TiZrNbMoW refractory high-entropy alloy coating, preparation method and application thereof

Abstract

The present application relates to the technical field of stainless steel surface protection, and specifically discloses a TiZrNbMoW refractory high-entropy alloy coating, a preparation method and application. The present application adopts a TiZr alloy target, a MoW alloy target and a single-element Nb target for multi-target magnetron co-sputtering to deposit a TiZrNbMoW refractory high-entropy alloy coating on a stainless steel substrate. The columnar crystals of the coating are densely arranged, and the coating has a single BCC solid solution phase typical of high-entropy alloys. The excellent mechanical properties and wear resistance of the refractory high-entropy alloy coating can achieve excellent surface protection of stainless steel. In addition, the element composition ratio and microstructure of the coating can be flexibly regulated by regulating the sputtering power of each target material, especially by regulating the content of Mo and W to further improve the mechanical properties and wear resistance of the coating, thereby optimizing the protective performance of the coating.

Description

Technical Field

[0001] This invention relates to the field of stainless steel surface protection technology, specifically to a TiZrNbMoW refractory high-entropy alloy coating, its preparation method, and its application. Background Technology

[0002] AISI 304 stainless steel is the most widely used austenitic stainless steel in industrial production and daily life due to its excellent corrosion resistance, hygiene and safety, ease of processing and formability, and wide-temperature range thermal stability. However, its inherent defects are becoming increasingly apparent in applications, such as low surface hardness and insufficient wear resistance, making it difficult to resist everyday mechanical damage. Furthermore, due to the high austenitic stability resulting from its high nickel and chromium content, AISI 304 stainless steel cannot achieve overall strengthening through phase transformation. In addition, while techniques such as solution strengthening, cold work hardening, and grain refinement improve strength, they also sacrifice the material's plasticity and corrosion resistance.

[0003] To overcome this bottleneck, surface engineering technology has become a key strategy and effective supplement for enhancing the surface properties of AISI 304 stainless steel, enabling surface repair without compromising the overall material shape and inherent properties. Among these, medium-to-high entropy alloys, with multiple elements as the main components, possess unique chemical and physical properties such as heterogeneity in local chemical composition, high-entropy thermodynamics, lattice distortion in structure, and hysteresis diffusion in kinetics. These properties result in superior mechanical properties, higher wear resistance, and excellent corrosion resistance, making them ideal for surface protective coatings and perfectly suited to the surface protection requirements of AISI 304 stainless steel.

[0004] For example, patent document CN 111074224 A discloses a corrosion-resistant high-entropy alloy nitride coating. This coating, (VAlTiCrMo)N, is deposited on a stainless steel substrate using a magnetron sputtering splicing target, achieving excellent high hardness and corrosion resistance. However, the splicing target used in this method is a composite target formed by vertically stacking V, Al, Ti, Cr, and Mo targets from top to bottom. This requires high sputtering power (1000 W–3000 W) and sputtering for 4–10 hours to complete the deposition of the high-entropy alloy nitride coating. This not only results in high energy consumption and low production efficiency but also significantly increases equipment wear and tear, leading to significantly higher maintenance costs and hindering long-term stable production. Patent document CN 113025953 A discloses a high-entropy alloy nitride composite coating, its preparation method, and its application. It utilizes a magnetron sputtering metal target and an AlCrWTiMo alloy target to deposit a metal transition layer and a (AlCrWTiMo)N coating on a stainless steel substrate. The deposition rate is fast, the efficiency is high, and it is easily industrialized. However, the magnetron sputtering target uses a high-entropy alloy palladium with a fixed composition, resulting in a fixed coating composition. This makes it difficult to further optimize the coating performance by adjusting the element ratios within the high-entropy alloy coating.

[0005] In summary, there is an urgent need to provide a novel high-entropy alloy coating, its preparation method, and its application, which can be rapidly deposited to achieve efficient production, while also allowing for easy control of the elemental composition ratio in the high-entropy alloy coating to further optimize the coating performance. Summary of the Invention

[0006] The purpose of this invention is to provide a novel high-entropy alloy coating, its preparation method, and its application, which enables rapid deposition and high-efficiency production while allowing for easy control of the elemental composition ratio in the high-entropy alloy coating to further optimize the coating performance.

[0007] In a first aspect, the present invention provides a TiZrNbMoW refractory high-entropy alloy coating, the coating being prepared by magnetron sputtering process, and the coating comprising, by atomic percentage: Ti: 10~27 at.%, Zr: 6~12 at.%, Nb: 10~22 at.%, Mo: 21~38 at.%, W: 18~37 at.%.

[0008] Optionally, the coating comprises, by atomic percentage: Ti: 10~17 at.%, Zr: 6~9 at.%, Nb: 10~16 at.%, Mo: 30~38 at.%, W: 28~37 at.%.

[0009] Optionally, the coating is a densely packed columnar crystal and the coating is a single body-centered cubic solid solution phase, the coating having a (110) texture.

[0010] Optionally, the body-centered cubic solid solution phase has a lattice constant of 3.23~3.30 Å and a grain size of 0.57~1.02 Å.

