The Additions of V and Cu on the Microstructure and Mechanical Properties of Mo-N Coatings

Abstract

Due to the excellent lubricity of V₂O₅ and soft metals, V and Cu have been added to Mo-N based coatings to further improve the tribological properties. In this study, the Mo-V-Cu-N coatings were deposited by high power impulse magnetron sputtering (HIPIMS). The effects of V and Cu on the microstructure and mechanical properties of Mo-N coatings were investigated. With increasing V/Cu content ratio, the deposition rate decreased from 15.4 to 6.5 nm/min, and the microstructure transformed from a featureless structure into a dense columnar structure. At low Cu contents, less than 6.5 at.%, the Mo-V-Cu-N coatings exhibited a single solid solution phase of c-Mo₂(V)N. When the Cu content reached 29.7 at.%, the Mo₄₅V₁Cu₃₀N₂₄ coating showed the lowest surface roughness of 2.0 nm, and the coating changed into a double-phase of c-Mo₂(V)N and c-Cu. The adhesion strength gradually increased from 32.2 to 87.8 N with an increasing V/Cu content ratio. Due to the microstructure densification, a maximum hardness of 27.3 GPa was achieved for the Mo₄₆V₁₅Cu₁N₃₈ coating, which was accompanied by a high compressive residual stress.

1. Introduction

Due to their superior physical, chemical, and mechanical properties, molybdenum nitrides have been widely studied and applied in recent decades, such as diffusion barriers for microelectronic devices [1], anode materials for Li-ion batteries [2], hard protective coatings for cutting tools and molds [3], and self-lubrication coatings for friction parts [4]. The Mo-N coatings have been successfully prepared by physical vapor deposition (PVD) techniques, and the Mo-N phase structure was mainly dependent on the deposition parameters, e.g., the nitrogen partial pressure, working temperature, and substrate bias voltage [5,6,7]. Due to the low solubility of nitrogen in the metal Mo [8,9], three stable phases, including the γ-Mo₂N (face-center-cubic) phase at high temperature, the β-Mo₂N (tetragonal) phase at low temperature, and the δ-MoN (hexagonal) phase, and one thermodynamically unstable phase, B1-MoN (face-center-cubic), appeared in the phase diagram of Mo-N. The transformation of phase structure directly affected the mechanical properties of Mo-N coatings. For instance, the Mo-N coatings containing the γ-Mo₂N phase exhibited the highest hardness, and then decreased when the B1-MoN phase appeared [10]. With increasing nitrogen partial pressure, the microstructure of the Mo-N coatings transformed from γ-Mo₂N to δ-MoN and B1-MoN, which significantly influenced the hardness and residual stress [11].

To further enhance the mechanical properties of Mo-N based hard coatings, three different structural and composition designs have been put forward. A nanocomposite structure has been formed by incorporating some nonmetallic elements (e.g., Si and C) into the Mo-N coatings [12]. For instance, Heo et al. [13] designed Mo-Si-N coatings and found that the hardness increased to a maximum value of 37 GPa at 10 at.% Si due to grain refinement. Second, by incorporating some soft metals (e.g., Ag and Cu) into the Mo-N coatings, a nanocomposite structure can be formed, which not only improves the coating hardness [14], but also reduces the friction coefficient due to the formation of lubrication oxides (e.g., MoO₃ and CuMoO₄) [15,16]. Third, Mo-N coatings alloy with some metallic elements (e.g., Al and V) and form a multicomponent structure, which enhances the hardness, wear resistance, and oxidation resistance, such as Mo-Al-N [17] and Mo-V-N [18]. Recently, a novel structure has been put forward by combining the advantages of nanocomposites and multicomponent structures. For instance, MoVN-Ag coatings have been prepared by magnetron sputtering [19]. The tribological properties were improved by the addition of Ag, but the mechanical properties were damaged, resulting in a low hardness of less than 5 GPa. The lubrication oxides MoO₃ (φ = 8.2) and V₂O₅ (φ = 10.2) had a low friction coefficient at high temperatures due to their high ionic potential. A combination of lubrication oxides (e.g., MoO₃ and V₂O₅) and soft metals like Cu with low shear strength exhibited excellent tribological performance in a wide temperature range [20,21]. In addition, due to a high degree of ionization, high power impulse magnetron sputtering (HIPIMS) has been widely applied for the deposition of hard coatings [22]. As a self-lubricating coating applied to cutting tools, the MoVN-Cu coatings have been deposited by HIPIMS in previous studies, which mainly focused on the influence of the deposition process, such as the nitrogen partial pressure [6,7] and charge voltage [20], and found that mainly B1-MoN and hex δ-MoN phases formed in the coatings. However, the γ-Mo₂N phase with the highest hardness [10] was not formed in the coatings. In addition, a higher Cu content led to a lower friction coefficient but reduced the wear resistance due to lower hardness and poor oxidation resistance [21].

