The Influence of Mo and Y on the Microstructure and Properties of TiZrHfNb Series Refractory High-Entropy Alloys

Abstract

TiZrHfNbMo refractory high-entropy alloy has poor plasticity and a relatively high density at room temperature, which limits its wide industrial application. To develop RHEAs featuring a simple structure, low density, and excellent overall performance. In this study, three refractory high-entropy alloys of Ti₂₂Zr₂₅Hf_(35−x)Nb₁₈Mo_x (x = 10, 15, 20) were preliminarily designed, and the effects of different Mo contents on their microstructure and properties were investigated. All three components are of a single-phase BCC structure. Room-temperature and high-temperature compression and friction wear tests show that, with the increase in Mo, the solid solution strengthening effect is enhanced. The room temperature yield strength increases by 35.1%, the high-temperature yield strength increases by 222.6%, and the wear rate decreases by 60.5%. However, the room temperature fracture strain of Mo20 is reduced to 24%. Therefore, Y was introduced into the Mo20 refractory high-entropy alloy at mass fractions of 0.1% and 0.2% to further enhance its plasticity. The experimental results show that the addition of Y, through grain refinement and solid solution strengthening, simultaneously enhances room-temperature plasticity (from 24% to 323%) and wear resistance. The findings furnish a theoretical framework for rational element selection in RHEAs design.

1. Introduction

In 2004, Ye et al. initially proposed the high-entropy alloy (HEA) system [1]. High-entropy alloys are generally composed of five or more alloying elements in equimolar or near-equimolar ratios, with a configurational entropy greater than 1.61R [2]. The multi-principal element properties of HEAs give them the unique high-entropy effect, sluggish diffusion effect, lattice distortion effect, and cocktail effect [3]. These four effects impart a simple crystal structure and superior mechanical properties [4]. With the ongoing advancement of research on high-entropy alloys, various research systems have been established. According to the constituent elements, they are categorized into lightweight high-entropy alloys, 3d group high-entropy alloys, refractory high-entropy alloys, and others [5,6]. Refractory high-entropy alloys (RHEAs) comprise refractory elements, including Ti, Zr, Hf, W, Nb, Mo, Ta, V, and others [7,8,9]. RHEAs have received widespread attention due to their superior wear resistance [10,11], mechanical properties [12,13], and promising potential for high-temperature applications [13,14,15], but their low plasticity and high density restrict their applications [16,17,18,19]. Among various refractory high-entropy alloy systems, it has been found that the TiZrHfNb series of refractory high-entropy alloys simultaneously exhibit room-temperature and high-temperature strength, as well as low density and solid matrix stability [20,21]. While retaining the aforementioned superior properties, they also exhibit moderate room-temperature plasticity and are anticipated to supplant nickel-based superalloys as the next generation of superalloys for aero engines. In addition, incorporating Mo elements into the TiZrHfNb series of refractory high-entropy alloys can improve their strength and high-temperature performance, but their room-temperature plasticity requires further enhancement.

Rare earth (RE) elements are recognized as “industrial vitamins” or “industrial monosodium glutamate”. RE elements can achieve the concurrent improvement of strength and plasticity through mechanisms such as grain refinement, solid solution strengthening, grain boundary purification, and suppression of brittle intermetallic compounds [22,23,24]. Jiang et al. developed the AlNbTiVY refractory high-entropy alloy incorporating rare earth elements with a density of just 5.18 g/cm³, and owing to the synergistic effect of lower-density elements, including Al, Ti, and Y, it exhibits a high specific strength [25]. Yao et al. incorporated the rare earth element Ce into Nb_0.5TiZrV_0.5. When the Ce content was 0.005 at.%, the average grain size was refined from 339.2 μm to 243.7 μm, and the ductility rose from 4.32% to 7.15% [26]. Zhang et al. added 0.3 at.% rare earth Y to CoCrFeNi, and both the Vickers hardness (from 146 HV to 400 HV) and yield strength (from 202 MPa to 1440 MPa) were substantially improved [27]. These findings demonstrate the feasibility of incorporating rare earths into high-entropy alloys to enhance their performance.

