The Effect of TiC and Zr Additions on the Microstructure and Mechanical Properties of Ti-30Mo Alloy

Abstract

In this study, Ti-30Mo-nTiC (0, 0.5 wt.%) and Ti-30Mo-0.5TiC-xZr (0, 4, 8, 12 wt.%) alloys were prepared using the powder metallurgy process with the addition of Zr at different rates. The effect of TiC and Zr additions on the microstructure and mechanical properties of the Ti-30Mo alloys were investigated, respectively. The results demonstrated that the addition of 0.5 wt.% TiC significantly improved the density, tensile strength and elastic modulus. The ultimate tensile strength, elongation and elastic modulus of the Ti-30Mo-0.5TiC alloy were determined to be 825 MPa, 7.3% and 112 GPa, respectively. For Ti-30Mo-0.5TiC-xZr alloys, the addition of Zr (8 wt.% or less) results in alloys having a high relative density (>98%), with the density of the alloy decreasing significantly when the Zr content is 12 wt.%. As the Zr content increases, the β phase lattice constant also increases along with the amount of carbide aggregation. This leads to a decrease in the alloy strength, with an increase in the alloy hardness. During high temperature tensile testing at 600 °C, the Ti-30Mo-0.5TiC alloy still had suitable mechanical properties, with its ultimate tensile strength and elongation being 472 MPa and 12.8%, respectively.

1. Introduction

Owing to the high specific strength, low density, enhanced corrosion resistance, excellent biocompatibility, and non-magnetic properties, titanium alloys are widely used in aerospace, petrochemical, energy, marine engineering, pharmaceuticals, and other industries [1,2,3,4]. In comparison to other titanium alloys, Ti-30Mo has demonstrated excellent corrosion resistance in the presence of hot concentrated hydrochloric acid, moderately concentrated sulfuric acid and acidic salt lake brine, giving it a wide application to various engineering fields and applications [5,6]. However, the low elastic modulus of β type Ti-Mo alloys (<80 GPa) [7] reflects their poor resistance to deformation. Meanwhile, due to rapid developments in offshore engineering, the Ti-30Mo alloy cannot meet the increasingly stringent structural requirements. Consequently, there is an urgent need to develop a Ti-30Mo alloy with excellent mechanical properties.

Compared to Ti, using TiH₂ powder as a raw material can significantly improve the diffusion process and accelerate sintering densification [8,9]. In addition, TiH₂ is less likely to react with oxygen and nitrogen in air during the preparation process, which will reduce the impurity content of the final product [10]. In general, titanium alloys can be strengthened by the addition of a second phase and alloying elements. Among the various nanopowder additives reported to strengthen alloys (Al₂O₃, ZrO₂, SiC, BC, Si₃N₄, TiN, TiB₂, etc. [11,12,13,14]), TiC is desirable due to compatibility with Ti and only having a small difference in linear expansion coefficients [15,16,17,18]. These nanopowders also have a high surface energy and can provide the driving force for atomic migration during sintering. Zr and Ti belong to the same group in the periodic table, have the same crystal structure, and are capable of infinite solid solution with each other. In addition, Zr is usually added to Ti alloys to improve the corrosion resistance of the material [19]^. In many studies, Zr has demonstrated a solid solution strengthening effect on Ti-Mo alloys [20]. For example, the tensile strength of the Ti-12Mo-5Zr alloy is approximately 1.0 GPa, with that of the Ti-12Mo-15Zr alloy being approximately 1.1 GPa [21]. However, there are few reports investigating the strengthening of a β type Ti-30Mo alloy through the addition of Zr.

The blended element method is beneficial for the preparation of high performance titanium alloys. In this study, TiH₂ powder was used as the titanium source during the ball milling process. In order to prepare a high performance alloy, Ti-30Mo-nTiC (n = 0, 0.5 wt.%) alloys were prepared by adding TiC nanoparticles, with Ti-30Mo-0.5TiC-xZr (x = 0, 4, 8, 12 wt.%) alloys prepared by adding ultra-fine Zr powder. The effects of both the TiC and Zr contents on the microstructure and mechanical properties of β type Ti-30Mo alloys were reported to provide a reference for expanding the application of these materials. It may be developed into an alternative structural material under high temperature conditions in offshore engineering, chemical industry and other fields.