To meet the high cutting performance of coated tools, the mechanical properties, including the hardness and adhesion strength, of Mo-V-Cu-N coatings should be improved further. However, few studies have focused on improving the mechanical properties of Mo-V-Cu-N coatings by optimizing the elemental composition and microstructure. In this study, the Mo-V-Cu-N coatings with various chemical compositions were deposited by the HIPIMS technique, and the effects of V and Cu on the microstructure evolution, surface roughness, residual stress, and mechanical properties of Mo-N based coatings were studied in detail.

2. Experimental Details

2.1. Coating Deposition

The Mo-V-Cu-N coatings were prepared on WC-Co cemented carbide and 316 L stainless steel substrates using high-power impulse magnetron sputtering (HIPIMS). A schematic diagram of the Mo-V-Cu target (69 mm × 443 mm, 99.9% purity) is shown in Figure 1. The sample distances from the top of the target were set at 30, 90, 150, 210, and 270 mm. To remove the scratches on the substrate surface, 1500 G and 3000 G diamond grinding discs were used to grind the substrate surface for 15 min, respectively. Then, a diamond polishing paste with a particle size of 1 μm was used to polish the substrate surface for 10 min to achieve the mirror polishing effect. All the polished substrates were ultrasonically cleaned, and then placed on a substrate holder that rotated at a speed of 4 r/min to keep all the substrates facing the target, whereas the turntable did not rotate. Before deposition, the chamber was pumped to less than 5.0 × 10⁻³ Pa and then heated to 200 °C. To clean surface contaminants, the substrate surface was etched by Ar⁺ plasma at a high bias voltage of 1000 V for 30 min. In addition, due to a high ionization rate and ion bombardment, the substrate surface was bombarded by Cr⁺ ions using arc ion plating (AIP) to enhance the adhesion strength, which operated at a high bias voltage of 800 V for 5 min. Then, a mixed gas of Ar (35 sccm) and N₂ (10 sccm) with a pressure of 0.6 Pa was introduced. The Mo-V-Cu-N coatings with various compositions were deposited by HIPIMS. The target power and duty cycle were kept at 1.5 kW and 2%, respectively. During deposition, the DC bias voltage of the substrate and the deposition time were set at 200 V and 400 min, respectively.

2.2. Coating Characterization

The surface and cross-section morphologies of as-deposited coatings were observed by using scanning electron microscopy (SEM, Nano430, Amsterdam, The Netherlands), and the elemental composition was measured by energy dispersive X-ray spectroscopy (EDS). The crystal structure was identified by X-ray diffraction (XRD, Bruker D8, Billerica, MA, USA) with Cu Kα radiation (λ = 1.5406 Å). The microstructure was characterized by transmission electron microscopy (TEM, Talos F200X, Thermo Fisher Scientific, Waltham MA, USA) operating at an accelerating voltage of 200 kV. The TEM samples were prepared by a focused ion beam (FIB, Auriga Carl Zeiss, Oberkochen, Germany) with Ga⁺ ions. The surface roughness was measured by atomic force microscopy (AFM, Bruker, Billerica, MA, USA) using a silicon nitride tip, which operated in a contact mode, and the scan area was 5 × 5 μm². The residual stress was measured by a film stress tester (FST-1000, Supro Instruments, Shenzhen, China) based on Stoney’s equation [23]:

$$\sigma =\frac{E_s}{6\left(1-{\upsilon }_s\right)}\frac{h_s^2}{h_c}\left(\frac{1}{R_c}-\frac{1}{R_0}\right)$$

(1)

where E_s, υ_s, and h_s refer to the elastic modulus (195 GPa), Poisson’s ratio (0.29), and thickness (0.8 mm) of the 316 L substrate, respectively. h_c and R_c refer to the thickness and curvature radius of the coatings, respectively.

