Oxidation-resistant Zr–Mo–Si–B coatings deposited by DCMS and HIPIMS methods

Introduction

Zirconium disilicide (ZrSi₂), owing to its high strength characteristics and oxidation resistance, is a promising material for use as a base in the development of high-temperature oxidation and wear-resistant coatings [1; 2]. The protective properties of ZrSi₂ are due to the formation of a surface layer of silicon dioxide (SiO₂), which prevents the diffusion of oxygen atoms into the material and possesses high viscosity, aiding in the healing of cracks formed during heating. In multicomponent and multiphase materials, the local rearrangement of particles caused by the phase transition from monoclinic zirconium oxide (ZrO₂) to tetragonal ZrO₂ serves as an additional mechanism for defect healing [3]. Moreover, ZrO₂, formed as a result of ZrSi₂ oxidation, has a high melting point and can react with SiO₂ to form zirconium silicate (ZrSiO₄), which is highly chemically stable and has low oxygen permeability [2; 4; 5].

Despite the high potential for practical application, information on the development of materials based on ZrSi₂ is limited. Known works [6; 7] describe ZrSi₂–MoSi₂–ZrB₂, ceramics obtained by self-propagating high-temperature synthesis (SHS). Due to the formation of a multilayer structure consisting of a ZrSiO₄ layer and sublayers based on ZrO₂ and SiO₂, these ceramics exhibit high oxidation resistance at 1650 °С.

In addition to the development of bulk materials based on ZrSi₂, researchers are also focusing on creating coatings with a similar composition. ZrSi₂-based coatings are widely used to protect zirconium alloys [8–10], C/C composites [11–14], and graphite [15; 16] from oxidation at high temperatures. Studies [8; 9] have shown that the deposition of ZrSi₂ coatings reduces the thickness of the oxide layer on zirconium alloys by a factor of 10 at temperatures of 1000 and 1200 °C. The introduction of molybdenum disilicide (MoSi₂), which has self-healing properties, into ZrSi₂ coatings is of interest [3]. The addition of MoSi₂ reduces the likelihood of the phase transition of ZrO₂ from the tetragonal to monoclinic phase, which causes volume expansion and leads to coating cracking [12]. In a study [17], ZrSi₂–MoSi₂–ZrB₂, coatings obtained by slip-firing method exhibited record oxidation resistance at temperatures of 1500–1700 °C. This can be explained by the following factors:

– formation of a thermally stable heterogeneous oxide film;

– dissolution of ZrO₂ in the borosilicate layer forming ZrSiO₄;

– formation of thermally stable molybdenum-based particles that hinder oxygen diffusion;

– high melting point and low thermal conductivity of the ZrO₂-based oxide layer, protecting the inner layers.

Despite the high oxidation resistance of ZrSi₂–MoSi₂–ZrB₂, coatings obtained by the slip-firing method, the application of this method is limited by the size of the treated parts. Additionally, it is not possible to accurately control the phase composition and thickness of the coating. In this regard, the direct current magnetron sputtering (DCMS) method appears promising, as it allows the production of defect-free coatings with a specified phase composition, uniform thickness, high adhesive strength, and the ability to reinforce complex-shaped products [18; 19].

The application of high-power impulse magnetron sputtering (HIPIMS) opens up additional possibilities for applying ceramic coatings [20; 21]. Due to its higher power, HIPIMS significantly increases the plasma density from ~10¹⁰ ions/cm3 for DCMS to 10¹³–10¹⁴ ions/cm³ for HIPIMS [22]. In the latter case, sputtered atoms are intensely ionized while passing through the plasma, and the flow consists predominantly of ions rather than atoms, as in DCMS. The increased ion/atom ratio in the flow characteristic of HIPIMS leads to a significant enhancement in the adhesive strength of the deposited coatings due to the formation of pseudo-diffusion layers and ion implantation effects during the preliminary etching of the substrate surface [22]. The increased density of the structure and adhesive strength improve the mechanical properties, wear resistance, and oxidation resistance of the coatings [23; 24].