2. Experimental Procedure

2.1. Material Preparation

Ti-30Mo-nTiC (n = 0, 0.5 wt.%) and Ti-30Mo-0.5TiC-xZr (x = 0, 4, 8, 12 wt.%) alloys with different Zr contents were prepared by the powder metallurgy method using TiH₂, Mo, Zr and TiC powder as raw materials. The morphology of the original powders is presented in Figure 1, with TiH₂ (particle size < 10 μm) and Zr powder (particle size < 1 μm) having irregular shapes. Mo (particle size < 2 μm) and TiC powder (particle size < 100 nm) were observed to be nearly spherical. The TiH₂, Mo, Zr and TiC powders were mechanically mixed in a ball mill (QM-3SP2, Chishun Tech, Nanjing, China) in proportion to the alloy composition under argon atmosphere for 5 h. During ball milling, the ball to powder weight ratio was 1:1, the cemented carbide ball as grinding ball, and the ball milling speed was 200 r/min. The ball-milled mixture was loaded into a rod-shaped rubber sheath, after which the rod-shaped forming blanks were prepared by pressing them under 170 MPa pressure for 60 s in a cold isostatic press. The green compacts were then sintered in two steps under high vacuum (below 10⁻³ MPa), H was removed completely by holding at 700 °C for 2 h. Finally, it was sintered at 1350 °C for 2 h, followed by furnace cooling. The sintered samples were prepared by wire cut electrical discharge machining (WEDM) to the dimensions required for the test. For convenience of expression, the prepared Ti-30Mo-0.5TiC-xZr alloys were expressed using the abbreviated form of TxZ, i.e., T0 for Ti-30Mo-0.5TiC, T4Z for Ti-30Mo-0.5TiC-4Zr, T8Z for Ti-30Mo-0.5TiC-8Zr, and T12Z for Ti-30Mo-0.5TiC-12Zr.

2.2. Characterization

The morphology of the powders was characterized by transmission electron microscopy (TEM, Titan G2 60-300, FEI, Hillsboro, OR, USA). The density values and porosity of the samples prepared were tested using an electronic balance (MSA324S-000-DU, Sartorius, Göttingen, Germany), and the theoretical density of the alloys was calculated according to the law of mixtures. Subsequently, the relative density of the samples was obtained through the ratio between the actual density and the theoretical density. The phase composition of the samples was tested using an X-ray diffractometer (XRD, Bruker D8 Advance) under the following conditions: operating voltage of 40 KV, current of 30 mA, anode target of Cu, Kα radiation (λ = 0.154 nm), and a test range of 20° to 80°. The morphology and microstructure of the sample and fracture surfaces were characterized by scanning electron microscopy (SEM, Nova Nano250, FEI, Hillsboro, OR, USA), and energy dispersive spectrometry (EDS) was undertaken for elemental analysis. Prior to XRD testing and SEM observation, samples were polished to 2000 mesh with SiC sandpaper and polished with 0.5 μm diamond polish to a reflective finish.

The mechanical properties of the alloy were characterized by performing hardness measurements, with tensile tests conducted at room temperature and 600 °C. A microindentation hardness tester (BUEHLER 5104, Buehler, Lake Bluff, IL, USA) was used to measure the Vickers hardness of the samples, and the test procedure was performed using a diamond indenter held at a load of 0.3 Kg for 10 s. For each sample, 10 individual microzones were tested and the reported result was the average value. Tensile tests (room temperature or 600 °C) were performed using an Instron 3369 mechanical testing machine with an initial strain rate of 1 mm/min. The shape and dimensions of the tensile specimens (in mm) are shown in Figure 2. The Young’s modulus of the samples was measured by mounting an extensometer during the room-temperature tensile test.

