study involved the fabrication of Ti–10Nb–3Mn–XTaC composites through the vacuum sintering process of powder metallurgy, where different proportions of TaC powders (X = 1, 2, and 3 mass%) were added. The composites were sintered under vacuum conditions at temperatures of 1225, 1250, 1275, and 1300°C for 1 hour each. The experimental findings indicate that the addition of 2 mass% TaC powders to the Ti–10Nb–3Mn alloys resulted in the optimal mechanical properties when sintered at 1275°C for 1 hour. The specimen exhibited a relative density of 96.24%, a hardness of 67.38 HRA, and achieved values of 1276.63 MPa for TRS (tensile rupture strength) and 40.37 GPa for the flexural modulus. The electron backscatter diffraction (EBSD) results revealed that the TaC facilitated the in-situ formation of TiC during the sintering process, which was uniformly dispersed in the Ti matrix. Additionally, the Ta atoms acted as β-stabilizing elements, forming a solid solution in the Ti matrix and improving both the microstructure and mechanical properties of the Ti alloys.
Over the past decades, titanium (Ti) and titanium alloys have been widely used in aerospace, petrochemical and medical implants due to their high strength, low elastic modulus, high corrosion resistance and biocompatibility.1,2) The manufacturing cost of Ti alloys is quite expensive, so how to reduce the cost of Ti alloys has become an important issue. In recent years, the vacuum sintering process of powder metallurgy (PM) technology has gradually been used to reduce the manufacturing cost of Ti alloys.3) Furthermore, the high-temperature β-phase of Ti alloys can be stabilized at room temperature by adding β-stabilizing elements such as niobium (Nb), tantalum (Ta), molybdenum (Mo) and manganese (Mn).4,5) D. Kalita and I. Çaha et al. discovered that Ti–Nb-based metastable β-Ti alloys are among the most promising candidates for long-term implantation because of their excellent corrosion resistance, biocompatibility and superior mechanical properties.6,7)
In addition, H. Dong et al. have pointed out that the addition of the alloying element manganese (Mn) can significantly improve the ductility of β-Ti alloys.5) Mn has been chosen as an alloying element by K. Cho and Y. Alshammari et al. due to its β-stabilizing effect and lower cost compared to other β-stabilizing elements.8,9) Consequently, Mn presents itself as a strong candidate for use as a low-cost β-stabilizing alloying element in binary alloys such as Ti–Mn and ternary alloys such as Ti–Nb–Mn. It is noteworthy that conventional high-strength Ti alloys contain β-stabilizing elements such as chromium (Cr), vanadium (V), niobium (Nb), molybdenum (Mo), iron (Fe), and manganese (Mn), which greatly enhance the fabrication ability and ductility of these alloys during hot and cold working operations.10,11) As a result, the Ti–Nb–Mn alloy has emerged as a novel β-Ti alloy with promising prospects.
Metal matrix composites (MMCs) offer advantages over alloys, including improved wear resistance, low thermal expansion coefficient, and good mechanical properties at elevated temperatures. These advantages are primarily attributed to the presence of reinforcements.12) Titanium matrix composites (TMCs) reinforced with ceramic particles have garnered considerable attention due to their high specific strength and stiffness, good chemical resistance, and promising mechanical properties.13) In addition, these composites typically use ceramic phases as reinforcing phases, such as TiC, TaC, NbC, WC, or carbon fibers, which are dispersed in a matrix phase and produce composites with exceptional mechanical strength and toughness.13) Among them, tantalum carbide (TaC) exhibits excellent physical and chemical properties, such as a high melting point, high hardness, and outstanding chemical stability. Therefore, TaC is suitable for enhancing the mechanical properties of composites and can be used as a ceramic reinforcement.14,15) Moreover, TaC acts as an effective grain growth inhibitor during the liquid phase sintering (LPS) of the powder metallurgy (PM) process, owing to its excellent solubility and mobility in the metal matrix phase.16,17)
Based on the literature review and research, it is evident that casting and subsequent thermomechanical treatments were commonly employed in the past for the preparation of titanium alloys. However, in recent years, these processes have been replaced by vacuum sintering using the powder metallurgy (PM) technique, which offers significant opportunities for developing and producing high-performance titanium alloys at a reduced cost. The PM process enables the production of complex-shaped components, minimizing material loss. Additionally, the addition of tantalum carbide (TaC) allows for the formation of titanium matrix composites (TMCs) that effectively enhance the mechanical properties of titanium alloys. In this study, a vacuum sintering process was utilized to fabricate a Ti–Nb–Mn alloy. The manganese content was adjusted to improve the mechanical properties of the Ti–10Nb alloy. Furthermore, TaC was added to the Ti–Nb–Mn alloy as a reinforcing carbide phase, based on the standard Gibbs free energy-temperature relationship of carbides, with the aim of enhancing the properties of the alloy.