The hardness and elastic modulus were determined by a nanoindentation tester (NHT², CSM, Peseux, Switzerland). To eliminate the substrate effect, a maximum load of 10 mN was applied for each test. At least five effective measurements were carried out for each sample. The adhesion strength of the coatings was evaluated by a scratch tester (RST, CSM, Peseux, Switzerland). For each sample, at least three effective measurements were made, which operated at a scratch length of 3 mm and a maximum load of 100 N.

3. Results and Discussion

3.1. Chemical Composition and Microstructure

According to the design of the sample positions in Figure 1, Mo-V-Cu-N coatings with various element compositions can be easily obtained. The chemical compositions and deposition rates of the coatings are summarized in

Table 1. Chemical composition and deposition rate of Mo-V-Cu-N coatings.

Sample	Coating	Chemical Composition (at.%)				N/(Mo + V) Ratio	Thickness (μm)	Deposition Rate (nm/min)
		Mo	V	Cu	N
S1	Mo45V1Cu30N24	45.4	1.0	29.7	23.9	0.52	6.2 ± 0.1	15.4 ± 0.3
S2	Mo58V4Cu7N31	58.5	3.6	6.5	31.4	0.51	4.7 ± 0.1	11.8 ± 0.3
S3	Mo53V8Cu2N37	52.7	8.2	1.8	37.3	0.61	4.5 ± 0.1	11.3 ± 0.2
S4	Mo46V15Cu1N38	46.5	14.6	1.2	37.7	0.62	3.3 ± 0.1	8.3 ± 0.2
S5	Mo38V19Cu1N42	37.6	19.0	0.9	42.5	0.75	2.6 ± 0.1	6.5 ± 0.2

. As the sample position gradually moved from S1 to S5, an increasing V content from 1.0 at.% to 19.0 at.% and a decreasing Cu content from 29.7 at.% to 0.9 at.% could be obtained. The Mo content increased from 45.4 at.% for S1 to 58.5 at.% for S2, and then decreased gradually to 37.6 at.% for S5. Similarly, the N content increased gradually from 23.9 at.% to 42.5 at.%, which exhibited a similar trend to the V content. Thus, the coatings changed from Mo₄₅V₁Cu₃₀N₂₄ to Mo₃₈V₁₉Cu₁N₄₂ based on the stoichiometric ratio. The calculated N/(Mo + V) atomic ratio for the Mo₄₅V₁Cu₃₀N₂₄ coating was 0.52, which was close to the stoichiometric ratio of the Mo2N phase. The N/(Mo + V) ratio gradually increased to 0.75 with increasing V content, indicating the presence of the VN phase. With an increasing V/Cu ratio, the deposition rate sharply decreased from 15.4 to 6.5 nm/min. A similar result was also found in the AlTiVN-Cu coatings, in which the deposition rate increased linearly with the increasing Cu/V ratio [24]. When compared to Mo and V metals, Cu has a much higher sputtering yield [25], resulting in an increase in the deposition rate.

Figure 2 displays the surface SEM morphologies of Mo-V-Cu-N coatings. In Figure 2a, the Mo₄₅V₁Cu₃₀N₂₄ coating presented a smoother surface than other compositions. However, some small microparticles were observed on the smooth surface, which is a typical growth defect that is induced by the deposition of arc ion plating (AIP) [26]. In our previous study [6], the Mo-Cu-V-N coatings deposited by HIPIMS exhibited a smooth surface without any microparticles observed. A similar phenomenon was also observed in Mo-Cu-N [15], Mo-V-N [18], and Mo-V-Ag-N [19] coatings that were deposited by magnetron sputtering. In this study, before coating deposition, the Cr⁺ ions were introduced to clean the substrate surface. Thus, these microparticles were mainly caused by the evaporation of the Cr target with AIP during the etching process. As V content increased, the microparticles on the coating surfaces increased, including both the microparticle sizes and numbers, which can be related to the decrease in the coating thickness. A similar phenomenon was found in our previous study [21]. These microparticles formed in the bottom layer were gradually covered with increasing coating thickness. In addition, with the increase of V content, the surface morphology changed from a smooth surface to a scale-like structure, as shown in Figure 2b,c, and then to a plate-like structure, as shown in Figure 2d,e. Compared to Cu, Mo and V have a higher ion/atom ratio in the plasma. More Mo and V ions can be accelerated to the coating surface under the bias voltage of the substrate, which enhances the ion bombardment effect during deposition. The formation of scale-like and plate-like structures would be induced by the enhanced ion bombardment and etching effect. A similar plate-like structure was also observed for the deposition of AlTiVCuN coatings due to strong ion bombardment [27].