The aim of this study was to develop Zr–Mo–Si–B system coatings using DCMS and HIPIMS methods, and to investigate their structure, mechanical properties, and high-temperature oxidation resistance.

Materials and methods

The coatings were produced using the DCMS and HIPIMS methods with a functionally graded target (ZrSi₂–ZrB₂–MoSi₂ )/Cr [25]. The coatings were produced using the DCMS and HIPIMS methods with a functionally graded target (ZrSi₂–ZrB₂–MoSi₂)/Cr [25]. Deposition was carried out in a vacuum chamber based on the UVN-2M pumping system (Quartz JSC, Russia) [26]. For the DCMS mode, a Pinnacle Plus power supply (Advanced Energy, USA) was used, while HIPIMS deposition was performed using a TruPlasma 4002 system (Trumpf, Germany). The average power for coating deposition using both methods was P = 1 kW. To compensate for the lower growth rate in the HIPIMS mode, experiments were conducted at an increased power of P = 2 kW. The HIPIMS frequency was 1 kHz, with a pulse duration of 50 μs, and the residual and working pressures of Ar (99.9995 %) were 3·10\(^–\)³ and 1·10\(^–\)¹ Pa, respectively. Polycrystalline Al₂O₃ plates (VOK-100-1 grade) and Si (KEF-100 grade) were used as substrates. To remove surface contaminants before sputtering, the substrates were cleaned ultrasonically in isopropyl alcohol for 5 min and then cleaned with Ar\(^+\) ions at 2 keV for 20 min using an ion source. The deposition time was 60 min for both modes (P = 1 kW) and 30 min for HIPIMS at P = 2 kW.

The magnetron discharge plasma was studied using optical emission spectroscopy with a PlasmaScope spectrometer (Horiba Jobin Yvon, France). The composition and structure of the coatings were analyzed using a scanning electron microscope (SEM) S-3400N (Hitachi, Japan) equipped with a Noran-7 Thermo energy-dispersive spectroscopy (EDS) attachment. The surface topography of the coatings was investigated using a WYKO-NT1100 optical profilometer (Veeco, USA). The phase composition of the coatings was determined by X-ray diffraction (XRD) using a D2 Phazer diffractometer (Bruker, Germany) with CuK_α radiation. Elemental profiles of the coatings were obtained by optical emission spectroscopy with a glow discharge (GDOES) using a Profiler 2 instrument (Horiba JY, France) [27].

Mechanical properties such as hardness (H), elastic modulus (E), and elastic recovery (W) were evaluated by nanoindentation using a Nanohardness Tester (CSM Instruments, Switzerland) at a load of 4 mN. To determine the oxidation kinetics, stepwise annealing was performed in air in a muffle furnace SNOL-7.2/1200 (AB Umega, Lithuania) at 1000 °C with hold times of 10, 30, 60, and 180 min. To assess the high-temperature oxidation resistance of the coatings, non-isothermal annealing was conducted in a muffle furnace TK 15.1800.DM.1F (Thermoceramics LLC, Russia) at temperatures of 1100, 1300, and 1500 °C with a hold time of 10 min. After annealing, the coatings were examined using SEM, EDS, and XRD methods. The designations of the coatings and their deposition modes are presented in the Table below.

Results and discussion

Plasma diagnostics during coating deposition

Fig. 1 presents the plasma spectra obtained by sputtering of a Zr–Si–Mo–B target in DCMS and HIPIMS modes at an average power of 1 kW, as well as in HIPIMS mode at P = 2 kW.