However, most studies mainly focus on the influence of different Mo contents on room-temperature strength and hardness, and there are almost no studies on adding rare earth Y to TiZrHfNbMo. This text investigates the effects and mechanisms of action of different Mo and Y contents on the microstructure and properties of RHEAs; the strength and plasticity were synergistically enhanced by regulating the content of Mo and Y. In this study, the TiZrHfNb refractory high-entropy alloy was selected as the matrix, and the Mo element and rare earth Y element were sequentially incorporated to optimize its composition. Initially, the optimal Mo content was determined, and subsequently, the rare earth element Y was further incorporated to investigate the mechanisms underlying the synergistic effect of the Mo element and rare earth Y on the properties of the high-entropy alloy. These findings can provide a theoretical foundation for the incorporation of elements in refractory high-entropy alloys.

2. Experimental Materials and Methods

2.1. Material Preparation

The materials used in this study are refractory high-entropy alloys prepared by Beijing Yanbang New Materials Technology Co., Ltd. (Beijing, China), including five refractory elements Ti, Zr, Hf, Nb, Mo, and rare earth element Y, with a purity of ≥99.9 wt%. In this experiment, three groups of Ti₂₂Zr₂₅Hf_(35−x)Nb₁₈Mo_x (x = 10, 15, 20) alloys with varying Mo contents and two groups of Ti₂₂Zr₂₅Hf₁₅Nb₁₈Mo₂₀ (0.1Y/0.2Y) alloys with different Y contents were prepared through vacuum arc melting. The furnace chamber is evacuated to 3 × 10⁻³ Pa, argon gas is injected with a purity of 99.9% to 0.5 atmospheres, and then the smelting process is started. Before smelting, an appropriate amount of alloy is supplemented according to the composition and the number of smelting times to make up for the loss during the smelting process. To achieve alloys with uniform composition, all alloy samples are subjected to repeated smelting at least 8 times. The total melting time is approximately 15 min, and the melting current is controlled at approximately 600 A. After smelting is complete, the alloy is allowed to cool in a water-cooled copper crucible for approximately 10 min. Once fully cooled, the alloy ingot is removed to obtain a button-shaped ingot with uniform composition. It was found in

Table 1. Theoretical components and practical components.

Alloys		Elements (at.%)
		Ti	Zr	Hf	Nb	Mo
Mo10	Design	22	25	25	18	10
	Actual	25.08	21.61	26.07	17.45	9.80
Mo15	Design	22	25	20	18	15
	Actual	24.84	21.00	19.78	17.81	16.57
Mo20	Design	22	25	15	18	20
	Actual	24.42	24.35	16.38	15.58	19.39
0.1Y	Design	22	25	15	18	20
	Actual	24.13	23.38	16.34	16.72	19.41
0.2Y	Design	22	25	15	18	20
	Actual	23.69	24.11	15.12	16.30	20.78

that the majority of elements deviated slightly from their actual contents, while Ti was slightly in excess, which might be due to excessive addition during the element supplementation process.

2.2. Microstructure Analysis

Before the microstructure analysis, the samples were sequentially mechanically ground with SiC sandpapers of #400, #800, #1200, and #2000, and subsequently, the sample surfaces were polished to a mirror-like finish with diamond polishing paste. The phases of various compositions were analyzed using an X-ray diffractometer (Cu-Kα, Ultima-IV, Rigaku Corporation, Tokyo, Japan). The 2θ range is from 20° to 100°, with a scanning speed of 2°/min and a scanning step size of 0.02°. The elemental distribution of the samples was observed using electron probe X-ray microanalyzer (EPMA, JXA-ISP100, Rigaku Corporation, Tokyo, Japan). The polished samples were etched with a corrosive solution of HF: HNO₃: H₂O = 1:2:4. Under room temperature conditions, the polishing surface of the sample is covered with the corrosive solution and etched for 2 to 5 s. The morphology of the etched samples was observed using a field emission scanning electron microscope (SEM, Zeiss-SIGMA, Carl Zeiss AG, Oberkochen, Germany), and the elemental distribution was analyzed using energy dispersive spectroscopy (EDS); the acceleration voltage is 20 kV. After mechanically polishing the samples, the microstructure of the as-cast RHEAs with various compositions was characterized by electron backscatter diffraction (EBSD, symsymmetry-S2, Carl Zeiss AG, Oberkochen, Germany), the acceleration voltage is 20 kV, and the scanning step size is 2.4 μm. The electrolyte comprises 6% perchloric acid, 35% n-butanol, and 59% methanol. At a temperature of 25 °C, the electrolytic polishing is performed for 5 to 10 s with a voltage of 25 V and a current of 1.5 A.