Our previous research found that adding 3 mass% Mn to Ti–10Nb alloy and vacuum sintering at 1250°C for 1 hour resulted in the Ti–10Nb–3Mn alloy with optimal mechanical properties. In order to improve the mechanical properties of the Ti–Nb–Mn alloys, this study utilized Ti–10Nb–3Mn alloy powders and added different ratios of refined TaC powders (1, 2, and 3 mass%) as a strengthening element to explore the effects of a series of vacuum sintering processes using the PM technique. The mean particle sizes of the Ti, Nb, Mn, and TaC powders in this experiment were about 22.3 ± 1.5, 7.9 ± 0.5, 13.3 ± 0.5, and 1.9 ± 0.5 µm, respectively. Figure 1 illustrates the morphology of the Ti, Nb, and Mn powders, which exhibited an irregularly shaped surface. In contrast, TaC displayed an irregularly shaped and polygonal appearance. Furthermore, the different amounts of TaC powder (1, 2, and 3 mass%) were mixed and added to Ti–10Nb–3Mn alloy powders, hereafter, designated as 1 TaC, 2 TaC, and 3 TaC, respectively, and the mixing experiment utilized a three-degree space rotary mixer (Turbula Shaker-Mixer, T2C). After milling, a 1:1 mixture of PVA (polyvinyl alcohol) was added as a binder to the powders. The resulting powder mixture was then compacted into green specimens measuring 40 × 6 × 6 mm3 under uniaxial pressure at 200 MPa for 300 seconds. To avoid the effects of variations in the vacuum processing conditions, the degree of vacuum was maintained at 1.33 × 10−3 Pa. At the same time, the heating and cooling of the specimens were all completed inside the vacuum sintering furnace (with a cooling rate of 20°C·min−1 below the sintering temperature). Subsequently, the alloys were vacuum sintered at 1225, 1250, 1275, and 1300°C for 1 h, respectively.
The morphology of the powders: (a) Ti, (b) Nb, (c) Mn, and (d) TiC, respectively.
To evaluate the effects of microstructure evolution on the TaC powders added to Ti–10Nb–3Mn alloys, this study conducted volume shrinkage, apparent porosity, relative density, Rockwell hardness (HRA), TRS (tensile rupture strength) and flexural modulus tests; XRD (Bruker/D2 Phaser) analysis; optical microscopy (OM, Nikon optiphot 66); scanning electron microscopy (SEM, Hitachi-S4700); and electron backscatter diffraction (EBSD, JSM-7800F) to make microstructure observations. EBSD is a scanning electron microscopy (SEM) technique used to study the crystallographic structure of materials. An apparent porosity test was conducted following the ASTM C373-88 standard, and the volume shrinkage was measured in accordance with the ASTM C830 standard. Both tests utilized the Archimedes principle for calculations. Hardness tests were measured by HRA with a loading of 588 N, which followed the ASTM E18-3 standard. Furthermore, the Hung Ta universal material test machine (HT-9501A) with a maximum load of 25 tons was used for TRS tests (ASTM B528-05). The TRS was obtained by the equation Rbm = 3FLk/2bh2, where Rbm is the TRS, which is determined as the fracture stress in the surface zone, and F is maximum fracture load. In this work, L was 30 mm, k was chamfer correction factor (normally 1.00–1.02), b and h were 5 mm, respectively. The sample dimensions of the TRS test were 5 × 5 × 40 mm3 and tested at least three pieces to get the average value.