Figure 3 and Figure 4 show three-dimensional AFM images of Mo-V-Cu-N coatings that were deposited on WC-Co cemented carbide and 316 L stainless steel substrates. Before deposition, the polished WC-Co and 316 L substrates showed extremely smooth surfaces with low surface roughnesses of 2.7 and 0.6 nm, respectively, as shown in Figure 3a and Figure 4a. After deposition, the coated substrates exhibited uneven three-dimensional morphologies, especially for the WC-Co substrates. During the deposition process, some atoms were sputtered and ionized, such as Mo⁺ and V⁺ ions, which reacted with nitrogen and then bombarded the substrate surfaces under a bias voltage, resulting in a typical atomic peening effect [28]. Therefore, with increasing V content, the (Mo + V) content also increased, which enhanced the ion bombardment, leading to an increase in the surface roughness. As shown in Figure 5, for the cemented carbide and stainless steel substrates, the surface roughness of the coatings presented a similar trend with increasing V content, which increased from 7.7 and 2.0 nm at 1.0 at.% V to 26.2 and 14.7 nm at 14.6 at.% V, respectively. When the V content increased to 19.0 at.%, the surface roughness decreased to 18.5 and 12.5 nm, respectively, indicating that the intensity of ion bombardment was weakened. Compared to the V atoms, Mo has a higher ion valence and relative atomic mass, contributing to much stronger ion bombardment during the deposition process. Thus, the reduction in surface roughness can be related to a sharp decrease in the Mo content in the coatings. Compared to the WC-Co substrates, the surface roughness of the coatings deposited on 316 L substrates was much lower, demonstrating that a polished substrate with low surface roughness contributed to the deposition of smooth coatings.

Figure 6 shows the cross-section morphologies of Mo-V-Cu-N coatings at different V contents. When the V content increased from 1.0 to 19.0 at.%, the coating thickness sharply decreased from 6.2 to 2.6 mm, which was mainly related to a sharp decrease in the Cu content in the coatings [24]. A similar phenomenon was also reported for the MoVN-Ag coatings [19], due to a high sputtering rate of Ag, the thickness gradually increased with increasing Ag content. All the cross-sections exhibited a smooth interface without obvious cracks or voids, implying that a good adhesion was obtained between the coatings and WC-Co substrates. In Figure 6a, the Mo₄₅V₁Cu₃₀N₂₄ coating exhibited a featureless structure. A similar structure was also reported in the AlTiVN-Cu coatings [24], in which the growth of columnar crystals was restricted by the addition of Cu and formed a featureless structure as the Cu content reached 22.6 at.%. In addition, a smooth surface can be clearly seen from the cross-section, which was consistent with above SEM and AFM results. With increasing V content, the coatings changed into a dense columnar structure, which was mainly induced by the enhanced ion bombardment. It has been reported that the adatom mobility was enhanced by ion bombardment, and promoted the migration of grain boundaries, then formed a dense microstructure [29].