2.3. Mechanical Property Tests and Room-Temperature Friction and Wear Experiments

The hardness test of the sample was conducted using a Vickers hardness tester (200HVS-5, WOW, Shanghai, China), with a load of 3000 g and a force holding time of 15 s. Seven points were tested at the center of the polished surface of the sample. To ensure accuracy, the highest and lowest values were excluded, and the average of the remaining measurements was calculated. Cylindrical specimens of Φ6 × 10 mm were cut out on the ingot by wire electrical discharge machining equipment. The room-temperature compression test and high-temperature compression test were carried out, respectively, by using the electronic universal testing machine (Zwick-Z250, WOW, Shanghai, China) and the Gleeble-3500 (Dynamic Systems Inc., Poestenkill, NY, USA) thermal simulation testing machine. The high-temperature compression test was conducted under vacuum conditions. Before the high-temperature compression test, heat the sample to 1000°C at a heating rate of 10°C/s in Gleeble-3500. After heating to 1000 °C, the temperature was maintained for 5 min to ensure uniform heating of all parts of the sample. After maintaining the temperature, the sample was compressed to 50% of its original height at a compression rate of 1 × 10⁻³/s. The sample size and compression rate at room-temperature compression are the same as those in the high-temperature compression experiment.

Square samples of 10 × 10 × 3 mm were extracted from the ingot using wire electrical discharge machining equipment. The room-temperature friction and wear test was performed using a friction and wear testing machine (Bruker CETR UMT-2, Bruker Corporation, Billerica, MA, USA). Before testing, the surface of the testing samples was polished down to a 2000-grit SiC paper to eliminate surface contamination. All friction tests were conducted at room temperature, and the friction time is 30 min. The friction mode was linear reciprocation, with a load of 5 N, a friction rate of 2 Hz, and a friction stroke of 2 mm, and the friction pair consisted of 6 mm silicon nitride ceramic balls. The three-dimensional topography of the worn surface was analyzed using the WIVS-type white light interference three-dimensional profilometer to obtain wear volume loss data. To prevent the inhomogeneity of the as-cast sample structure from affecting the accuracy of the test, the mechanical property test was repeated three times.

3. Results and Discussion

3.1. The Microstructure and Properties of Mo_x

3.1.1. Phase and Microstructure Analysis of Mo_x

Figure 1 displays the XRD patterns of Mo10, Mo15, and Mo20. The figure reveals that the alloys with three compositions all exhibit a single-phase body-centered cubic (BCC) structure. The peak positions of the XRD pattern shift rightward with increasing Mo content. The lattice parameter of these components is as follows: Mo10 (d = 3.378), Mo15 (d = 3.376), and Mo20 (d = 3.329). This is primarily because the atomic radius of Mo is smaller than that of other constituent atoms. With increasing Mo content, the average atomic radius of RHEAs decreases, resulting in a decreased lattice constant, which causes the peak positions of the XRD pattern to shift rightward.

When designing alloys, thermodynamic parameters can be evaluated to forecast phase formation. The high configurational entropy of high-entropy alloys can lower the Gibbs free energy of solid solutions and facilitate solid solution formation, but configurational entropy and Gibbs free energy are not the sole criteria for facilitating solid solution formation. Yang et al. defined the atomic size difference (δ) and parameter (Ω) as the key parameters for forecasting the structure of high-entropy alloys [28]. Within a certain range, the larger the atomic size difference, the better the solid solution strengthening effect. Research indicates that when Ω ≥ 1.1 and δ ≤ 6.6%, stable solid solution phases are readily formed. Zhang et al. found that when −20 ≤ ∆H_mix ≤ 5 kJ/mol and δ ≤ 6.6%, high-entropy alloys are likely to form solid solution structures [29]. Guo et al. proposed the valence electron concentration (VEC) standard, which was found to be a physical parameter for the stability of BCC or FCC solid solution phases [30]. The BCC phase is stable when VEC ≤ 6.87, and in single-phase RHEAs, the ductility can be enhanced by reducing VEC by regulating the element content [31].

$${∆\mathrm{S}}_{\mathrm{m}\mathrm{i}\mathrm{x}}=-\mathrm{R}{\sum }_{\mathrm{i}=1}^{\mathrm{n}}\left({\mathrm{c}}_{\mathrm{i}}{\mathrm{l}\mathrm{n}\mathrm{c}}_{\mathrm{i}}\right)$$