3. Results and Discussion

3.1. Analysis of Microstructure and Mechanical Properties of Ti-30Mo-nTiC Alloys

3.1.1. Microstructure Analysis

The relative densities and porosities of the Ti-30Mo and Ti-30Mo-0.5TiC alloys are reported in

Table 1. Relative density and porosity of Ti-30Mo-nTiC alloys.

Alloys	Density (g/cm3)	Relative Density (%)	Porosity (%)
Ti-30Mo	4.88 ± 0.02	91.2 ± 0.4	11 ± 0.3
Ti-30Mo-0.5TiC	5.34 ± 0.01	98.3 ± 0.2	0.3 ± 0.1

. For the Ti-30Mo alloy, the density and relative density are low. The sintering temperature required for the alloy to reach the density was increased due to the elevated melting point of Mo, with the holding time of 2 h at a sintering temperature of 1350 °C insufficient to obtain a high relative density. However, after the addition of TiC, the relative density of the alloy increased significantly with a decrease in the porosity. This demonstrated that the addition of TiC can promote the sintering densification process and significantly increase the relative density of Ti-30Mo alloy.

The metallographic photographs of the Ti-30Mo and Ti-30Mo-0.5TiC alloys without etching are presented in Figure 3. In Figure 3a, numerous black spots are distributed throughout the image, indicating the presence of pores in the Ti-30Mo alloy. In comparison, there are very few pores in the sample with TiC added, demonstrating that this addition can reduce the porosity of the alloy. These images align well with the porosity measurements reported in Table 1.

Figure 4 shows the EDS results of the corresponding areas of the Ti-30Mo-0.5TiC alloy. Area1 is the matrix phase and Area2 is the particle phase, approximately 10–20 µm in size. It can be seen that the matrix phase contained the elements Ti and Mo in addition to C (Figure 4b). The particle phase (Area2) mainly contained Ti and C, with trace amounts of Mo present, indicating that during the sintering process TiC diffuses into the Ti-Mo matrix with this reaction accelerating sintering densification [22] and increasing the relative density of the alloy.

The binary phase diagram of the Ti-C low carbon region [23] is presented in Figure 5 where it can be seen that both α-Ti and β-Ti phases of TiC had a degree of solid solution, with the solid solution region presented in the light blue (Figure 5). When TiC was added (0.5 wt.%), the molar percentage of TiC in the Ti-30Mo system was approximately 0.55% and the atomic percentage of C would be lower, resulting in the ratio of Ti to C in the single-phase solid solution zone (Figure 5). The evidence demonstrated that the microstructure of the Ti-30Mo alloy was a single β phase, and the solid solution diffusion of TiC in the β-phase increased the sintering density of the alloy.

3.1.2. Mechanical Properties Analysis

The mechanical properties of the Ti-30Mo and Ti-30Mo-0.5TiC alloys at room temperature are reported in

Table 2. Mechanical properties of Ti-30Mo-nTiC alloys at room temperature.

Alloys	HV0.3	Young’s Modulus (GPa)	Ultimate Tensile Strength (MPa)	Elongation (%)
Ti-30Mo	152 ± 11	84 ± 2	654 ± 27	3.8 ± 0.4
Ti-30Mo-0.5TiC	291 ± 4	112 ± 1	825 ± 12	7.3 ± 0.2

. It is evident that the hardness, elastic modulus, tensile strength and elongation of the Ti-30Mo-0.5TiC are significantly better than those of Ti-30Mo, indicating that the addition of trace quantities of TiC is effective in improving the mechanical properties. The microhardness of the matrix was measured to have increased from 152 HV_0.3 in Ti-30Mo to 291 HV_0.3 in Ti-30Mo-0.5TiC, likely due to the significantly higher relative density of Ti-30Mo-0.5TiC. There is also evidence that TiC diffuses into the matrix, and the solid solution of C in the lattice increases the microhardness (Figure 4). The elastic modulus is generally related to the phase composition of the material. β-type Ti-Mo has a low elastic modulus (<80 GPa) [7], and the TiC phase has a relatively high elastic modulus (450 GPa) [24] according to the law of mixtures, as presented in Equation (1).