Flexural modulus was measured as the elastic modulus of the bending deformation of the material, while the elastic modulus was calculated from the TRS test results. This measurement followed the ASTM E2769-16M, and the elastic modulus was obtained by the equation E = PL3/4δt3w, where E was the flexural modulus, P was the maximum fracture load (N), L was span length (this study was 30 mm), δ was displacement (mm), t was thickness (mm), and w was width (mm), respectively.
In the research, corrosion potential (Potential Stat Chi 601) analysis uses three electrodes and follows ASTM G59-97: the reference electrode is a saturated of silver-silver chloride electrode, the auxiliary electrode uses a platinum electrode, and the working electrode is connected to the test specimens.1,13) The contact area of the specimen was 0.78 cm2. 1 N H2SO4 was used as the corrosive solvent, which was kept at room temperature. A scanning speed of 0.01 V·s−1, initial potential of −1.5 V, and the final potential of 3.5 V were controlled. The polarization curve was obtained by Corr-View software to analyze and compare the corrosion current (Icorr), corrosion potential (Ecorr) and polarization resistance (Rp) of sintered Ti–10Nb–3Mn–XTaC composite materials.
Figure 2 illustrates the volume shrinkage, apparent porosity, and relative density of Ti–10Nb–3Mn alloys with varying mass percentages of TaC under different sintering temperatures. It can be observed from Fig. 2(a) that the volume shrinkage gradually decreased as the TaC content increased (1 → 2 → 3 mass%). Conversely, an increase in sintering temperature (1225 → 1250 → 1275 → 1300°C) led to a significant increase in volume shrinkage. At the same sintering temperature, the volume shrinkage decreased with increasing TaC content. It is reasonable to suggest that this is due to the lower amount of TaC added and the higher ductility of pure manganese and pure niobium metals. The higher ductility enables a greater amount of plastic deformation during the compaction process, resulting in smaller variations in volumetric shrinkage during the subsequent vacuum sintering process. However, with an increase in sintering temperature, the volume shrinkage significantly increased due to the stronger sintering force that brought the powders closer together. It is evident that higher sintering temperatures effectively improved the volume shrinkage. Notably, the sample with 1% TaC sintered at 1300°C exhibited the highest volume shrinkage of 37.96%.
Comparison of the (a) volume shrinkage, and (b) apparent porosity and relative density of various mass% TaC added to Ti–10Nb–3Mn alloys by the different sintering temperatures.
Figure 2(b) displays that the apparent porosity of specimens with 1 TaC, 2 TaC, and 3 TaC after 1 hour of sintering at various temperatures. It can be observed that, under the same conditions (sintering temperature or TaC addition), there was an inverse relationship between the apparent porosity and volume shrinkage. Furthermore, all the specimens exhibited a significant decrease in apparent porosity as the sintering temperature increased. However, it is important to note that the apparent porosity increased with higher TaC content at the same sintering temperature. Among the specimens, the 1 TaC sample sintered at 1300°C had the lowest apparent porosity (0.21%).
In this experiment, as the sintering temperature increased, the diffusion of atoms between the powders was accelerated. This resulted in an increase in the neck size between the powders during the sintering process. As a consequence, the pores between the powders became smaller, leading to a decrease in apparent porosity and an increase in sintered density. Clearly, increasing the sintering temperature effectively reduced the apparent porosity, which is advantageous for the sintering densification of Ti–10Nb–3Mn–XTaC composites.