(1)

$${\Omega }_{ij}=4{\mathsf{\Delta }\mathrm{H}}_{\mathrm{A}\mathrm{B}}^{\mathrm{m}\mathrm{i}\mathrm{x}}$$

(2)

$${∆\mathrm{H}}_{\mathrm{m}\mathrm{i}\mathrm{x}}={\sum }_{\mathrm{i}=1,\mathrm{i}\ne \mathrm{j}}^{\mathrm{n}}{\mathsf{\Omega }}_{\mathrm{i}\mathrm{j}}{\mathrm{c}}_{\mathrm{i}}{\mathrm{c}}_{\mathrm{j}}$$

(3)

$$\mathsf{\delta }=\sqrt{{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{(1-{\displaystyle \frac{{\mathrm{r}}_{\mathrm{i}}}{\overline{\mathrm{r}}}})}^2}$$

(4)

$$\mathsf{\Omega }={\displaystyle \frac{{\mathrm{T}}_{\mathrm{m}}{∆\mathrm{S}}_{\mathrm{m}\mathrm{i}\mathrm{x}}}{\left|{∆\mathrm{H}}_{\mathrm{m}\mathrm{i}\mathrm{x}}\right|}}$$

(5)

$$\mathrm{V}\mathrm{E}\mathrm{C}={\sum }_{\mathrm{i}=1}^{\mathrm{N}}{\mathrm{c}}_{\mathrm{i}}{\mathrm{V}\mathrm{E}\mathrm{C}}_{\mathrm{i}}$$

(6)

$${\mathrm{T}}_{\mathrm{m}}={\sum }_{\mathrm{i}=1}^{\mathrm{n}}{\mathrm{c}}_{\mathrm{i}}{\left({\mathrm{T}}_{\mathrm{m}}\right)}_{\mathrm{i}}$$

(7)

c_i represents the molar percentage of the i-th element, R is the gas constant, 8.314 JK⁻¹mol⁻¹, and ΔS_mix is the configurational entropy. ${\mathsf{\Omega }}_{\mathrm{i}\mathrm{j}}=4{∆\mathrm{H}}_{\mathrm{A}\mathrm{B}}^{\mathrm{m}\mathrm{i}\mathrm{x}}$ represents the interaction parameter between the i-th and j-th elements, and ${∆H}_{AB}^{mix}$ is the mixing enthalpy of the binary alloy. T_mi represents the melting point of the i-th element of the alloy. r_i represents the atomic radius of the i-th element, and $\overline{\mathrm{r}}$ represents the average atomic radius. The calculation results are presented in

Table 2. Parameters for the formation of single-phase BCC solid solutions in high-entropy alloys.

Alloys	ΔSmix	∆Hmix	δ	Ω	VEC	ρ	Tm
Mo10	13.01	−0.027	5.19	1132.7	4.38	8.52	2095
Mo15	13.26	−0.943	5.42	33.6	4.48	8.36	2114
Mo20	13.26	−1.779	5.57	17.9	4.58	8.21	2134

. All parameters satisfy the criteria for forming single-phase BCC solid solutions and align with the results in Figure 1. Furthermore, they exhibit a high melting point (≥1750 °C) and a low density (≤8.5 g/cm³).

Figure 2a–c displays the microstructure of three alloy compositions after etching. The figure reveals that the microstructure of all three alloys is dendritic. As presented in

Table 3. Mo_x dendrites and element contents between dendrites.

Alloy Regions		Elements (at.%)
		Ti	Zr	Hf	Nb	Mo
Mo10	DR	25.10	21.83	26.31	17.34	9.42
	ID	25.62	21.87	26.29	16.98	9.24
Mo15	DR	24.53	20.08	20.23	18.42	16.72
	ID	25.43	21.79	19.89	16.82	16.17
Mo20	DR	25.65	22.80	15.35	16.71	19.49
	ID	26.53	25.93	15.18	14.19	18.17

, some elements are biased, the alloy dendrites (DR) are enriched with Nb and Mo elements, while the interdendrites (ID) are enriched with Ti and Zr elements. The distribution of Hf elements is comparatively uniform. This is attributed to the differences in the melting points of the elements. The high-melting-point Nb and Mo elements solidify within the dendrites first, and the low-melting-point Ti and Zr elements solidify within the interdendrites subsequently.