[0011] Optionally, the coating has a hardness of 12–19 GPa and an elastic modulus of 180–265 GPa.

[0012] Optionally, the coating has a coefficient of friction between 1.16 and 1.18, and a wear rate of 1.4 × 10⁻⁶. -5 mm 3 / (N·m) or less.

[0013] Optionally, the coating frictionally forms a self-lubricating phase MoO3 and / or WO3.

[0014] Secondly, the present invention provides a method for preparing the aforementioned TiZrNbMoW refractory high-entropy alloy coating. The method involves multi-target magnetron co-sputtering using a TiZr alloy target, a MoW alloy target, and a single-element Nb target to deposit the coating on a substrate. Argon is used as the working gas. The DC sputtering power of the TiZr alloy target is 100-300 W, the DC sputtering power of the MoW alloy target is 100-500 W, the DC sputtering power of the single-element Nb target is 50-150 W, the RF bias power of the substrate is 30-70 W, the working flow rate of the argon gas is 30-50 sccm, and the deposition time is 30-150 min.

[0015] Optionally, the DC sputtering power of the TiZr alloy target is 150~250 W, the DC sputtering power of the MoW alloy target is 200~450 W, the DC sputtering power of the elemental Nb target is 80~120 W, the RF bias power of the substrate is 45~55 W, the working flow rate of the argon gas is 35~45 sccm, and the deposition time is 40~90 min.

[0016] Thirdly, the application of the aforementioned TiZrNbMoW refractory high-entropy alloy coating in the field of substrate surface protection is provided, wherein the substrate includes stainless steel.

[0017] In summary, the present invention has the following beneficial effects:

[0018] 1. This invention provides a novel TiZrNbMoW refractory high-entropy alloy (RHEAs) coating, its preparation method, and its applications. It enables rapid deposition and high-efficiency production, while also allowing for easy control of the elemental composition ratio in the RHEA coating to further optimize coating performance. Specifically, by independently controlling the sputtering power of the MoW alloy target, the atomic percentages of Mo and W in the RHEA coating are precisely controlled during the deposition process, thereby obtaining a series of coatings with different elemental ratios. Optimizing the elemental composition ratio and microstructure of the coating improves its mechanical and tribological properties.

[0019] 2. This invention provides a novel TiZrNbMoW refractory high-entropy alloy coating, its preparation method, and its application. The prepared coating exhibits regular columnar crystal growth characteristics with dense arrangement. The synergistic addition of Mo and W promotes the formation of a single body-centered cubic (BCC) solid solution phase in the alloy. The introduction of high shear modulus Mo and W elements promotes lattice distortion, thereby increasing the lattice frictional stress and yield strength in the alloy. The coating hardness increases from 8 GPa to 18.4 GPa with increasing MoW content, and the elastic modulus increases from 126.8 GPa to 260.1 GPa. The wear resistance of the coating is significantly improved, with the wear rate decreasing from 3.07 × 10⁻⁶. -5 mm 3 / (N•m) decreased to 0.86×10 -6 mm 3 / (N•m). Attached Figure Description

[0020] To more clearly illustrate the technical solutions in the embodiments of the present invention, the accompanying drawings used in the description of the embodiments will be briefly introduced below. Obviously, the accompanying drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0021] Figure 1 This is a schematic diagram of the multi-target co-deposition magnetron sputtering system of Embodiment 1 of the present invention;

[0022] Figure 2 These are (ae) cross-sectional SEM images of the sample coatings of Examples 1-4 and Comparative Example 1 of the present invention; (fj) surface SEM images; (ko) EDS elemental distribution of the surface;

[0023] Figure 3 These are XRD patterns of the sample coatings from Examples 1-4 and Comparative Example 1 of the present invention;

[0024] Figure 4 These are analytical diagrams of the lattice constant (a) and grain size (D) of the sample coatings of Examples 1-4 and Comparative Example 1 of the present invention;

[0025] Figure 5 The following are (a) nanoindentation load-displacement curves and (b) curves showing the changes in nanohardness and elastic modulus with indentation depth for the substrate AISI 304 of this invention, the sample coatings of Examples 1-4 and Comparative Example 1.

[0026] Figure 6 The following are the relationships between (a) hardness and elastic modulus of the substrate AISI 304 of this invention, and the sample coatings of Examples 1-4 and Comparative Example 1, and (b) H / E and H content. 3 / E 2 The ratio;

[0027] Figure 7 The following are (a) curves showing the coefficient of friction as a function of distance for the sample coatings of the AISI 304 substrate of the present invention, Examples 1-4 and Comparative Example 1; and (b) coefficient of friction values.

[0028] Figure 8 (af) SEM images of the worn surfaces of the sample coatings of the AISI 304 substrate of the present invention, Examples 1-4 and Comparative Example 1; (gl) elemental distribution of the wear scratch region;

[0029] Figure 9 This is the wear track profile of the wear surface of the sample coating of the AISI 304 substrate of the present invention, Examples 1-4 and Comparative Example 1. Detailed Implementation

[0030] This invention provides a TiZrNbMoW refractory high-entropy alloy coating, its preparation method, and its application. To enable those skilled in the art to better understand the technical solutions of this invention, the technical solutions in the embodiments of this invention will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some embodiments of this invention, and not all embodiments.