E = E_Ti-Mo × V_Ti-Mo + E_TiC × V_TiC

(1)

where E_Ti-Mo, E_TiC, V_Ti-Mo and V_TiC represent the modulus of Ti-Mo matrix and TiC and the volume fraction of Ti-Mo matrix and TiC, respectively. The theoretical volume fraction of TiC can be calculated according to Equation (2).

$${\mathrm{V}}_{\mathrm{TiC}}=\frac{{\mathrm{v}}_{\mathrm{TiC}}}{{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{v}}_{\mathrm{i}}} \times 100\%$$

(2)

where V_TiC is the theoretical volume fraction of TiC, v_TiC is the volume of TiC, and v_i is the volume of each component.

The volume fraction of the 0.5 wt.% TiC calculated according to Equation (2) is approximately 0.55%, with the elastic modulus of the Ti-30Mo alloy experimentally measured to be 84 GPa. The Ti-30Mo-0.5TiC alloy elastic modulus was calculated to be 86.0 GPa, while the elastic modulus of Ti-30Mo-0.5TiC alloy was experimentally measured to be 112 GPa, significantly larger than the calculated result. In fact, the mixture calculation is an ideal calculation resulting from the simple superposition of the physical and chemical properties of each component, without considering the mutual diffusion reaction between components. This difference further confirms that TiC diffuses into the matrix, thereby significantly increasing the elastic modulus of Ti-30Mo alloy.

The ultimate tensile strength and elongation of Ti-30Mo-0.5TiC were 825 MPa and 7.3%, respectively, which are 26.1 and 92.1% higher when compared to Ti-30Mo (Table 2). The significant increase in the ultimate tensile strength and elongation of the alloy is due to the dispersion strengthening effect of the TiC particles in addition to the significant increase in relative density.

3.2. Analysis of Microstructure and Mechanical Properties of Ti-30Mo-0.5TiC-xZr Alloys

3.2.1. Microstructure Analysis

The relative density of the Ti-30Mo-0.5TiC-xZr alloys with variations in Zr content is presented in Figure 6. It can be seen that the relative density of the alloy was basically stable when the Zr content ranged from 0–8 wt.%, but when the Zr content was 12 wt.%, there was a significant decrease in the relative density. This is likely due to the large difference in the atomic radii of Zr and Ti (0.216 nm for Zr and 0.147 nm for Ti), leading to a high resistance to the replacement of Ti atoms in the lattice by Zr atoms. Once Ti atoms are replaced by Zr, the lattice distortion energy will be increased, resulting in a high-energy unstable state. In general, the system will eventually shift to a lower energy state spontaneously, leading to increased resistance to atomic substitution. This is not conducive to diffusion of atoms, so sintering of the alloy becomes more difficult to densify after the addition of Zr. The higher the Zr content, the higher the sintering temperature required to achieve densification. A sintering temperature of 1350 °C for 2 h, is sufficient to densify the alloy at a Zr addition of 8 wt.% or less, while Ti12Z presents additional challenges.

The XRD patterns of the Ti-30Mo-0.5TiC-xZr alloys are presented in Figure 7. The alloy mainly consists of a single β phase, with no trace of TiC detected. This is due to the addition of a relatively high content of Mo. According to the Ti-Mo binary phase diagram, the critical Mo content for complete stabilization of the β phase at room temperature (300 k) is 11.2 at.% [23]. In addition, Zr in β-Ti alloys also has the effect of enhancing the stability of the β phase [25,26]. This is reflected in the XRD peaks, as when the content of Zr in the alloy increases, the diffraction peaks are slightly shifted to a lower 2θ angle. This occurs due to the atomic radius of Zr being larger than that of Ti, and the solid solution of Zr in the Ti lattice expanding its lattice constant.

The SEM imaging in backscattered electron mode of the Ti-30Mo-0.5TiC-xZr alloys is presented in Figure 8. The alloys are composed of the matrix phase with carbide particle inclusions present. There is a tendency for the carbide particles to aggregate as the Zr content increases. This is due to the large difference in atomic radius of Zr replacement solid solution in the Ti-Mo lattice, making a large elastic distortion around the Zr atom. With the increase of Zr content, the elastic distortion generates more stress field regions, making the diffusion of TiC difficult and enhancing aggregation of the carbide particles.