Figure 7 presents the XRD patterns of Mo-V-Cu-N coatings at different V contents. Among these peaks, all the sharp peaks of the cemented carbide substrate corresponded to the hexagonal close-packed WC (PDF#51-0939) and face-center-cubic (FCC) TiC (PDF#32-1383) phases. For the coatings, three broad diffraction peaks at approximately 37.9°, 43.1°, and 75.4° corresponded to the (111), (200), and (311) planes of the face-center-cubic (FCC) phase, respectively. Due to the high content of Mo, the Mo-N phase is mainly formed in the coatings. Moreover, due to the similar atomic radius of V (0.135 nm) and Mo (0.136 nm), the V atoms tended to dissolve in the Mo-N lattice and formed a single solid solution phase. A similar solid solution phase of c-Mo-V-N was found in the Mo₆₇V₄N₂₉ coating [18] and the V_0.57Mo_0.43N_0.95 coating [30]. Regarding the existence of Cu in the coatings, the chemical bonding states of the Mo-Cu-V-N coatings have been analyzed by XPS in previous work [7], and it was found that Cu existed as a metallic species instead of copper nitride in the coatings. Furthermore, it has been reported that when the Cu content in the Mo-Cu-N coatings exceeded 14.0at.%, a double-phase of B1-MoN and c-Cu formed [15]. However, for the Mo₄₅V₁Cu₃₀N₂₄ coating with a high Cu content of 29.7 at.%, the (111) diffraction peak of the Cu phase at approximately 43.3° was overlapped with the (200) peak of the Mo-V-N phase, and no other diffraction peaks of the Cu phase could be seen in the XRD pattern, which would be further analyzed by TEM.

With increasing V content, the diffraction peaks of nitride shifted toward lower angles, implying an increase in the lattice parameter. As listed in

Table 2. Lattice parameters of Mo-V-Cu-N coatings with various V contents.

Planes	Lattice Parameter a0 (Å)
	1.0 at.% V	3.6 at.% V	8.2 at.% V	14.6 at.% V	19.0 at.% V
(111)	4.112	4.164	4.143	4.143	4.167
(200)	4.194	4.227	4.225	4.198	4.184
(311)	4.180	4.200	4.200	4.190	4.190
Mean	4.162	4.197	4.189	4.177	4.180
Stdev	0.044	0.032	0.042	0.030	0.012

, the lattice parameters of the coatings were obtained from the (111), (200), and (311) peaks. With the increase of V content, all the lattice parameters were ranged from 4.162 Å to 4.197 Å, corresponding to the FCC γ-Mo₂N phase (PDF#25-1366), which is a stable phase with a face-center-cubic structure. With increasing N content from 5.1 at.% to 50.8 at.%, the phase structure of Mo-N coatings transformed from α-Mo to γ-Mo₂N phase, and then to B1-MoN phase [10]. However, when the N content increased above 55.5 at.%, the phase structure of Mo-Cu-V-N coatings transformed from B1-MoN into δ-MoN phase [6]. In this study, the N contents in the coatings varied from 23.9 to 42.5 at.%, then a γ-Mo₂N phase can be formed. The Mo₄₅V₁Cu₃₀N₂₄ coating showed a strong preferred orientation, which corresponded to the (200) plane of Mo-V-N phase and/or the (111) plane of Cu phase. With increasing V content, a weak (220) diffraction peak appeared at 62.2°, and the coatings changed into no obvious preferred orientation.

Figure 8 presents the cross-sectional TEM and HRTEM images of the Mo₄₅V₁Cu₃₀N₂₄ coating. As shown in Figure 8a, a clear interface between the coating and substrate can be identified in the HAADF image. The EDX element mappings showed that all the elements, including Mo, V, Cu, and N, were uniformly distributed in the cross-section and formed a typical nanocomposite structure. From the TEM bright-field image in Figure 8b, a dense columnar structure formed along the growth direction. The SAED pattern in Figure 8c showed incontinuous diffraction points belonging to the (200), (222), and (420) planes of the c-Mo₂(V)N phase, and the (111), (220), and (311) planes of the c-Cu phase, respectively. The d-spacings of the Mo2(V)N and Cu phases in

Table 3. Measured d-spacing of the Mo₄₅V₁Cu₃₀N₂₄ coating according to the SAED pattern.