3.1.2. Mechanical Performance of Mo_x

As presented in Figure 3, with the incorporation of Mo, both the room/high-temperature yield strengths and hardness of Ti₂₂Zr₂₅Hf_(35−x)Nb₁₈Mo_x increase.

Table 4. Yield strength and fracture strain of Mo_x.

Alloys	HV3	σ0.2-RT (MPa)	σ0.2-HT (MPa)	ε (%)
Mo10	346	1206	131	-
Mo15	375	1424	309	-
Mo20	412	1629	429	24

indicates that the room-temperature yield strength σ_0.2-RT of Mo20 is 423 MPa higher than that of Mo10, and the solid solution strengthening effect of Mo is substantially enhanced in a high-temperature environment (1000 °C), with the high-temperature yield strength σ_0.2-HT increasing from 131 MPa to 429 MPa. The hardness also increased from 346 HV to 412 HV with the increase in Mo content. In single-phase BCC solid solutions, solid solution strengthening is the main strengthening mechanism [32,33]. Owing to the substantial atomic size difference between Mo and other constituent elements, a substantial lattice distortion occurs, substantially enhancing the solid solution strengthening effect and elevating its yield strength. Furthermore, because the Mo element exhibits superior high-temperature performance, its high-temperature yield strength is substantially improved in high-temperature compression tests. However, with increasing Mo content, plasticity concurrently decreases. At room temperature, the fracture strain ε of Mo20 is 24% and the yield strength σ_0.2-RT is 1629 Mpa, which is consistent with the results of high δ and high VEC corresponding to the high solid solution strengthening effect and low plasticity in Table 2.

3.1.3. Room-Temperature Friction Performance of Mo_x

In Figure 4a–c, it is observed that the wear trajectories of the three compositions are comparatively uneven, which may be attributed to the uneven microstructure or defects in the as-cast material. Varying local wear rates result in wear grooves of different depths.

In Figure 5a,b, it is found that with the increase in the Mo element, both the wear depth and wear rate (calculated by Formula (8)) decrease. This is mainly because Mo can enhance the solid solution strengthening effect, and the lattice distortion produced by solid solution strengthening can hinder the movement of dislocations during the deformation process, thereby strengthening the alloy matrix and synergically enhancing the wear resistance [34]. In Figure 5c, it is noted that the coefficient of friction varies substantially during the initial stage of friction. This is attributed to the initial stage of friction, where the Si₃N₄ ceramic balls of the friction pair interact with the RHEAs surface, producing debris. The debris forms a third body between the frictional surfaces, elevating point contact friction and the coefficient of friction. With increasing friction time, the friction achieves a stable state [35]. However, the friction coefficient of Mo10 continues to vary significantly in the later stage of friction, primarily because numerous new deformation layers interact with the friction pair during the friction process, rendering the friction coefficient unstable. In addition, with increasing Mo content, the coefficient of friction decreases (from 0.686 to 0.664), primarily because increasing Mo content promotes the formation of the oxide layer. The continuous and dense oxide glaze layer serves a lubricating function, thereby lowering the coefficient of friction and minimizing wear [36].

$$W={\displaystyle \frac{V}{LF}}$$

(8)

In the formula, V is the wear volume, mm³; L is the friction length, m; and F is the load, N.

Figure 5. Mo_x: (a) friction depth, (b) wear rate, and (c) friction coefficient.

Figure 5. Mo_x: (a) friction depth, (b) wear rate, and (c) friction coefficient.

After room-temperature friction testing, many furrows parallel to the sliding direction are observed in the wear trajectories presented in Figure 6a–c, and a delamination phenomenon is apparent in the frictional trajectories, indicating that all compositions exhibit abrasive and adhesive wear mechanisms. During the friction testing, relatively hard wear fragments and free oxide particles engage with Si₃N₄ ceramic balls, inducing micro-cutting on the frictional surface, resulting in abrasive wear. In Figure 6a,b, it is observed that Mo10 and Mo15 exhibit more plastic deformation layers and fewer furrows and oxide layers. This is primarily attributed to the frictional surfaces of alloys with greater plasticity (Mo10 and Mo15) being susceptible to interacting with Si₃N₄ ceramic balls, inducing plastic deformation, resulting in adhesive wear. The primary mechanism for Mo10 and Mo15 is adhesive wear. In Figure 6c, it is observed that Mo20 exhibits more furrows and oxide layers with fewer deformation layers. This is attributed to the heat generated during friction testing. During the friction and wear testing, metal elements react with oxygen to form fine-grained oxide particles, which are subsequently compacted into oxide layers during the friction testing. With increasing Mo composition, the wear mechanism progressively transitions from adhesive wear to abrasive and oxidative wear. In Figure 5b, it is observed that the improvement in wear performance from Mo15 to Mo20 is comparatively modest. This is attributed not only to the comparatively high oxidative wear of Mo20, but also to the presence of the alloy’s plastic deformation film, which partially mitigates the direct plowing effect of hard particles on the alloy, thereby reducing abrasive wear [37].