[0031] In some embodiments of the present invention, a TiZrNbMoW refractory high-entropy alloy coating is provided. The coating is prepared by magnetron sputtering and comprises, by atomic percentage: Ti: 10~27 at.%, Zr: 6~12 at.%, Nb: 10~22 at.%, Mo: 21~38 at.%, W: 18~37 at.%; preferably, Ti: 10~17 at.%, Zr: 6~9 at.%, Nb: 10~16 at.%, Mo: 30~38 at.%, W: 28~37 at.%; more preferably, Ti: 10~13.5 at.%, Zr: 6~7.7 at.%, Nb: 10~11.8 at.%, Mo: 34.5~37.5 at.%, W: 33.0~36.5 at.%.

[0032] In some embodiments of the present invention, the coating is a densely packed columnar crystal and the coating is a single body-centered cubic solid solution phase, the coating having a (110) texture.

[0033] In some embodiments of the present invention, the body-centered cubic solid solution phase has a lattice constant of 3.23~3.30 Å, a grain size of 0.57~1.02 Å, and an interplanar spacing of 2.29~2.33 Å; preferably, the lattice constant is 3.23~3.24 Å, the grain size is 0.57~0.66 Å, and the interplanar spacing is 2.29~2.30 Å.

[0034] In some embodiments of the present invention, the coating has a hardness of 12-19 GPa and an elastic modulus of 180-265 GPa; preferably, the coating has a hardness of 16-19 GPa and an elastic modulus of 225-265 GPa; more preferably, the coating has a hardness of 18-19 GPa and an elastic modulus of 250-265 GPa.

[0035] In some embodiments of the present invention, the total solid solution strengthening Δσ of the coating SS In the middle, modulus mismatch strengthens Δσ G The contribution rate is 66% to 89%.

[0036] In some embodiments of the present invention, the coefficient of friction of the coating is between 1.16 and 1.18, and the wear volume is 0.43~0.54×10⁻⁶. -4 mm 3 The wear rate is 1.4 × 10⁻⁶. -5 mm 3 / (N•m) or less; preferably, the wear rate is 0.86 ~ 1.07 × 10 -5 mm 3 / (N•m).

[0037] In some embodiments of the present invention, the coating frictionally forms a self-lubricating phase MoO3 and / or WO3.

[0038] In some embodiments of the present invention, the present invention provides a method for preparing the aforementioned TiZrNbMoW refractory high-entropy alloy coating. The method involves multi-target magnetron co-sputtering using a TiZr alloy target, a MoW alloy target, and a single-element Nb target to deposit the coating on a substrate. Argon is used as the working gas. The DC sputtering power of the TiZr alloy target is 100-300 W, the DC sputtering power of the MoW alloy target is 100-500 W, the DC sputtering power of the single-element Nb target is 50-150 W, the RF bias power of the substrate is 30-70 W, the working flow rate of the argon gas is 30-50 sccm, and the deposition time is 30-150 min.

[0039] In some embodiments of the present invention, the DC sputtering power of the TiZr alloy target is 150~250 W, the DC sputtering power of the MoW alloy target is 200~450 W, the DC sputtering power of the elemental Nb target is 80~120 W, the RF bias power of the substrate is 45~55 W, the working flow rate of the argon gas is 35~45 sccm, and the deposition time is 40~90 min.

[0040] Example 1: This example provides a TiZrNbMoW refractory high-entropy alloy coating, its preparation method, and its application. A multi-target co-deposition magnetron sputtering system (M230, China Weikai Technology) is used, employing high-purity TiZr alloy targets (purity ≥99.95%), MoW alloy targets (purity ≥99.95%), and elemental Nb targets (purity ≥99.9%) with equal atomic ratios for co-sputtering. The target-substrate distance is fixed at 15 cm. A schematic diagram of the multi-target co-deposition magnetron sputtering system is shown below. Figure 1 As shown. Before deposition, all substrates were ultrasonically cleaned sequentially in acetone and ethanol for 30 min each; simultaneously, Ar... + Ion pre-sputtering of the target material was performed for 30 min to reduce surface contamination and oxide residue. The substrate was then fed into the sample stage of the deposition chamber via a robotic arm, and the chamber was evacuated to a base vacuum of 5.0 × 10⁻⁶. -6 Torr (approximately 6.7 × 10⁻⁶) -4High-purity argon gas (purity ≥99.999%) was introduced into the vacuum deposition chamber via a precision gas path system. Plasma was generated by applying a bias voltage, and residual contaminants on the substrate surface were removed by plasma pre-cleaning. Subsequently, coating deposition was performed at room temperature (25 °C). The RF substrate bias power was set to 50 W, and the sample stage rotated the substrate at a constant speed of 50 rpm. A TiZrNbMoW sample coating, named MoW100, was prepared using a MoW sputtering power of 100 W. The specific magnetron sputtering deposition parameters are shown in Table 1. It should be noted that single-crystal Si(111) wafers and mirror-polished AISI 304 stainless steel substrates were used for sample coating preparation in the same magnetron sputtering process. The sample coating deposited on the Si substrate was used for composition and microstructure characterization, while the coating deposited on the AISI 304 substrate was used for mechanical and tribological property testing.