Previous research has indicated that increasing the sintering temperature can enhance thermal energy.1,2,13) This study further confirms that raising the sintering temperature effectively increases the driving force for Ti–10Nb–3Mn–XTaC composites, resulting in reduced porosity and increased sintered density. The experimental results demonstrate that the relative density decreases with an increase in TaC content, but increases with an increase in sintering temperature. As mentioned earlier, higher sintering temperatures effectively reduce the apparent porosity in this study. However, adding excessive TaC powders does not effectively improve the volume shrinkage, apparent porosity, or relative density. In other words, composite specimens with high TaC content require higher temperatures to achieve complete densification. Therefore, when the sintering temperature was increased to 1300°C, the relative density exhibited a significant increase. Consequently, the 1 TaC specimen sintered at 1300°C exhibited the lowest apparent porosity (0.21%) and the highest relative density (96.92%).
Figure 3 illustrates the results of hardness, TRS, and flexural modulus tests for Ti–10Nb–3Mn alloys with varying mass percentages of TaC under different sintering temperatures. The hardness values of the Ti–10Nb–3Mn–XTaC specimens demonstrated a noticeable improvement as the amount of added TaC increased. Furthermore, the hardness exhibited a significant increase with higher sintering temperatures, as shown in Fig. 3(a), with the highest hardness value observed in the 3 TaC specimen sintered at 1300°C (68.97 HRA). TaC carbide, known for its high hardness (>20 GPa), can effectively enhance the hardness of the composite material when incorporated into the alloys.
Comparison of the mechanical properties of various mass% TaC added to Ti–10Nb–3Mn alloys after sintering at different temperatures for 1 h: (a) hardness, (b) TRS and flexural modulus, respectively.
As previously mentioned, Ta can function as a β-stabilizing element in the titanium matrix, forming a homogeneous solid solution that contributes to strengthening effects. It is reasonable to speculate that the solid solution of Ta enhances the stability of the β-phase, resulting in harder β-phases. Additionally, the in-situ reaction (TaC + Ti → Ta + TiC) generates a significant amount of TiC (confirmed by subsequent EBSD analysis), which precipitates at the grain boundaries. This TiC precipitation effectively impedes dislocation movement, leading to dispersion strengthening effects. Consequently, the 3 TaC specimens exhibited the highest hardness (68.97 HRA) after 1 hour of sintering at 1300°C. Similarly, the 1 TaC and 2 TaC samples also demonstrated good hardness values (68.03 and 68.45 HRA, respectively). Although increasing the porosity and decreasing the density of sintered materials are capable of decreasing plastic deformation resistance and hardness, the results show that the TiC precipitation in the Ti–10Nb–3Mn–XTaC specimens has a more dominant effect on hardness.
Figure 3(b) also presents the TRS and flexural modulus values of the Ti–10Nb–3Mn–XTaC specimens at various sintering temperatures. It is noteworthy that the TRS value of the Ti–10Nb–3Mn–XTaC specimens exhibited an initial increase followed by a decline as the sintering temperature was raised. Importantly, an excessive addition of TaC (3 mass%) led to a decrease in TRS value. This observation suggests that with an increased TaC content, more β-phase can be formed (as confirmed by subsequent XRD analysis), and the β-phase possesses greater ductility compared to the α-phase. In other words, an appropriate increase in the proportion of β-phase contributes to enhanced alloy ductility, leading to improved toughness and strength. However, as the TaC content increased from 2 mass% to 3 mass%, more noticeable pores were formed (as shown in Fig. 2(b)), resulting in reduced strength.
Among all the different TaC additions and sintering temperatures, the specimen with 2 TaC sintered at 1275°C exhibited the highest TRS (1276.63 MPa). This result can be further compared with literature values for Ti–14Nb (949 MPa)4) and Ti–10Mn (1223 MPa)9) alloys after the sintering process. Remarkably, the 2 TaC specimens in this study actually demonstrated higher strength.
