Influence of Si Content on the Microstructure and Properties of Hydrogenated Amorphous Carbon Films Deposited by Magnetron Sputtering Technique

Abstract

Hydrogenated amorphous carbon (a-C:H) films are widely valued for their excellent mechanical strength and low friction, but their performance significantly degrades at elevated temperatures, limiting practical applications in aerospace environments. In this work, we aimed to enhance the high-temperature tribological behavior of a-C:H films through controlled silicon (Si) doping. A series of a-C:H:Si films with varying Si contents were fabricated via direct current magnetron sputtering, and their microstructure, mechanical properties, and friction behavior were systematically evaluated from room temperature up to 400 °C. Results show that moderate Si doping (8.3 at.%) substantially enhances hardness and wear resistance, while enabling ultralow friction (as low as 0.0034) at 400 °C. This superior performance is attributed to the synergistic effects of transfer layer formation, preferential Si oxidation, and tribo-induced graphitization. This study provides new insights into the high-temperature lubrication mechanisms of Si-doped a-C:H films and demonstrates the critical role of Si content optimization, highlighting a viable strategy for extending the thermal stability and lifespan of solid-lubricating films.

1. Introduction

With the rapid advancement of aviation technology, increasing demands have been placed on the high-temperature lubrication of critical aero-engine components. Conventional liquid lubricants undergo rapid thermal degradation above approximately 350 °C, rendering them unsuitable for such extreme environments [1,2]. Hydrogenated amorphous carbon films (a-C:H), depending on their excellent mechanical strength, chemical inertness, and intrinsically low friction coefficient, have emerged as promising solid-lubrication candidates for aerospace and other high-performance applications [3].

Amorphous carbon lacks a long-range crystalline structure and comprises a metastable mixture of sp²- and sp³-bonded carbon. Under elevated temperatures, this metastability promotes conversion of sp³-C into sp²-C and concurrent hydrogen desorption, leading to structural graphitization and diminished wear resistance. Indeed, non-equilibrium magnetron-sputtered diamond-like carbon (DLC) films prepared by Ni et al. exhibited pronounced frictional degradation beginning at 240 °C and complete failure at 300 °C [4], while microwave sheath-voltage plasma deposition by Deng et al. showed onset of property decay at 200 °C and rapid loss of lubricity upon short-term sliding at 300 °C [5]. Erdemiret and coworkers further reported that DLC’s wear rate at 250 °C (1.39 × 10⁻⁶ mm³ N⁻¹ m⁻¹) is orders of magnitude higher than at room temperature (0.000186 × 10⁻⁶ mm³ N⁻¹ m⁻¹) [6]. Although Liu et al. demonstrated relatively stable wear behavior up to 200 °C in humid air against Al₂O₃, the films suffered catastrophic failure by graphitization above 300 °C despite a friction coefficient reduction from 0.15 to as low as 0.02, indicating a severely limited high-temperature lifetime [7].

Elemental doping (e.g., F [8], Si [9], Zr [10], Cr [11,12], W [13], Mo [14]) has proven effective at enhancing both mechanical properties and high-temperature tribological performance of amorphous carbon. Silicon, in particular, offers unique benefits: the longer Si-C bond relative to C-C relieves intrinsic compressive stress [15], preferentially increases the sp³-bonded fraction, and reduces sp² cluster size, thereby fortifying the film matrix [16]. HiPIMS-deposited Si-doped a-C films reported by Zhang et al. showed elevated sp³ content and reduced residual stress due to Si substitution in sp² sites [17]. Previous studies reported that Si doping levels of approximately 2–4 at.% could significantly reduce the wear rate of a-C:H films in the 10⁻⁷ mm³/N·m order of magnitude and lower the friction coefficient to values as low as 0.04 at elevated temperatures [18,19], due to suppressed oxidation and the formation of Si-containing transfer layers [18,19,20,21].

Despite extensive efforts, a comprehensive understanding of how varying Si content influences the microstructure and high-temperature tribological behavior of hydrogenated amorphous carbon films remains limited. In this study, we systematically investigate a-C:H:Si films prepared by DC magnetron sputtering with Si contents ranging from 0 to 14.7 at.%, focusing on their structural evolution, mechanical performance, and friction behavior at room temperature, 300 °C, and 400 °C. The objective of this work is to elucidate the role of Si doping in enhancing the high-temperature durability of a-C:H films and to establish correlations between microstructure and tribological response, which have not been comprehensively addressed in prior literature.