Plasma analysis in the wavelength range of 200 to 880 nm was conducted for the elements present in the target, as well as for the working gas (Ar). The main characteristic line positions, according to the Quantum XP software (Horiba JY, France), are 252, 288, and 386 nm for Si; 399 nm for Zr; 313 and 380 nm for Mo; 250 nm for B; and 603, 697, 707, 750, 801, 811, and 842 nm for Ar. When transitioning from DCMS to HIPIMS mode (with an average power of 1 kW in both cases), the intensities of the Mo, B, and Si peaks increased by 1.7, 3.7, and 2.5 times, respectively. The most significant change was observed for the Zr signal, with peak intensities increasing 14-fold, reaching a maximum when transitioning from DCMS to HIPIMS at P = 1 kW. The intensity of the Ar peaks during HIPIMS deposition decreased by 1.3 to 3.7 times compared to DCMS. Increasing the power from 1 to 2 kW during HIPIMS deposition resulted in an increase in the intensity of the Si, Mo, and B lines by 1.6, 1.7, and 1.5 times, respectively, while the intensity of the Zr remained unchanged.

Thus, transitioning from DCMS to HIPIMS at an average power of 1 kW allows for the generation of plasma predominantly consisting of target material ions [28; 29], with the integrated intensity, calculated using MagicPlot Pro software, increasing threefold. Increasing the power from 1 to 2 kW during HIPIMS operation results in a further twofold increase in this indicator.

Study of coating structure and mechanical characteristics

Typical SEM images of the structure, elemental profiles, and three-dimensional surface images (using coating 1 as an example) are shown in Fig. 2.

All samples exhibited a dense, defect-free structure and a similar roughness parameter R_a in the range of 7–8 nm. The structure of the initial coatings was studied in more detail in [25]. According to GDOES data (Fig. 2, b), all elements were uniformly distributed across the thickness of the coatings (refer to the Table).

The thickness of coatings 1, 2, and 3 reached 8.4, 5.7, and 5.6 μm, respectively. The growth rate of sample 1 was 140 nm/min, but switching to HIPIMS mode reduced it by 1.5 times. This reduction can be attributed to self-sputtering effects and the decreased overall sputtering time in HIPIMS pulse mode [30; 31]. It is worth noting that increasing the power during coating deposition under these conditions does not affect the composition or growth rate of the coatings.

According to our previously obtained data [25], all coatings, regardless of the deposition mode, are X-ray amorphous.

Nanoindentation results (refer to the Table) showed that all samples had similar mechanical characteristics: hardness of 11–14 GPa, elastic modulus of 177–208 GPa, and elastic recovery of 37–42 %. It is noteworthy that a hardness level of 10–15 GPa is characteristic of silicide ceramics based on ZrSi₂ [1; 32].

Study of high-temperature
oxidation resistance of coatings

Fig. 3 shows the specific mass change (∆m/S) of the coatings as a function of holding time, as well as the appearance of the coatings before and after annealing at 1000 ºC. For all samples, an increase in the parameter ∆m/S was observed for holding times up to 30 min, associated with the growth of the oxide film on the surface of the coatings. The decrease in this parameter observed for all coatings during the holding period of 30–60 min may be related to the partial evaporation of MoO_x. In the range of 60–180 min, the value of ∆m/S increased by 25 and 50 % for samples 1 and 3, respectively, while for coating 2 it remained unchanged. The maximum values of ∆m/S, 11.2 and 10.5 mg/cm², were recorded for samples 1 and 2, while coating 3, produced in HIPIMS mode at an average power of 2 kW, exhibited the lowest specific mass change of 5.3 mg/cm², indicating its superior oxidation resistance under stepwise annealing at 1000 °C. The visual analysis of the samples (Fig. 3, b) showed that coatings 1 and 3 maintained their integrity during the tests, with no visible damage (delamination or cracking), whereas coating 2 partially delaminated during the testing.

To determine the maximum working temperature at which the coatings retain their protective properties, non-isothermal annealing was conducted in the temperature range of 1100–1500 °C. Fig. 4 presents SEM images of the coating surfaces after annealing at 1100 °C. All samples showed areas of damage (blistering), which could be due to the formation and evaporation of MoO_x, as well as phase transformations accompanied by volume changes of the structural components. Similar effects were previously observed in studies of Mo–(Hf, Zr)–Si–B coatings [26]. The size of the local damage areas decreased by 1.3 and 4.0 times when transitioning from coating 1 to coatings 2 and 3, respectively. Particles of ZrO_x and ZrSiO_x, sized 0.4–1.2 µm, were identified on the surfaces of all samples. According to EDS analysis at points on the surfaces of coatings 1 and 2, the formation of MoO_x particles was observed.