The flexural modulus values of Ti–10Nb–3Mn–XTaC specimens after various sintering temperatures are depicted in Fig. 3(b). X.Y. Xu et al. noted that the addition of TaC can enhance the plastic deformation ability of materials.18) From the experimental results in Fig. 3(b), it is evident that the flexural modulus and TRS values of Ti–10Nb–3Mn–XTaC composites exhibit a similar trend. Specifically, at the same temperature, they initially increase slowly and then decrease with an increase in TaC content. Additionally, under the same TaC addition ratio, the flexural modulus also shows a rising and falling trend as the sintering temperature is increased. In general, a smaller flexural modulus indicates lower stiffness and poorer resistance to deformation, whereas a higher flexural modulus corresponds to greater stiffness and increased resistance to deformation. Notably, the specimen with 2 TaC sintered at 1275°C displayed the highest flexural modulus (40.37 GPa) in this experiment, indicating superior deformation resistance.
Based on the abovementioned discussion and test results, it is justifiable to conclude that the Ti–10Nb–3Mn–XTaC specimens sintered at 1275°C with a 2 mass% TaC addition exhibit favorable sintering characteristics and mechanical properties. In the subsequent experiments, microstructure analysis and corrosion testing will be performed on the Ti–10Nb–3Mn–XTaC composites using samples sintered at the optimal temperature of 1275°C.
3.2 Effect of microstructure and corrosion behaviors on sintered Ti–10Nb–3Mn–XTaC compositesFigure 4 reveals the XRD patterns of various amounts of TaC added to Ti–10Nb–3Mn alloys after sintering at 1275°C for 1 hour. The results show that the main diffraction peaks consist of α-Ti, β-Ti, and TiC. It can be found that TiC diffraction peaks are more shifted towards higher angles, which is the same phenomenon that occurred in our previous study.1,2,13) The main reason is that Nb, Ta, and Mn elements are solid-dissolved into the β phase, which causes internal distortion and interplanar spacing (d) to decrease after sintering. According to the Bragg’s law (nλ = 2d sin θ), as d decreases gradually, θ will increase, thus these peaks shift to high angles.19) It is worth noting that the peaks corresponding to tantalum carbide (TaC, PDF#35-0801) or tantalum (Ta, PDF#04-0788) were not detected and were replaced by the peak of titanium carbide (TiC, PDF#32-1383). Moreover, the XRD analysis did not detect any signals of oxide. In other words, there was no systematic variation of vacuum in the composition of the specimens, particularly with regards to the O contents. According to the Gibbs free energy theory, a lower free energy indicates an easier reaction process. Since the standard free energy of TiC formation is smaller than that of TaC, in situ TaC decomposes into carbon and tantalum during the high-temperature sintering process.20) For example, carbon reacts with titanium to form titanium carbide (TaC + Ti → Ta + TiC), and tantalum acts as a solid solution in titanium, stabilizing the β-phase. This result is supported by the fact that the peak value of TiC increases with an increase in TaC content (as shown in Fig. 4). Furthermore, it can be observed that as the TaC amount increases (1 → 2 → 3 mass%), the peak value of the α phase slightly declines, while the peak value of the β-phase significantly rises. This phenomenon is consistent across samples sintered at temperatures of 1225 and 1250°C and 1300°C.
XRD pattern of various mass% TaC added to Ti–10Nb–3Mn alloys after sintering at 1275°C for 1 h.
As previously mentioned, titanium alloys possess not only high strength and hardness but also excellent corrosion resistance. Therefore, in this experiment, a 1 N H2SO4 solution was utilized for corrosion testing to simulate and evaluate the corrosion resistance of Ti–10Nb–3Mn–XTaC composites in harsh environments. Figure 5 displays the Tafel slope results of the 1275°C-sintered 1 TaC, 2 TaC, and 3 TaC specimens after undergoing 1 N H2SO4 corrosion testing. It is observed that the polarization-corrosion curve shifts to the left with an increase in TaC content, indicating an enhancement in the corrosion resistance of the Ti–10Nb–3Mn–XTaC composites. Titanium readily reacts with oxygen to form a protective oxide film (TiO2), which can quickly regenerate even if it peels off. Additionally, in this study, the addition of 10 mass% Nb element to the titanium alloy promotes the formation of an oxide layer composed of Nb2O5, thereby reinforcing the passivation film of the composite material and effectively improving its corrosion resistance.18)
Tafel results of various mass% TaC added to Ti–10Nb–3Mn alloys after sintering at 1275°C for 1 h.