[0041] Examples 2-4: The difference between Examples 2-4 and Example 1 is that the TiZrNbMoW sample coatings in Examples 2, 3, and 4 were prepared with MoW sputtering powers of 200 W, 300 W, and 400 W, respectively, and were named MoW200, MoW300, and MoW400, respectively.

[0042] Comparative Example 1

[0043] The difference between Comparative Example 1 and Example 1 is that Comparative Example 1 prepared a TiZrNb sample coating with a MoW sputtering power of 0 W, which was named MoW0.

[0044] For the specific target power parameters of the coatings prepared in Examples 1-4 and Comparative Example 1, please refer to Table 1. The coating thickness was controlled to be around 1.2 μm by adjusting the deposition time.

[0045] Table 1. Deposition parameters of the coating

[0046] I. Microstructure testing and characterization

[0047] Field emission scanning electron microscopy (FESEM) was used to observe the surface and cross-sectional morphology of the sample coatings on the Si(111) wafer substrates of Examples 1-4 and Comparative Example 1 using secondary electron signal imaging. Simultaneously, elemental distribution and chemical composition information were obtained using a matching X-ray energy dispersive spectroscopy (EDS). The test results are as follows: Figure 2 As shown in (ae), all samples (MoW0 to MoW400) exhibit a typical columnar crystal growth structure with dense grains and high orientation consistency. Compared with the undoped Mo / W coating, the introduction of Mo and W transforms the surface morphology of the coating from a loose porous structure to a dense layered structure (see [reference]). Figure 2 (f–j)). Figure 2 The EDS elemental distribution results of (ko) visually present the uniform distribution of Ti, Zr, Nb, Mo, and W elements in the coating, with no obvious elemental segregation or phase separation observed. This is consistent with the core characteristic of high-entropy alloys: "high mixing entropy drives uniform elemental solid solution." The specific EDS test results and stoichiometry of the sample coatings in Examples 1-4 and Comparative Example 1 are shown in Table 2. As can be seen from Table 2, with the increase of MoW sputtering power, the atomic percentages of Mo and W elements in the coating increase, while the atomic percentages of Ti, Zr, and Nb decrease. This compositional evolution stems from the sputtering flux competition during the multi-target co-sputtering process: with fixed power for the TiZr and Nb targets, increasing the MoW target power enhances the sputtering yield of Mo and W, increasing their content and diluting Ti, Zr, and Nb; simultaneously, at high power, the deposited Mo and W atoms have higher kinetic energy, triggering slight resputtering. Lighter elements (Ti and Zr) are preferentially sputtered due to their lower sputtering thresholds, further reducing their content in the coating.

[0048] Table 2. EDS results (atomic ratio) and stoichiometry of the coatings

[0049]

[0050] To confirm the true structure of the coatings, the crystal structure of the sample coatings from Examples 1-4 and Comparative Example 1 was analyzed using grazing incidence X-ray diffraction (GIXRD). The incident angle for GIXRD testing was 2°, the scanning range was 2θ = 30°–90°, and the scanning rate was 5°·min. -1 The operating voltage is 40 kV and the current is 40 mA. Crystal structure information of the material is obtained by analyzing the position and intensity of diffraction peaks. The test results are as follows: Figure 3As shown, all sample coatings exhibited only clear body-centered cubic (BCC) diffraction peaks within the detection range of 30° to 90°, corresponding to the (110), (200), (211), and (220) crystal planes, and no characteristic peaks of other crystal structures were observed. This indicates that within the composition range of the present invention, the addition of Mo and W together promotes the formation of a single BCC solid solution phase in the alloy, exhibiting typical high-entropy alloy characteristics. Meanwhile, the intensity of the (110) diffraction peak in the MoW200, MoW300, and MoW400 coatings was significantly higher than that in MoW0 and MoW100, and it dominated all crystal plane diffraction peaks, indicating that the coating formed a strong (110) preferred orientation. This preferred orientation gradually increased with the increase of MoW target power, reflecting the gradual evolution of the coating microstructure: on the one hand, the significant (110) texture corresponds to optimized columnar crystal growth and higher coating density, because the (110) plane is the closest packing plane in the BCC structure, and its preferred orientation can reduce the interface energy and promote close atomic packing; on the other hand, this strong (110) texture is directly related to the increase in coating hardness, because in BCC metals and alloys, <110> The crystal orientation has a high atomic line density, is the dominant slip system direction, and has strong interatomic bonding, which together enhances the theoretical shear strength and elastic modulus. Based on Bragg's law, the grain size (D), lattice constant (a), interplanar spacing (d), and 2θ angle corresponding to the (110) crystal plane of the MoW0~MoW400 coatings were calculated, and the results are listed in Table 3. All grains are nanoscale, and the lattice constant gradually decreases from 3.40 Å in MoW0 to 3.23 Å ​​in MoW400, reflecting lattice shrinkage in the coating. The main reason for this is that the Ti and Zr atoms with larger atomic radii in the matrix are replaced by Mo and W atoms with smaller radii. Figure 4 The lattice constant (a) and grain size (D) of the TiZrNbMoW coating are shown, confirming that both decrease monotonically with increasing MoW target sputtering power, indicating a regular evolution of the microstructure with deposition energy. Grain refinement and the persistent atomic size mismatch between Mo / W and the matrix elements (Ti, Zr, Nb) jointly induce significant lattice distortion, which is the key solid solution strengthening mechanism for improving the mechanical properties of the coating.