Planes	Measured d-Spacing (Å)		Standard d-Spacing (Å)
	Mo2(V)N	Cu	Mo2N (PDF#25-1366)	Cu (PDF#85-1326)
(111)	-	2.084	-	2.087
(200)	2.077	-	2.081	-
(220)	-	1.278	-	1.278
(222)	1.185	-	1.202	-
(311)	-	1.091	-	1.090
(400)	1.043	-	1.041	-
(420)	0.918	-	0.931	-

were in good agreement with that of the standard Mo₂N (PDF#25-1366) and Cu (PDF#85-1326) phases, implying that a double-phase of c-Mo₂(V)N and c-Cu existed in the coating. In addition, based on the high-resolution HRTEM image and FFT pattern in Figure 8d, the lattice fringe spacing in region A was 0.2076 nm, which can be identified as the (200) plane of the c-Mo₂(V)N phase. In region B, two lattice fringes with spacings of 0.2083 and 0.1832 nm corresponded to the (111) and (200) planes of the c-Cu phase, respectively. Thus, it demonstrated the existence of a double-phase structure of c-Mo₂(V)N and c-Cu in the Mo₄₅V₁Cu₃₀N₂₄ coatings.

Figure 9 displays the cross-sectional TEM and HRTEM images of the Mo₃₈V₁₉Cu₁N₄₂ coating. From the HAADF image and corresponding EDX element mapping in Figure 9a,b, it can be seen that Cu element was uniformly distributed in the cross-section. In Figure 9c, some incontinuous diffraction rings can be observed in the SAED pattern, corresponding to the (111), (200), (220) and (311) planes, implying that a single solid solution phase of c-Mo₂(V)N formed in the coating, which was consistent with above XRD results. However, no diffraction rings of the Cu phase appeared in the SAED pattern. In addition, based on previous XPS analysis of the Mo-Cu-V-N coatings [7], Cu was found to exist as a metallic species instead of Cu-N or Mo-V-Cu-N in the coatings. Similar results have been found in the Ti-Cu-N [31] and TiAlN-Cu [32] coatings that Cu existed as a metallic species rather than dissolved into the Ti-N and Ti-Al-N lattices. Due to Cu being immiscible in the hard phase (MeN, Me = Ti, Zr, Mo, etc.), soft metallic copper is often added as a second phase into the transition metal nitride and formed a typical nanocomposite structure [33]. In addition, due to an excellent diffusion barrier capability against Cu [34], most Cu atoms existed at the grain boundary rather than in the Mo₂N grains [14]. A similar result was also found in the Mo-Cu-N coatings that Cu atoms existed as an amorphous phase at the grain boundary when Cu content was below 11 at.% [15]. In this study, due to the low contents (0.9–6.5 at.%), the metallic Cu might also exist as an amorphous phase in the Mo-V-Cu-N coatings. From the HRTEM image in Figure 9d,e, two lattice fringes with spacings of 0.2395 and 0.2078 nm can be identified. These values of d-spacing were between those of the γ-Mo₂N (0.2404 and 0.2081 nm) and c-VN (0.2390 and 0.2070 nm) phases, respectively. This demonstrated that V atoms dissolved into the γ-Mo₂N lattice, and formed a solid solution phase of c-Mo₂(V)N. A schematic diagram of the (220) plane of the Mo₂(V)N phase is presented in Figure 9f.

3.2. Mechanical Properties

In general, as for nitride coatings, the residual stress includes intrinsic and extrinsic stresses [35]. Among which, the intrinsic residual stress is often generated by deposition parameters (e.g., target power, duty cycle, and deposition temperature). The extrinsic residual stress is mainly generated by the mismatch of thermal expansion between the coatings and substrates, which often corresponds to the thermal stress. In this study, the residual stresses of nitride coatings were calculated based on Equation (1). Figure 10 shows the residual stress of the Mo-V-Cu-N coatings as a function of V content. The Mo₄₅V₁Cu₃₀N₂₄ coating exhibited a relatively low compressive residual stress of 1.46 GPa, which was mainly caused by the addition of a high content of soft metal Cu [36]. With increasing V content, the compressive residual stress sharply increased to 4.20 GPa at 14.6 at.% and then decreased to 4.00 GPa at 19.0 at.%. As discussed above, with increasing V content, the ion bombardment and defect concentration were enhanced, leading to an increase in the residual stress [37]. In contrast, due to the presence of a compliant copper phase, the increase in Cu content in the coatings resulted in a relaxation of the residual stress [24,36].