The experimental results also demonstrated a significant passivation phenomenon in all the specimens, indicating excellent corrosion resistance of the sintered Ti–10Nb–3Mn–XTaC composites. With an increase in TaC content, the corrosion current slightly decreased (5.96 → 5.19 → 4.83 × 10−4 A·cm−2), while the polarization resistance gradually increased (8.81 → 9.41 → 9.99 × 102 Ω·cm2). The results indicate that the 3 TaC specimen exhibited the lowest corrosion current and the highest polarization resistance values, suggesting optimal corrosion resistance. However, the 2 TaC sample showed an appropriate corrosion current and suitable polarization resistance values, demonstrating excellent corrosion resistance.
Figure 6 displays the OM and SEM images of Ti–10Nb–3Mn alloys with various mass% of TaC after sintering at 1275°C for 1 hour. In this study, the alloy prepared is the β-phase + TiC at high temperatures, and the α-phase is formed during the subsequent cooling. Hence, in this context, the definition of the grain boundary pertains to the β-phase matrix. It can be observed that all specimens exhibited typical Widmanstätten-like needle structures.13,21) The 1TaC specimen showed a slender and flaky α-phase, along with a coarser distribution of α-phase in the β-phase matrix. Additionally, coarser α-phase precipitates were observed from the grain boundaries of the β-phase, as depicted in Fig. 6(a). In contrast, the microstructure of the 2 TaC specimen appeared slightly refined, with the flaky α-phase on the grain boundary being more discontinuous and slender compared to the 1 TaC sample. Furthermore, the presence of α-phase within the β-phase matrix became less noticeable, as shown in Fig. 6(b). As for the 3 TaC specimen, the content of the β-phase matrix increased, and with an increase in TaC addition, more TiC carbides precipitated along the grain boundaries, as illustrated in Fig. 6(c). Based on the OM observations, the addition of an appropriate content of TaC (2 mass%) noticeably reduced the presence of slender and flaky α-phase, while increasing the β-phase content and the suitable amount of TiC precipitates. These observations effectively contribute to the improvement of strength in the Ti–10Nb–3Mn–XTaC composites.
The OM and SEM images of various mass% TaC added to Ti–10Nb–3Mn alloys after sintering at 1275°C for 1 h: (a) OM image of 1 TaC, (b) OM image of 2 TaC, (c) OM image of 3 TaC, and (d) SEM images of 2 TaC.
Furthermore, when observing the SEM images of Ti–10Nb–3Mn–xTaC composites, a similar pattern is observed, with the sample Ti–10Nb–3Mn–2TaC serving as an explanatory example. Figure 6(d) presents the SEM image of Ti–10Nb–3Mn–2TaC composites after sintering at 1275°C for 1 hour, and the corresponding EDS analysis results are listed in Table 1. In this study, the constituent elements of Location 1 were identified as solid-solution phases of Ti and Nb, respectively. The presence of dark, discontinuous precipitates is reasonable to be attributed to the slender flaky α-phase. Moreover, based on the previous XRD analysis, the intensity of the diffraction peak associated with the β-phase primarily depended on the TaC content, indicating that an increase in TaC content led to the generation of more tantalum atoms through in situ reductions. Consequently, this increase enhanced the intensity of the β-phase diffraction peak. Therefore, it is reasonable to suggest that Location 2 comprises Ti, Nb, Mn, and Ta elements, with tantalum being a β-phase stabilizing element, implying that the observed precipitates are likely to be in the β-phase state.