Table 3. The results of EDS elemental analysis for each area in Figure 8.

Alloys	Areas	Elements (at. %)
		Ti	Mo	C	Zr
T0	Area1	74.7	12.3	13.0	-
	Area2	69.0	0.3	30.7	-
T4Z	Area1	71.1	12.6	14.2	2.1
	Area2	66.3	0.2	31.8	1.7
T8Z	Area1	66.7	12.1	17.7	3.5
	Area2	64.4	0.3	32.1	3.2
T12Z	Area1	58.8	11.4	24.5	5.3
	Area2	52.3	0.3	32.5	4.9

shows the EDS element analysis results of the corresponding areas in Figure 8, showing the element contents of the matrix phase and particle phase of the TxZ alloys, of which area1 is the matrix phase and area2 is the particle phase. It can be seen that Mo is mainly distributed in the matrix phase, while the content in the particle phase is very small. With the addition of Zr content, the content of the C element in the particle phase gradually increases, indicating that the addition of Zr will hinder the diffusion of C element, which shows that Zr addition will increase the degree of carbide aggregation.

SEM mapping was performed on the Ti-30Mo-0.5TiC-12Z alloy to confirm the distribution of key elements. Figure 9 presents the SEM mapping image of the area in Figure 8d and the distribution of each element. The color intensity represents the mass fraction of the elements, with purple indicating Ti, blue Mo, green C and tan Zr. In various studies, the low diffusion rate of Mo has been reported to produce a bias of Mo in Ti matrices [27]. In this study all elements show a uniform distribution in the matrix.

The fracture morphology of Ti-30Mo-0.5TiC-xZr alloys after tensile testing at room temperature is presented in Figure 10. In Figure 10a–e, all images show a fracture through the crystal. As shown by the arrows in Figure 10b, a small number of dimples can be seen in T0, confirming that the T0 alloy has good toughness characteristics. The number of dimples decreased with the addition of Zr, indicating an increase in alloy brittleness. In Figure 10a–d, there are small particles located around the large grains of the matrix, as shown in the red dashed box. These may be carbides distributed on grain boundaries which will cause brittle fractures as cracks expand along the weakened grain boundaries. In Figure 10e, the small particles are no longer distinguishable due to the presence of a large number of pores.

3.2.2. Analysis of Room Temperature Mechanical Properties

The variation in microhardness and Young’s modulus of the alloy with added Zr is reported in Figure 11. As the Zr content was increased, the microhardness of the alloy gradually increased from 304 HV_0.3 in T0 to 334 HV_0.3 in T12Z. Zr atoms dissolved in the matrix lattice will lead to lattice distortion, increasing the resistance to dislocation motion and making slip difficult. This will increase the hardness of the β-Ti matrix. As the Zr content was increased, the elastic modulus of the alloy was measured to decrease gradually from 112 GPa in T0 to 94 GPa in T12Z. The elastic modulus depends on the magnitude of the interatomic bonding force. As the alloy primarily consists of a single β phase, the solid solution of Zr becomes the main factor affecting the elastic modulus of the material. The solid solution of Zr has a large atomic radii in the Ti lattice, reducing the interatomic bonding force and decreasing the elastic modulus.

The variation of tensile strength and elongation of Ti-30Mo-0.5TiC-xZr alloy with Zr content is presented in Figure 12. The ultimate tensile strength and elongation of the alloy show a significant decrease as Zr content is elevated, with the tensile strength and elongation decreasing from 846 MPa and 7.8% in T0 to 364 MPa and 3.6% in T12Z, respectively. Although the alloy has a high relative density in the presence of small amounts of Zr (4, 8 wt.%), the addition of Zr will make the carbides tend to aggregate and grow. The carbides tend to aggregate in defective areas, such as grain boundaries and pores, weakening the grain boundary strength and reducing the tensile strength and elongation of the alloy.