Figure 11 presents the hardness and elastic modulus of Mo-V-Cu-N coatings as a function of V content. At 1.0 at.% V, the Mo₄₅V₁Cu₃₀N₂₄ coating showed an extremely low hardness of 11.5 GPa and an elastic modulus of 243.1 GPa. When the V content was increased, the hardness and elastic modulus increased to 27.3 GPa and 374.3 GPa at 14.6 at.%, and then decreased to 26.4 GPa and 350.9 GPa at 19.0 at.%, respectively. The initial increase in the hardness and elastic modulus can be explained by the following aspects: Firstly, the increase in hardness can be caused by microstructure densification [16]. With increasing V content, the microstructure of the coatings transformed from a featureless structure into a dense columnar structure, as shown in Figure 6; Secondly, the hardness was also enhanced by the formation of the c-Mo₂(V)N solid solution phase. A similar solid solution hardening effect was also observed in the Mo-V-N coatings [18], in which the Mo₅₂V₂₂N₂₆ coating exhibited the highest hardness due to the solution strengthening effect; Thirdly, the enhanced hardness can also be related to an increase in the residual stress. Generally, a higher hardness in nanocomposite coatings is always accompanied by a higher compressive residual stress [32]. In addition, as a soft metal, the reduction of Cu content also contributed to an increase in the coating hardness and elastic modulus [14]. However, when the V content increased to 19.0 at.%, the slight decrease in the hardness would be related to the relaxation of residual stress [38]. In addition, the H³/E*² ratio of the coatings can also be used to evaluate the resistance against crack initiation and propagation [39,40], where E* = E/(1 − ν²) is the effective Young’s modulus and ν is the Poisson’s ratio. With increasing V content, the H³/E*² ratio increased from 0.023 to a maximum value of 0.136, and then slightly decreased to 0.131.

To evaluate the adhesion strength between the coating and substrate, the scratch tracks were characterized by optical microscopy (OM), as shown in Figure 12. Among which, the adhesive failure mode L_C2 refers to adhesive chipping at track edges, which is often used to determine the adhesion strength between the coating and substrate. In Figure 12a, at a low V content of 1.0 at.%, many adhesive chips appeared along the track edges, indicating that the Mo₄₅V₁Cu₃₀N₂₄ coating showed a low adhesion strength, which can be mainly related to the low H³/E*² ratio of 0.023. A similar result was also reported in a previous study [24]. The AlTiVN-Cu (46.7 at.%) coating exhibited a low adhesion strength of 6.6 N due to a low H³/E*² ratio of 0.008. With increasing V content, fewer chips along the track edges were observed, indicating that the adhesion strength was significantly improved. As for the Mo₃₈V₁₉Cu₁N₄₂ coating, it can be seen that some black adhesive chips along the track edges were distributed continuously in the interval of 80 to 100 N, and the white arrow corresponded to the critical load to adhesive failure, thus a high adhesion strength was achieved. In Figure 12b, the adhesion strength gradually increased from 32.2 to 87.8 N with increasing V content. The increase in adhesion strength would be attributed to the combined effects of the high H³/E*² ratio and appropriate compressive residual stress [41]. In addition, the enhanced adhesion strength was also related to the reduction of Cu content. Due to low adhesion energy, the addition of Cu to the Mo-N coatings adversely affected the coating/substrate bond [36].

4. Conclusions

In this study, the Mo-V-Cu-N coatings with various chemical compositions were deposited by HIPIMS, and the additions of V and Cu on the microstructure evolution, surface roughness, and mechanical properties of Mo-N based coatings were explored. The main results can be concluded as follows:

(1)

With an increasing V/Cu content ratio, the deposition rate decreased from 15.4 to 6.5 nm/min, and the coatings transformed from Mo₄₅V₁Cu₃₀N₂₄ to Mo₃₈V₁₉Cu₁N₄₂;

(2)

At low Cu contents, less than 6.5 at.%, the coatings exhibited a single solid solution phase of c-Mo₂(V)N. When the Cu content reached 29.7 at.%, the Mo₄₅V₁Cu₃₀N₂₄ coating showed a double-phase of c-Mo₂(V)N and c-Cu;

(3)

Compared to the WC-Co substrates, the surface roughness of the coatings that were deposited on 316 L substrates was much lower, and the Mo₄₅V₁Cu₃₀N₂₄ coating showed the lowest surface roughness of 2.0 nm;

(4)

With increasing V/Cu content ratio, the microstructure transformed from a featureless structure into a dense columnar structure, contributing to an increase in the coating hardness and adhesion strength.