In the composition of Location 3, only Ti and C element signals were detected, as shown in Table 1. Among them, Ti accounts for 71.28 at%, and carbon accounts for 28.72 at%. Therefore, it can be preliminarily determined that these precipitates are likely to be TiC, which precipitates in the form of particles along the grain boundaries. However, since the EDS composition analysis results show a Ti/C ratio of 3:1, a range of non-stoichiometric TiCx compounds (between ∼32 at% C and ∼50 at% C in the Ti–C phase diagram) could precipitate on cooling,22) which seem to mean there is some uncertainty and measurement errors in the EDS analysis. Thus, we subsequently use more accurate EBSD to confirm the TiC phase (as see in Fig. 7). Additionally, the compositions of Location 4 consist of Ti, Nb, and C elements, respectively. It is worth noting that the niobium content is higher than that of carbon, with niobium accounting for 10.46 at% and carbon accounting for 4.23 at%. However, the precipitates are still predominantly composed of TiC. Although the reason for the niobium element signal is unknown, it is speculated that Nb may gradually dissolve into the matrix phase. Consequently, the EDS results confirm the previous discussions in this paper, indicating the reaction of carbon with titanium to form titanium carbide, while tantalum acts as a solid solution in titanium, stabilizing the β-phase. Furthermore, this study utilized EBSD to further confirm the existence and crystallographic orientation of TiC precipitates.
The EBSD analysis for various mass% TaC added to Ti–10Nb–3Mn alloys after sintering at 1275°C for 1 h: (a), (c), (e) image quality, (b), (d) and (f) phase mapping, which are 1 TaC, 2 TaC and 3 TaC, respectively.
To further understand the phase distribution and crystal orientation, EBSD analysis was conducted on the 1275°C-sintered Ti–10Nb–3Mn–XTaC specimens. Figure 7 presents the phase distribution results of the EBSD analysis for Ti–10Nb–3Mn alloys with various mass% of TaC after sintering at 1275°C for 1 hour. Figures 7(a), 7(c), and 7(e) show the image quality of the 1 TaC, 2 TaC, and 3 TaC specimens, respectively, while Figs. 7(b), 7(d), and 7(f) reveal the phase mapping of the 1 TaC, 2 TaC, and 3 TaC samples, respectively. The EBSD analysis clearly displays the α-phase (indicated by the red color), the β-phase (indicated by the blue color), and the TiC phase (indicated by the yellow color) in the 1275°C-sintered Ti–10Nb–3Mn–XTaC composites, with the carbides predominantly precipitating along the grain boundaries. It is noteworthy that the proportion of β-phase and TiC phase increases with an increase in TaC content. This finding is consistent with the previous EDS analysis, which confirms the presence of TiC carbides. Moreover, it aligns with the results of XRD analysis (Fig. 4), indicating that as the TaC content increases, the proportions of β-phase and TiC phase also increase. Therefore, it is established that tantalum atoms, generated through the decomposition of TaC, effectively dissolve in Ti alloys, leading to an increase in the β-phase content, while carbon atoms undergo a substitution reaction with titanium, resulting in the precipitation of TiC particles.
Figure 8 reveals the crystal orientation map (COM) from the EBSD analysis of the 1275°C-sintered 2 TaC specimens. Figure 8(b) shows the matrix of β-Ti grains with a BCC structure, displaying different color blocks in the three crystal directions of [001], [111], and [101]. These color variations represent the distribution of crystal grains. While continuous α-phase precipitates are observed at the grain boundaries, the β-Ti phase exhibits random orientation. Moreover, the slender flaky α-Ti with an HCP structure is represented by different color blocks in the three crystal directions of [0001], $[\bar{1}2\bar{1}0]$, and $[01\bar{1}0]$, as shown in Fig. 8(c). This indicates that the α-Ti structure also exhibits random orientation. It is logical to propose that this phenomenon can be attributed to the cooling phase during vacuum sintering. During this phase, the α phase precipitates from the β phase, leading to a matrix within the β phase that is characterized by random orientation. Consequently, the α phase also displays an orientation that is similarly random in terms of crystal structure. Besides, as depicted in Figs. 8(b) and 8(c), the [101] orientation of the β-Ti phase aligns in parallel (//) with the [0001], $[\bar{1}2\bar{1}0]$, and $[01\bar{1}0]$ orientations of the α-Ti phase. Consequently, the α-Ti phase that precipitates within the β-Ti matrix exhibits three distinct grain orientations. This signifies that no specific crystallographic orientation exists for the β and α phases within the context of this study.