3.2.3. High Temperature Tensile Properties and Fracture Model at 600 °C

The stress–strain curves of T0 and T4Z alloys in tensile tests at 600 °C are reported in Figure 13. The addition of Zr reduced the strength and elongation at high temperature. The tensile strength and elongation at the fracture of the T0 alloy at 600 °C decreased from 472 MPa and 12.8% to 364 MPa and 7.9% in T4Z, respectively. The atomic diffusion rate is higher at the elevated temperature in comparison to room temperature, with the dislocation lines more likely to break free from the pinned stress field and move during dislocation slip. This results in the strength of the alloy at high temperature being lower than that at room temperature, while the elongation is higher. A slight sawtooth fluctuation is produced on the stress–strain curve of T0 alloy, with this characteristic enhanced on the curve of the T4Z alloy. This phenomenon can be explained by dynamic strain aging (DSA) [28], where the microscopic mechanism of DSA results from the interaction between solute atomic gas clusters and movable dislocations. In this situation, this would represent the solid solution atoms moving towards the dislocations during diffusion, causing the material strength to increase briefly, following which the movable dislocations continue to move away from the solute atoms and the material strength decreases. Compared with T0, the stress field generated by the solid solution in the matrix lattice has a fixed effect on the dislocations due to the large size difference between the solute atoms Zr and Ti in T4Z alloy. Once the Zr atoms are freed from the fixed gas clusters, the dislocations will move more easily, resulting in a greater stress drop and more obvious sawtooth fluctuations in the stress–strain curve. In addition, during the earlier stage of deformation, the dislocation density increases abruptly, leading to hardening. However, with the extension of the deformation time, the temperature increases the atomic driving force and reduces the resistance to dislocation motion, while dynamic reversion and recrystallization occur, making dynamic softening dominant, reducing the deformation resistance and leading to a larger sawtooth fluctuation in the stress–strain curve.

The fracture morphology of the T0 and T4Z alloys after tensile test at 600 °C is presented in Figure 14. Numerous small dimples are evident in Figure 14a–d, indicating that the fracture mode of both alloys is dominated by ductile fracturing. The fracture patterns present in Figure 14b are deeper than those in Figure 14d, indicating that the Ti-30Mo-0.5TiC alloy has a relatively improved high-temperature toughness. This is consistent with the results reported in Figure 13. It is worth noting that there are flat and fine sections present, as indicated by the red arrows. In Figure 14b, there are flat and fine sections, while in Figure 14d, the flat sections are coarser, indicating that brittle fracturing occurred in these regions and that the addition of Zr caused aggregation of these brittle phases. These flattened sections may be the result of brittle fracturing of the carbide particles, with aggregation of brittle carbide particles reducing the strength of the alloy at high temperature.

4. Conclusions

In this study, the β type Ti-30Mo-nTiC (n = 0, 0.5 wt.%) and Ti-30Mo-0.5TiC-xZr (x = 0, 4, 8, 12 wt.%) alloys were prepared using a powder metallurgy process, with the effects of TiC and Zr additions on the microstructure and mechanical properties investigated. The key findings of this work are summarized below:

• The diffusion and particle strengthening effect of TiC with the addition of 0.5 wt.% TiC to Ti-30Mo significantly increased the relative density of the alloy, while its elastic modulus increased from 84 GPa to 112 GPa and tensile strength increased from 654 MPa to 825 MPa. Meanwhile, the microhardness and elongation of Ti-30Mo-0.5TiC alloy were 291 HV_0.3 and 7.3%, respectively.
• The β-phase lattice constant of Ti-30Mo-0.5TiC-xZr alloys increased with the increase of Zr content. This led to an increased amount of carbide particle aggregation, decreasing the tensile strength and elongation of the alloys.
• During high-temperature tensile testing at 600 °C, both Ti-30Mo-0.5TiC and Ti-30Mo-0.5TiC-4Zr alloys exhibited a distinct ductility fracture pattern, but Ti-30Mo-0.5TiC had superior high-temperature properties with a tensile strength and elongation of 472 MPa and 12.8%, respectively.