2. Experimental

2.1. Film Deposition

In this work, a-C:H and Si-doped a-C:H (a-C:H:Si) films were deposited onto 9Cr18 stainless-steel substrates (for friction tests) and Silicon wafers (for SEM, Raman, TEM, and hardness tests) using a dual-source DC magnetron sputtering system (Figure 1). Prior to insertion into the vacuum chamber, all substrates underwent a rigorous cleaning protocol to ensure surface cleanliness and promote strong film adhesion. First, the 9Cr18 discs were mechanically polished with successive grades of SiC abrasive paper, then mirror-finished. They were subsequently ultrasonically degreased in metal-detergent solution, rinsed in anhydrous ethanol for 20 min, and finally in deionized water for another 20 min. After each bath, the substrates were blow-dried with hot air and loaded into the sputtering chamber. Silicon wafers were cleaned in an identical fashion (40 min in ethanol, 20 min in deionized water) to monitor film thickness and composition. Once the base pressure in the chamber fell below 5 × 10⁻³ Pa, the sputtering targets were pre-cleaned by Ar plasma. The substrates were then etched in situ by Ar⁺ ion bombardment at a flow rate of 500 sccm (99.999% Ar) under −1000 V bias for 20 min to remove any residual oxides and adsorbates. During deposition, a rectangular graphite target (300 mm × 75 mm × 6 mm) and a silicon target of identical size were powered independently by two DC magnetrons. The substrates were aligned parallel to the targets. The reactive gas mixture of Ar and C₂H₂ was maintained at a flow ratio of 4:1 (92 sccm Ar, 23 sccm C₂H₂). The graphite target power was fixed at 1000 W, while the substrate bias was fixed at −100 V. Si incorporation was controlled by varying the Si-target power between 50, 100, and 200 W. A thin Si interlayer was deposited to improve the adhesion between the substrate and films. The distance between the target and substrates was approximately 80 mm. All films were grown for 120 min. The deposition parameters of films are provided in

Table 1. Parameters of film deposition.

Parameters	a-C:H:Si
Base pressure (Pa)	5 × 10−3
Working pressure (Pa)	~0.6
Ar/C2H2 ratio	4:1
Working temperature (°C)	120
Bias voltage (V) Substrate rotation (r/min)	−100 3
Deposition time (min)	120
Carbon target power (W)	1000
Silicon target power (W)	0/50/100/200

.

2.2. Film Characterization

The as-deposited films’ surface morphology, cross-sectional thickness, and elemental distributions were examined using a field-emission scanning electron microscope (FE-SEM; TESCAN MIRA4, Brno, Czechia) equipped with an energy-dispersive X-ray spectrometer (EDS). Film bonding configurations, both before and after high-temperature tribological testing, were probed by Raman spectroscopy (Renishaw InVia, λ = 532 nm, Wotton-under-Edge, UK). Detailed microstructural information was obtained by high-resolution transmission electron microscopy (HRTEM; Talos F200X, Thermo Fisher Scientific, Waltham, MA, USA). Mechanical properties were characterized via Knoop microhardness testing (MH-5LD, Shanghai Optical Instrument Factory, Shanghai, China) under a load of 25 gf and a dwell time of 15 s. To minimize measurement scatter, at least five indentations were made at different locations on each sample; reported values thus represent the composite hardness of the coating and substrate.

Room-temperature tribological performance was evaluated on an A & P TRB³ ball-on-disk tribometer. A 9Cr18 stainless-steel ball (6.35 mm in diameter) served as the counterpart material. The wear test was conducted under a normal load of 5 N at 400 rpm and a track radius of 6 mm for 60 min in an ambient environment with a relative humidity of 60%. High-temperature friction tests were carried out on a UMT-3 tribometer (BrukerCETR, Billerica, MA, USA) equipped with a furnace module. Tests were performed at 300 °C and 400 °C under a 5 N load, 200 rpm (track radius 5 mm), using a 6.35 mm Al₂O₃ ceramic ball as the counterpart. Wear-track morphology and worn volume were measured by optical metallography and white-light interferometry. Wear rates were calculated according to the Archard equation:

k = V/F·S,

(1)

where k is the wear rate (mm³·N⁻¹·m⁻¹), V the worn volume (mm³), F the normal load (N), and S the sliding distance (m) [22].

3. Results and Discussion

Figure 2a presents the elemental composition of a-C:H:Si films deposited at various target powers of 50 W, 100 W, and 200 W, as determined by EDS. Carbon originates from both the C₂H₂ precursor gas and the graphite sputtering target, silicon from the Si target, and oxygen from residual O₂ in the deposition chamber as well as post-deposition atmospheric adsorption. Since hydrogen cannot be quantified by EDS, the reported atom fractions are normalized to the sum of C + Si + O. As the Si target power increases from 50 W to 200 W, the Si content rises from 4.6 at.% to 14.7 at.%, while C correspondingly decreases from 89.8 at.% to 83.4 at.%. The increase in Si content is attributed to the deeper sputtering of Si atoms at higher power, yielding a greater silicon flux to the growing film. Figure 2b shows the corresponding thicknesses and deposition rates of films. Compared with the undoped a-C:H film, all a-C:H:Si coatings exhibit lower thicknesses and deposition rates under identical sputtering conditions. However, as the Si target power increases, both thickness and deposition rate increase gradually. This behavior arises from a balance between two competing effects: on one hand, energetic Si ions under negative substrate bias can sputter loosely bound hydrocarbon species from the film’s surface, reducing film growth; on the other hand, higher Si target power enhances plasma density and precursor ionization, which promotes film growth. At low Si power, the sputtering (etching) effect dominates, whereas at higher power, the increased ionization and film-forming flux prevail, resulting in thicker films and faster deposition.