The crystal orientation map of the EBSD analysis for Ti–10Nb–3Mn–2TaC composites after sintering at 1275°C for 1 h: (a) all phases, (b) β-Ti, (c) α-Ti, and (d) TiC, respectively.
Figure 8(d) shows a few TiC structures that also lack a specific crystal orientation. However, the observed microstructures distinctly exhibit former β grains, with the crystal orientation showing a significant correlation between these former β grains and the internally formed α-phase. Additionally, the microstructure reveals the presence of the α-phase and TiC formation occurring at the grain boundaries. It is worth mentioning that due to the low solubility of carbon in the Ti alloy (only 0.2%), TiC primarily distributes on the larger carbide particles. Consequently, only a small amount of TiC is dispersed within the β-Ti matrix. This finding aligns with our previous research.1)
To observe whether the TiC precipitates exhibit specific crystal orientations, the distribution map of crystal orientations for the 1275°C-sintered 2 TaC composite was generated using Channel 5 software. A pole figure was then obtained through a complex calculation method. The pole figure was used to analyze the crystal orientations, revealing a high density of data points on the {100}, {110}, and {111} planes, as shown in Fig. 9. This reconfirms that the TiC structures in this study do not exhibit specific crystal orientations. This result is consistent with the previous crystal orientation map obtained from the EBSD analysis.
The TiC pole figure of EBSD analysis for Ti–10Nb–3Mn–2TaC composites after sintering at 1275°C for 1 h.
Based on the above discussion and results, it can be concluded that the microstructure of vacuum-sintered Ti–10Nb–3Mn–XTaC composites, with the addition of an appropriate amount of TaC powder, consists of both α-phase and β-phase structures, along with TiC precipitates.
In this study, compared to the Ti–10Nb–3Mn alloy, the addition of an appropriate amount (2 mass%) of TaC to Ti–10Nb–3Mn–XTaC composites proved effective in reducing the apparent porosity (from 0.43% to 0.34%) and increasing the relative density (from 93.74% to 96.24%). These improvements are beneficial for sintering densification, leading to enhanced mechanical properties and corrosion resistance of the composite materials. The 2 TaC specimens exhibited the desired hardness (from 63.86 HRA to 67.38 HRA), the TRS (from 1671.83 MPa to 1276.63 MPa), and flexural modulus (from 36.69 GPa to 40.37 GPa) after sintering at 1275°C for 1 hour. Additionally, the 2 TaC specimens demonstrated favorable corrosion resistance, with only slight variation in the polarization resistance (from 9.96 × 102 Ω·cm2 to 9.41 × 102 Ω·cm2).
The XRD analysis revealed that the main precipitates in the Ti–10Nb–3Mn–XTaC composites after various sintering temperatures were α-Ti, β-Ti, and TiC. The EBSD results further confirmed the presence of α, β, and TiC phases in the matrix of the Ti–10Nb–3Mn–XTaC composites sintered at 1275°C. The α-Ti phase exhibited a hexagonal close-packed (HCP) structure with crystal directions of [0001], $[\bar{1}2\bar{1}0]$, and $[01\bar{1}0]$. The β-Ti phase, with a body-centered cubic (BCC) structure, showed crystal directions of [001], [111], and [101]. Additionally, the TiC pole figure indicated hot spots on the {100}, {110}, and {111} planes, confirming that the TiC structure had no specific crystal orientation in this study.
This research is supported by the Ministry of Science and Technology of the Republic of China under Grant No. MOST 111-2221-E-027-100-. And the publication founds are subsidized by the National Science and Technology Council of the Republic of China under Grant No. NSTC 112-2221-E-027-035-. The authors would like to express their appreciations for ASSAB Steels Taiwan Co., Ltd. Furthermore, thanks to Prof. H.C. Lin and Mr. C.Y. Kao of Instrumentation Center, National Taiwan University for EBSD experiments.