[0051] Table 3. Lattice constant and grain size of the coating

[0052]

[0053] II. Mechanical property testing

[0054] The mechanical properties of the TiZrNbMoW coatings on the AISI 304 substrate and in Examples 1-4 and Comparative Example 1 were systematically characterized using the nanoindentation continuous stiffness method (CSM). The hardness (H) and elastic modulus (E) of the TiZrNbMoW were obtained using a nanoindenter (Keysight Technologies G200) equipped with a standard Berkovich indenter in CSM mode. The hardness (H) and elastic modulus (E) of the coating were obtained by analyzing the load-displacement curves during loading and unloading. To reduce random errors, six different locations on the coating surface were selected for testing, with a maximum indentation depth of 200 nm. To minimize the influence of the substrate on the nanoindentation test, the H and E values ​​at an indentation depth of approximately 1 / 10 of the coating thickness (120 nm) were averaged.

[0055] Test results are as follows Figure 5 As shown, Figure 5 (a) Nanoindentation load-displacement curves of the AISI 304 substrate, samples from Examples 1-4, and Comparative Example 1 are presented. The curves of all sample coatings shown are smooth and continuous, without obvious abrupt changes or steps, indicating that no brittle failure behaviors such as coating cracking or large-scale delamination occurred during the indentation process, demonstrating that the coating has good bonding strength with the substrate and the material's own plastic deformation capability. Figure 5 (b) The hardness (H) of the coatings in the AISI 304 substrate, Examples 1-4, and Comparative Example 1 is shown as a function of indentation depth. As the indentation depth increases, the hardness curves gradually tend towards a relatively stable plateau. The average values ​​within these plateau regions can reliably characterize the intrinsic (bulk) properties of the coating. Therefore, this invention selects values ​​measured at an indentation depth of approximately 1 / 10 (120 nm) of the coating thickness for analysis and comparison.

[0056] Figure 6 (a) This demonstrates the relationship between nanohardness (H) and elastic modulus (E) of the coatings on the AISI 304 substrate, Examples 1-4, and Comparative Example 1, as a function of MoW content. Figure 6 As shown in (a) and Table 4, the nanoindentation hardness (H) of the TiZrNbMoW coatings in Examples 1-4 is significantly higher than that of the AISI 304 substrate (4.0±0.3 GPa). Specifically, the hardness H increases from 8.0±0.2 GPa in MoW0 to 18.5±0.5 GPa in MoW300, and stabilizes at 18.4±0.3 GPa in MoW400, with an increase of 98.0%~357.2%. The elastic modulus E increases from 126.8±1.7 GPa to 260.1±4.4 GPa, which is significantly higher than that of the AISI 304 substrate.

[0057] Table 4. Hardness, elastic modulus, H / E and H of the coating 3 / E 2

[0058]

[0059] It is worth noting that the H and E values ​​of MoW300 and MoW400 are basically the same. Combining Tables 2 and 4, it can be seen that when the Mo / W content exceeds about 35 at.%, the lattice distortion caused by the substitution of Ti / Zr / Nb matrix atoms by small-radius Mo / W atoms has a decreasing effect on the strengthening of mechanical properties.

[0060] In addition, tribological properties are closely related to the hardness / modulus ratio (H / E) and the resistance to plastic deformation (H³ / E²). The higher the value of this parameter, the better the wear resistance of the material. Figure 6 (b) shows that the H / E (0.073) and H³ / E² (0.010 GPa) of MoW300 both reached their peak values, indicating that MoW300 has the best tribological properties and providing a theoretical basis for subsequent tribological test results.

[0061] To quantitatively reveal the strengthening mechanism behind the hardness increase, a decoupled analysis of the solid solution strengthening contribution was performed, including: based on the optimized Labusch model, namely the Toda–Caraballo model. The lattice deformation of the alloy originates from the differences in atomic size and shear modulus (G) among the constituent elements. This is achieved through atomic size mismatch (Δσ). r ) and shear modulus mismatch (Δσ) G The contribution of solid solution strengthening can be quantitatively assessed using the following formula:

[0062]

[0063] Where A is a dimensionless constant, G is the shear modulus of the alloy, δ is the estimated average difference of atomic size or shear modulus, and c is the solute concentration. For BCC solid solution HEAs, the atomic size (δ... rij ) and atomic modulus (δ Gij The average difference of ) is defined by the formula and definition respectively:

[0064]

[0065]

[0066] in and These represent the atomic size difference and shear modulus difference between atoms i and j, respectively. Meanwhile, δi represents the difference in atomic size and shear modulus between different atomic pairs. rij and δ Gij The calculated values ​​are shown in Tables 5 and 6.