Cross-sectional SEM images and EDS maps of the studied coatings after annealing at 1100 °C are presented in Fig. 5. The obtained data showed that an oxide film of equal thickness (~5.0 µm) formed on the surfaces of coatings 1 and 2, consisting of two layers:

– the upper layer (~2.6 µm thick) consisting of an amorphous phase a-Si:B and nanocrystallites of nc-ZrO₂, nc-ZrSiO₄;

– the а-Si:В:O layer (~2.4 µm thick).

A porous oxide layer of а-Si:В:O + nc-ZrO₂ + nc-ZrSiO₄ with a thickness of 4.8 µm formed on the surface of sample 3. Notably, the formation of ZrSiO_x and ZrO₂ effectively enhances the protective properties of the oxide film by increasing stability and reducing defects in the borosilicate glass layer [33; 34]. After annealing at 1300 °C, the surfaces of samples 1, 2, and 3 exhibited blistering areas of approximately 500, 250, and 15 µm, respectively (Fig. 6), as well as cracks, which could be attributed to oxidation and crystallization processes occurring in the coatings during heating [35; 36]. It is worth mentioning that the presence of cracks facilitates oxygen diffusion into the material, potentially compromising its oxidation resistance [37; 38]. It is worth noting that particles of ZrSiO₄ formed in the crack area on the surface of coating 1, which may provide additional healing effects [33].

Fig. 7 shows cross-sectional SEM images of the coatings after annealing at 1300 °C. According to EDS data (mapping and point analysis), the cross-sectional fracture of sample 1 revealed three layers in the oxide film with a thickness of ~6.2 µm:

– the upper porous layer of а-Si:В:O + nc-ZrO₂ + nc-ZrSiO⁴ with a thickness of ~3.0 µm and pore sizes of 0.2–0.6 µm;

– a 200 nm thick layer containing MoO_x particles;

– a layer at the “oxide film-coating” interface, consisting a-Si:B:O and nc-ZrO₂.

The oxide film on the surface of sample 2 consisted of а-Si:В:O + nc-ZrO₂ + nc-ZrSiO₄ with a thickness of ~5.2 µm. The layer near the surface had a porous structure with pore sizes of 0.2–1.2 µm, while the lower layer was dense (see Fig. 7). For coating 3, the thickness of the porous oxide layer а-Si:В:O + nc-ZrO₂ + nc-ZrSiO₄ was minimal, measuring 4.9 µm. The unoxidized layers of all samples contained crystallites of ZrSi₂ and MoSi₂ phases.

The results after oxidation annealing at 1300 °C demonstrated that using the HIPIMS method instead of DCMS reduced the thickness of the surface oxygen-containing layer by 16 %. Increasing the power from 1 to 2 kW during HIPIMS deposition further reduced the thickness of the oxide layer by 6 %. This effect may be related to the densification of the structure and the elimination of structural defects in the coatings when transitioning from DCMS to HIPIMS [39; 40]. It is noteworthy that increasing the power during HIPIMS sputtering enhances ionization rates and ion energy, which positively affects the quality of the resulting coatings [41; 42].

Increasing the annealing temperature to 1500 °C led to the complete oxidation of coatings 2 and 3 due to their small thickness. On the surface of sample 1, in addition to characteristic blistering areas, crack formation was observed (Fig. 8). At higher SEM magnification (×2000), grains of ZrO₂ (the brightest areas in the SEM images) with sizes of 0.6–1.8 µm and grains of ZrSiO₄ with sizes of 1.2–5.0 µm were identified in an amorphous Si:B matrix (the darkest areas).

According to SEM images and EDS maps of the cross-sectional fracture (Fig. 8), coating 1 retained an oxygen-free layer, indicating its high oxidation resistance at 1500 °C. The thickness of the protective layer а-Si:В:O + nc-ZrO₂ + nc-ZrSiO₄, formed on the surface during oxidation was ~8.0 µm. The unoxidized layer contained MoSi₂ grains sized 3.6 to 5.0 µm and Zr₅Si₃ grains sized 0.2 to 1.0 µm. The superior oxidation resistance of coating 1, produced in DCMS mode, can be attributed to its maximum thickness.