3.2. The Microstructure and Properties of Mo20 (0.1Y/0.2Y)

3.2.1. Phase and Microstructure Analysis of Mo20 (0.1Y/0.2Y)

As presented in Figure 7, the incorporation of rare earth Y does not alter the phase structure of Mo20. However, with the incorporation of the Y element, the peak position of the XRD pattern shifts leftward. The lattice parameter of these components is as follows: Mo20 (d = 3.329), 0.1Y (d = 3.341), and 0.2Y (d = 3.356). Since the atomic radius of the Y element is larger than that of other refractory elements, when Y is dissolved in the BCC matrix, it increases the lattice constant, resulting in a leftward shift in the XRD pattern peak positions for Mo20.

Fine particles were observed in the SEM images of Figure 8a,b and were identified as Y₂O₃ particles by EDS in Figure 8f,g. With the incorporation of rare earth Y, the abundance of Y₂O₃ also increased, and some Y₂O₃ particles began to coarsen. Moreover, in Figure 8c,d, it was observed that the majority of Y₂O₃ was distributed within the interdendritic regions, primarily attributed to the exclusion of Y atoms from the main dendrite trunk of the BCC phase during the solidification process, resulting in their enrichment in the interdendritic region [38]. Y atoms exhibit a high affinity for oxygen [39], reacting with residual oxygen in the raw materials, thereby forming Y₂O₃ particles. In Figure 8e, a shrinkage defect in the 0.2Y composition was observed, which may be attributed to the incorporation of the Y element widening the solidification temperature range and resulting in the shrinkage defect in the later stages of solidification.

Figure 9 shows the EBSD spectra of Mo20, 0.1Y, and 0.2Y. With the incorporation of rare earth Y, the average grain size of Mo20 is reduced from 131.5 μm to 54.5 μm. The incorporation of rare earth Y facilitates heterogeneous nucleation conditions, thereby decreasing the grain size [40]. The effect of grain size variation on alloy properties was investigated using the Hall–Petch equation [41].

$${\sigma }_y={\sigma }_0+{\displaystyle \frac{K_{HP}}{\sqrt{D}}}$$

(9)

$$\mathsf{\Delta }{\sigma }_y=K_{HP}({\displaystyle \frac{1}{\sqrt{D_x}}}-{\displaystyle \frac{1}{\sqrt{D_0}}})$$

(10)

In the formula, σ_y represents the yield strength, σ₀ is the frictional force required for a single dislocation to move, and K_HP is the Hall–Petch constant, which measures the grain boundary’s hindrance to dislocation slip. The K_HP of high-entropy alloys is typically several hundred [42,43,44]. D is the average grain size, and Δσ_y represents the yield strength resulting from fine-grained strengthening. D₀ is the original grain size, and D_x is the refined grain size.

As shown in Figure 10a–c, the EPMA of the three components all present a dendritic structure, and the element distribution results are consistent with those in Table 2. However, it was found that the low-melting-point Ti is more evenly distributed than other elements. This might be due to the higher temperature of vacuum arc melting, and at high temperatures, the diffusion coefficient of Ti is higher than that of other elements, making its distribution more uniform. In addition, Ti has an atomic radius of only 146 pm and an electron negativity of 1.54, which are both closer to Nb and Mo; thus, Ti is more accepted by the BCC lattice since other Hume-Rothery factors are the same for (Nb,Mo)-Ti and (Nb,Mo)-Zr (crystal structure, valence electron concentration). Therefore, the actual tendency of Ti to be rejected by the prior BCC dendrites is milder despite its lowest melting temperature, leading to the observed homogeneous distribution.