[0067] Shear modulus (G) is determined by the formula G=E. r / [2(1+γ)], calculated using the elastic modulus measured by nanoindentation, where γ=0.36, E r It represents the elastic modulus.

[0068] Table 5. Atomic size difference (δ) between i and j atom pairs in the coating rij )

[0069]

[0070] Table 6. Difference in atomic shear modulus (δ) between atomic pairs i and j in the coating Gij )

[0071]

[0072] Table 7 summarizes the effects of atomic size mismatch and shear modulus mismatch in TiZrNbMoW coatings, and provides the key mismatch parameters (δ) for each composition. r δ G δ SS The calculated value of ). For the undoped MoW0 sample, the total solid solution strengthening (Δσ) SS The strength reached 66.8 MPa, of which atomic size mismatch contributed (Δσ) r The contribution of 44.3 MPa to the shear modulus mismatch is approximately Δσ. G The shear modulus is nearly twice that of MoW100 (22.5 MPa). The introduction of Mo / W makes shear modulus mismatch the dominant strengthening mechanism: in MoW100, Δσ SS It surged to 757.1 MPa, where Δσ G (673.3 MPa) accounted for 89%. At higher target power, the Δσ of MoW200... SS 1210.8 MPa (Δσ) G The peak pressure was 898.5 MPa, accounting for 74%; while MoW300 reached a peak pressure of 1529.8 MPa (Δσ). G (1064.8 MPa, accounting for 70%), while Δσ r (465.0 MPa) increases due to lattice distortion. For MoW400 (Mo / W content > 35 at.%), Δσ SS It reached 1715.7 MPa, but Δσ G The growth rate of [specific modulus] slowed down (1128.7 MPa, accounting for 66%), exhibiting a saturation effect. Overall, the modulus mismatch strengthening (Δσ) driven by the high shear modulus of Mo and W [is evident]. GThe modulus mismatch remains the primary strengthening mechanism, contributing 66%–89%, consistent with the results of hardness and modulus tests (see Table 4). This shift from size mismatch-dominated to modulus mismatch-dominated mechanisms explains the hardness enhancement observed in the experiments, while modulus mismatch strengthening tends to saturate at high Mo / W concentrations.

[0073] Table 7. Atomic size mismatch and modulus mismatch of coatings

[0074]

[0075] III. Tribological property testing

[0076] The tribological properties of the AISI304 substrate and the TiZrNbMoW alloy coatings of Examples 1-4 and Comparative Example 1 were evaluated using a disc-ball tribometer (TRB3, Anton Paar) in reciprocating mode at room temperature. The friction pair, amplitude, total sliding distance, normal load, and constant frequency were Al2O3 (Φ=6 mm), 2.5 mm, 20 m, 0.5 N, and 0.5 Hz, respectively. After testing, the wear track morphology of the coatings was characterized using scanning electron microscopy (SEM), the elemental distribution of the worn surface was obtained using energy dispersive X-ray spectroscopy (EDS), and an electrical signal was generated by linear contact sliding of a diamond probe on the sample surface using a 3D surface profilometer (DektakXT, Bruker) to reflect the surface profile curve of the coating. The cross-section and three-dimensional profile (3D) of the wear track were obtained by scanning. The cross-sectional data were imported into Origin 2021 (9.8) software for integral calculation of the wear volume, and the wear rate was calculated according to the following formula:

[0077]

[0078] Where W is the wear rate (mm) 3 / (N·m)), V is the wear volume (mm²) 3 F is the applied load (N), and S is the total sliding distance (m).

[0079] Figure 7(a) Curves showing the coefficient of friction (COF) as a function of sliding distance for the substrate AISI304, the TiZrNbMoW coatings of Examples 1-4, and Comparative Example 1. All coatings showed a rapid increase in COF during the initial running-in phase; however, AISI304, MoW0, and MoW100 exhibited significant and drastic fluctuations in COF, caused by the rupture of the surface oxide layer during wear. This resulted in a change in the contact pattern from "coating-dual surface" to a direct "substrate-dual surface" contact, indicating coating failure due to insufficient hardness and suggesting poor wear resistance. In contrast, MoW200, MoW300, and MoW400 did not exhibit such fluctuations after surface stabilization. Their COF, after a rapid initial increase, remained stable at around 1.2 throughout the process, indicating the formation of a stable and reliable tribological system.

[0080] Figure 7 (b) shows the steady-state friction coefficient. The results indicate that the friction coefficient of MoW200–MoW400 is stable at 1.16–1.18. However, the lower average friction coefficients of AISI304 and MoW0 / MoW100 are not meaningful due to the premature failure of the coating during the wear process.