X-ray diffractograms of coatings 1–3 annealed at 1100 °C are presented in Fig. 9, a. Peaks corresponding to the phases o-ZrSi₂ (ICDD 032-1499) and t-MoSi₂ (ICDD 41-0612), indicative of the unoxidized layer, were identified in all samples. Peaks corresponding to the oxide phases m-ZrO₂ (ICDD 07-0343) and t-ZrSiO₄ (ICDD 06-0266) were also observed. Coatings 1 and 2 exhibited peaks from the m-MoO₂ phase (ICDD 76-1807).

The crystallite sizes of the main phases, determined using the Debye–Scherrer formula, were similar for all coatings: 50–57 nm for o-ZrSi₂ (131) and t-MoSi₂ (200), 47–50 nm and 57–70 nm for m-ZrO₂ (111) and t-ZrSiO₄ (200), respectively. The crystallite sizes of the m-MoO₂ (200) phase for coatings 1 and 2 were 61 and 68 nm, respectively.

X-ray diffractograms of coatings annealed at 1300 °C (Fig. 9, b) revealed peaks from the hexagonal phase h-Zr₅Si₃ (ICDD 79-4988), in addition to the phases t-MoSi₂, m-ZrO₂, and t-ZrSiO₄, indicating a phase transition from the orthorhombic o-ZrSi₂ phase to the hexagonal h-Zr₅Si₃ phase. The crystallite size of h-Zr₅Si₃ phase increased from 57 to 77 nm when transitioning from coating 1 to coatings 2 and 3.

After heating sample 1 to 1500 °C, the crystallite sizes of the t-ZrSiO₄ and m-ZrO₂ phases were 50 and 43 nm, respectively, and the sizes of the h-Zr₅Si₃ and t-MoSi₂ phases were 85 and 77 nm, respectively. A partial transition of the t-MoSi₂ phase to t-Mo₅Si₃ was observed (Fig. 9, c). The crystallite size of the t-Mo₅Si₃ phase, determined from the (002) line, was 57 nm.

Thus, among all the samples studied, coating 1 exhibited the smallest crystallite sizes for both the main phases (h-Zr₅Si₃ and t-MoSi₂) and the oxygen-containing phases (t-ZrSiO₄ and m-ZrO₂), indicating its higher thermal stability and enhanced protective properties. This assertion is well supported by the fact that only this sample withstood annealing at 1500 °C. It is noteworthy that Zr–Mo–Si–B coatings are comparable in oxidation resistance to previously studied ZrB₂ and ZrSiB coatings [43; 44].

Conclusions

1. Coatings of the Zr–Mo–Si–B system were produced using DCMS and HIPIMS methods with a functionally graded SHS target (ZrSi₂–ZrB₂–MoSi₂)/Cr. Transitioning from DCMS to HIPIMS at P = 1 kW led to an increase in the ionization degree of target component atoms and a threefold rise in the integral intensity of the spectra. Increasing the power from 1 to 2 kW during HIPIMS deposition further doubled this parameter.

2. Regardless of the deposition method (DCMS or HIPIMS), the coatings were X-ray amorphous with a homogeneous structure and uniform elemental distribution throughout their thickness. They exhibited similar values of hardness (11–14 GPa), elastic modulus (177–208 GPa), and elastic recovery (37–42 %). All coatings demonstrated high oxidation resistance in air at t = 1100 and 1300 °C. Transitioning from DCMS to HIPIMS reduced defectiveness and decreased the oxidation depth by 16–21 % at 1300 °C.

3. The 8 µm thick coating showed the highest oxidation resistance at 1500 °C, which is attributed to its thermal stability and the formation of a dense oxide film on the surface with an amorphous matrix of a-Si:B and nanocrystallites of t-ZrSiO₄ and m-ZrO₂.