3.2.2. Mechanical Performance of Mo20 (0.1Y/0.2Y)

As illustrated in Figure 11a,b, as well as the high-temperature true stress–strain curves, it was observed in Figure 11a that the addition of Y initially enhances both the yield strength and plasticity, followed by a subsequent decline. Firstly, this is attributed to the relatively large atomic radius of rare earth Y, which, through its solid solution in the BCC matrix, strengthens the solid solution strengthening effect. Secondly, rare earth elements facilitate nucleation sites for RHEAs, thereby refining the grain structure. As shown in

Table 5. Yield strength and fracture strain of Ti₂₂Zr₂₅Hf₁₅Nb₁₈Mo₂₀ (0.1Y/0.2Y).

Alloys	HV3	σ0.2-RT (MPa)	σ0.2-HT (MPa)	ε (%)
Mo20	412	1629	429	24
0.1Y	420	1693	419	32
0.2Y	429	1526	392	23.5

, this results in enhanced strength (1693 MPa) and plasticity (32%) for 0.1Y. Nevertheless, the reduction in room-temperature yield strength and plasticity at 0.2Y deviates from the calculation results of the Hall–Petch formula in (9) and (10) above, likely attributable to performance degradation caused by the shrinkage defect in Figure 8e. In the high-temperature compression curve in Figure 11b, it was observed that the yield strengths of 0.1Y and 0.2Y both exhibited a slight reduction. The primary reason is that the Y element refines the RHEAs grains. Under high-temperature conditions, alloys with high grain boundary density are susceptible to grain boundary sliding [45], which results in local stress concentration and accelerates material softening. As shown in Figure 11c, its Vickers hardness increases with the addition of Y, mainly due to fine-grained strengthening and solid solution strengthening.

3.2.3. Mo20 (0.1Y/0.2Y) Room-Temperature Friction and Wear Test

As illustrated in Figure 12, it is found that as Y increases, the friction trajectory is as uneven as that of different Mo components. Rare earth Y does not improve the wear unevenness.

In Figure 13a,b, it was observed that after adding rare earth Y, both the wear depth and wear rate of Mo20 exhibited a reduction. The synergistic effect of fine grain strengthening and solid solution strengthening increases the strength of Mo20. The coefficient of friction in Figure 13c progressively stabilizes as friction time increases. The average coefficients of friction at 0.1Y and 0.2Y are 0.664 and 0.658, respectively. The addition of rare earth Y has a relatively small influence on the coefficient of friction.

As illustrated in Figure 14a,b, pronounced furrows, delamination, and a significant oxide layer are observed in the microscopic morphology of the friction trajectories of 0.1Y and 0.2Y components. The primary wear mechanisms of these two alloys are abrasive wear and oxidation wear. The addition of rare earth Y induces grain refinement and solid solution strengthening, enhancing the strength of Mo20, and consequently improving its wear performance. Nevertheless, due to the minimal addition of Y, its effect on the wear mechanisms of oxidation wear and abrasive wear remains limited, preserving the original wear mechanisms.

4. Conclusions

In this research, three compositions of Ti₂₂Zr₂₅Hf_(35−x)Nb₁₈Mo_x were prepared. Among the refractory high-entropy alloys with Mo contents of 10, 15, and 20, mechanical property tests revealed that Ti₂₂Zr₂₅Hf₁₅Nb₁₈Mo₂₀ exhibits a superior comprehensive performance, though its plasticity is lower than that of the other two compositions. Its mechanical properties can be further improved by incorporating rare earth Y, with the optimal content of Y determined through experimentation.

(1)

With the increase in Mo element content, the solid solution strengthening effect of RHEAs enhances. Lattice distortion can impede dislocation movement, and its strength increases significantly in high-temperature environments. The room-temperature yield strength and high-temperature yield strength of Mo20 have increased by 35.1% and 227.5%, respectively, compared with Mo10. The high-temperature yield strength has increased significantly. However, its plasticity decreases, and the plasticity can be further improved by adding rare earth Y.

(2)

Incorporating an optimal amount of rare earth Y facilitates grain refinement and solid solution strengthening. With the incorporation of rare earth Y, the average grain size of Mo20 decreased from 131.5 μm to 54.5 μm. Moreover, the hardness and wear resistance of Mo20 also increase with the addition of Y.

(3)

With the addition of Y, it was found that shrinkage defects were discovered in 0.2Y, which deteriorated its mechanical properties. However, in 0.1Y, its mechanical properties were all improved. The breaking strain of 0.1Y was increased to 32%, the yield strength was raised to 1693 MPa, and the wear rate was reduced to 1.026 × 10⁻⁴ mm³.