[0081] Figure 8 The results of tribological tests on AISI 304 substrate and TiZrNbMoW coatings with different MoW contents are presented, including the wear surface morphology and the distribution of O and Fe elements in the wear track area. The results show that the AISI 304 substrate, MoW0, and MoW100 coating samples exhibit deep and continuous ploughing grooves, large-area delamination, severe plastic deformation, and large-scale oxygen-rich wear areas (see...). Figure 8 (a–c) and (g–i) confirm that the wear mechanism is mainly oxidation-abrasive wear; elemental surface scanning further shows that Ti, Zr, Nb, Mo, and W elements are lost in the wear track area, while a strong Fe signal appears, indicating that the substrate is exposed and the coating integrity is poor. The presence of hard iron oxides at the wear interface of MoW0 and MoW100 coated samples dominates the friction process and induces severe wear, but their wear degree is still lower than that of AISI304. In contrast, the wear characteristics of MoW200, MoW300, and MoW400 are shallower and smoother (see Figure 8(d–f) and (j–l)). No Fe signal was detected within the wear tracks, confirming that the coating remained intact and no substrate exposure occurred. The oxygen signal was enhanced at the edges of the wear tracks and weakened in the central region, indicating that oxygen-containing wear products were generated in situ during friction. Driven by shear stress, the Mo / W-rich oxygen-containing lubricating layer migrated towards the edges of the wear tracks, placing the central region of the wear tracks in a dynamic equilibrium between lubricant layer formation and removal. The MoO3 / WO3 oxide generated by the reaction of the coating surface with oxygen / moisture in the environment acts as a solid lubricant to reduce wear. This transformation changed the friction interface from being dominated by hard iron oxides to being dominated by self-lubricating oxides. The elemental surface scan results showed that Ti, Zr, Nb, Mo, and W elements were almost not lost, confirming that the high Mo / W content (see Table 2) and hardness exceeding 16 GPa (see Table 4) provided reliable protection for the substrate.

[0082] Figure 9 The wear track profiles of the AISI304 substrate and TiZrNbMoW coatings with different MoW contents are presented. The wear depth decreases in the order of AISI304 > MoW0 > MoW100 > MoW200 > MoW300 > MoW400. The wear track width and maximum wear track depth of MoW200, MoW300, and MoW400 are at a low level and similar in value, significantly lower than those of AISI304, MoW0, and MoW100. This indicates that the wear rate decreases sharply between MoW100 and MoW200, which means that the wear mechanism at the contact interface changes. To avoid catastrophic failure and ensure a long service life, the Mo and W contents in the TiZrNbMoW coating system must both exceed 20 at.% (see Table 2) to provide sufficient mechanical stability under long-term frictional loads.

[0083] Figure 9 The wear volume and wear rate of the uncoated AISI304 substrate and the TiZrNbMoW coating were calculated by three-dimensional wear track profile analysis, and the results are summarized in Table 8. The maximum wear depth of AISI304 was approximately 3.86 µm, which exceeded the coating thickness. The maximum wear depth of MoW0, MoW100, MoW200, MoW300, and MoW400 decreased sequentially, to approximately 1.95, 1.34, 0.26, 0.24, and 0.30 µm, respectively. The wear track width of all TiZrNbMoW coated samples was significantly smaller than that of the AISI304 substrate. Under the same load and sliding conditions, the wear volume and wear rate decreased continuously with the increase of MoW target power. Among them, MoW300 had the shallowest and narrowest wear morphology (see Table 8). Figure 9 (e) The wear volume is the smallest, and the wear rate is the lowest, only (8.6±1.9)×10. -8 mm³ / (N·m) confirms that it has the best wear resistance.

[0084] Table 8. Average coefficient of friction, wear volume, and wear rate of the coating

[0085]

[0086] Hardness is the dominant factor determining the wear resistance of a material. In this invention, the incorporation of Mo and W effectively improves the coating hardness, inhibits abrasive plowing and surface delamination, thereby reducing wear and improving frictional stability. MoW300 exhibits the best tribological properties and reliability. Its highest hardness measured by nanoindentation reaches 18.5 ± 0.5 GPa (see Table 4), which is highly consistent with the wear behavior, further verifying the positive correlation between hardness and wear resistance.

[0087] In summary, the improved tribological properties of the TiZrNbMoW coating provided by this invention stem from a synergistic strengthening mechanism, mainly including the following four points:

[0088] (1) Hardness was improved by adding Mo / W (see Table 4). The coating hardness increased from 8 GPa in MoW0 to 18.5 GPa in MoW300, which reduced the indentation depth of the abrasive grains and changed the wear morphology from deep furrows in AISI 304 / MoW0 / MoW100 to shallow scratches in MoW200 / MoW300 / MoW400. This result is based on wear depth data (see Table 8) and SEM observation (see Table 9). Figure 8 and 9 This has been confirmed;

[0089] (2) Microstructure optimization: The coating has dense columnar grains (see Figure 2 (SEM morphology) and strong (110) texture (see SEM morphology) and strong (110) texture (see SEM morphology) Figure 3 (GIXRD structure), which improves grain boundary bonding and shear resistance, thereby suppressing fatigue-induced delamination under cyclic loading (see...). Figure 8 (af));

[0090] (3) The Mo / W-rich oxygen-containing lubricating layer formed during the friction process can effectively avoid direct contact at the friction interface, significantly reducing the coefficient of friction and wear rate. At the same time, MoW300 has the highest H³ / E² (0.010±0.003 GPa) and H / E (0.073±0.004) ratios (see Table 4), and the improvement of these parameters is beneficial to enhancing wear resistance;

[0091] (4) EDS energy dispersive spectroscopy analysis after wear (see Figure 8 (jl) No elemental segregation or loss was found, confirming that the coating has good compositional stability and structural integrity, and can effectively protect the substrate during friction.

[0092] This multi-scale mechanism combines atomic-scale strengthening effects with microstructure control, enabling the coating to achieve excellent macroscopic tribological properties.

[0093] This invention demonstrates that co-doping with Mo and W elements in a TiZrNb coating and controlling the sputtering power of the MoW alloy target can significantly improve the hardness and wear resistance of the stainless steel substrate.

[0094] First, increasing the sputtering power from 100 W to 400 W can increase the Mo and W contents from 21.3 at.% and 18.7 at.% to 37.3 at.% and 36.3 at.%, respectively, and synergistically produce two key effects:

[0095] (1) It forms a dense columnar microstructure and is stable as a single body-centered cubic (BCC) solid solution phase;

[0096] (2) Lattice distortion caused by high shear modulus of Mo / W.

[0097] Ultimately, the coating hardness increased from 8.0±0.2 GPa for MoW0 to 18.5±0.5 GPa for MoW300, a 4.6-fold increase compared to untreated stainless steel.

[0098] Furthermore, as the Mo / W content increases, the dominant strengthening mechanism changes from atomic size mismatch to modulus mismatch, a shift confirmed by calculations using the Toda-Caraballo model.

[0099] This specific embodiment is merely an explanation of the present invention and is not intended to limit the invention. After reading this specification, those skilled in the art can make modifications to this embodiment without contributing any inventive step, but such modifications are protected by patent law as long as they are within the scope of the claims of the present invention.

Claims

1. A TiZrNbMoW refractory high-entropy alloy coating, characterized in that, The coating is prepared by magnetron sputtering and comprises, by atomic percentage: Ti: 10~27 at.%, Zr: 6~12 at.%, Nb: 10~22 at.%, Mo: 21~38 at.%, W: 18~37 at.%.

2. The TiZrNbMoW refractory high-entropy alloy coating according to claim 1, characterized in that, The coating comprises, by atomic percentage: Ti: 10~17 at.%, Zr: 6~9 at.%, Nb: 10~16 at.%, Mo: 30~38 at.%, W: 28~37 at.%.

3. The TiZrNbMoW refractory high-entropy alloy coating according to claim 1 or 2, characterized in that, The coating is composed of densely arranged columnar crystals and is a single body-centered cubic solid solution phase, and the coating has a (110) texture.

4. The TiZrNbMoW refractory high-entropy alloy coating according to claim 3, characterized in that, The body-centered cubic solid solution phase has a lattice constant of 3.23–3.30 Å and a grain size of 0.57–1.02 Å.

5. The TiZrNbMoW refractory high-entropy alloy coating according to claim 1 or 2, characterized in that, The coating has a hardness of 12–19 GPa and an elastic modulus of 180–265 GPa.

6. The TiZrNbMoW refractory high-entropy alloy coating according to claim 1 or 2, characterized in that, The coating has a coefficient of friction between 1.16 and 1.18, and a wear rate of 1.4 × 10⁻⁶. -5 mm 3 / (N·m) or less.

7. The TiZrNbMoW refractory high-entropy alloy coating according to claim 1 or 2, characterized in that, The coating friction forms a self-lubricating phase MoO3 and / or WO3.

8. A method for preparing a TiZrNbMoW refractory high-entropy alloy coating according to any one of claims 1 to 7, characterized in that, The coating was deposited on a substrate using multi-target magnetron co-sputtering with TiZr alloy targets, MoW alloy targets, and elemental Nb targets. Argon was used as the working gas. The DC sputtering power of the TiZr alloy target was 100~300 W, the DC sputtering power of the MoW alloy target was 100~500 W, the DC sputtering power of the elemental Nb target was 50~150 W, the RF bias power of the substrate was 30~70 W, the working flow rate of the argon gas was 30~50 sccm, and the deposition time was 30~150 min.

9. The method for preparing the TiZrNbMoW refractory high-entropy alloy coating according to claim 8, characterized in that, The DC sputtering power of the TiZr alloy target is 150~250 W, the DC sputtering power of the MoW alloy target is 200~450 W, the DC sputtering power of the elemental Nb target is 80~120 W, the RF bias power of the substrate is 45~55 W, the working flow rate of the argon gas is 35~45 sccm, and the deposition time is 40~90 min.

10. The application of a TiZrNbMoW refractory high-entropy alloy coating according to any one of claims 1 to 7 in the field of substrate surface protection, wherein the substrate comprises stainless steel.

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