Surface and cross-sectional SEM micrographs of the as-deposited films are presented in Figure 3. All films exhibit a compact columnar structure without apparent surface cracks, delamination, or large voids. The film surfaces are characterized by a typical nodular (spherical-cluster) morphology. Compared to the undoped a-C:H coating, the a-C:H:Si films display noticeably larger, more three-dimensional clusters, whose size increases with Si content. Cross-section images further reveal that the columnar structure of the a-C:H:Si films is significantly more compact than that of a-C:H film. This densification is attributed to the increased discharge power applied to the Si target, which enhances the overall sputtering flux and promotes higher energy ion bombardment at the substrate surface, thereby improving surface mobility of adatoms and leading to a denser microstructure [23].

For further investigation of the microstructure of the film, the TEM technique was adopted. As shown in Figure 4a, the low-magnification cross-sectional TEM image of the a-C:H:Si (8.3 at.% Si) film reveals two distinct layers: a thin interfacial transition zone and the overlying a-C:H:Si film. The interface is continuous and defect-free, with no evidence of delamination or cracking. Figure 4b presents elemental-mapping EDS images, demonstrating a uniform spatial distribution of C and Si throughout the film; the modest O signal arises from residual O₂ in the deposition chamber. The high-resolution TEM micrograph in Figure 4c displays a featureless, glassy morphology characteristic of amorphous materials. Consistently, the selected-area electron diffraction (SAED) pattern in Figure 4d exhibits only diffuse halos and no discrete diffraction rings, confirming the fully amorphous nature of the a-C:H:Si film, as previously reported [17,24]. Additionally, the appearance of short-range, curved chain-like motifs within the amorphous matrix indicates a substantial fraction of sp²-C bonds [25].

Figure 5 presents the Raman spectra of the a-C:H:Si films. Each spectrum exhibits a broad envelope around 1500 cm⁻¹ that can be deconvoluted into two Gaussian components: The D peak near 1350 cm⁻¹, arising from the stretching vibration of sp²-C bond sites in rings and chains, and the G peak near 1580 cm⁻¹, corresponding to the in-plane breathing mode of sp²-C bond in graphitic rings [26,27] (Figure 5a). Three key parameters are used to quantify structural disorder: the G-peak full width at half maximum (FWHM), which reflects bond-angle and bond-length distortions; the G-peak position, which correlates with the size and topology of sp² clusters; and the intensity ratio I_D/I_G, which indicates the relative abundance of ring-like versus chain-like sp²-C configurations, since a low I_D/I_G suggests predominantly chain-bound sp² carbon with limited π-electron delocalization in rings [28].

As Si content increases, the G-peak FWHM first decreases from 158.94 cm⁻¹ to 148.79 cm⁻¹ and then rises to 156.33 cm⁻¹, indicating an initial reduction in bond disorder followed by a slight reintroduction of structural distortion at the highest Si level. The G-peak position shifts upward from 1545.75 cm⁻¹ to 1552.52 cm⁻¹ before decreasing to 1548.19 cm⁻¹, suggesting a nonmonotonic variation in sp² cluster size and shape distribution. Likewise, I_D/I_G increases from 1.117 to 1.367, then decreases to 1.250, implying that moderate Si incorporation promotes the formation of larger graphitic clusters, whereas excessive Si begins to disrupt ring-like order. Collectively, these trends demonstrate that Si doping exerts a complex influence on the short- and medium-range ordering of the amorphous carbon network.

Figure 6 illustrates the Knoop microhardness of a-C:H:Si films with various Si contents. The undoped a-C:H coating exhibits a hardness of 557.0 ± 23.1 HK_0.025. With 4.6 at.% Si, the hardness increases to 673.3 ± 23.0 HK_0.025, reaching a maximum of 864.0 ± 42.1 HK_0.025 at 8.3 at.% Si. Further increasing the Si content beyond this level results in a gradual decline in hardness. The initial hardness enhancement is attributed to Si’s preferential bonding with sp³-hybridized carbon, which raises the overall sp³ fraction and thus reinforces the film’s resistance to indentation. However, excessive Si doping leads to increased free volume and progressive stress relaxation within the amorphous matrix, both of which contribute to a reduction in film hardness at higher Si concentrations [16,17].

Figure 7a shows the friction coefficient of a-C:H:Si films with varying Si contents, measured at room temperature. All films exhibit an initial running-in period of approximately 400 s, during which the friction coefficient stabilizes near ~0.12, followed by a gradual, slight increase over the remainder of the 60 min test. Figure 7b summarizes the average friction coefficients and wear rates extracted from these tests. Incorporation of Si causes a modest increase in the mean friction coefficient, which rises from 0.116 for the undoped a-C:H film to a maximum of 0.126 at 8.3 at.% Si, then decreases slightly at higher Si levels. In contrast, the wear rate of film decreases initially with Si addition, reaching a minimum value of 0.78 × 10⁻⁷ mm³·N⁻¹·m⁻¹ at 8.3 at.% Si. Then increasing to 0.99 × 10⁻⁷ mm³·N⁻¹·m⁻¹ at 14.7 at.% Si. The combination of the highest hardness and the lowest wear rate was observed at 8.3 at.% Si confirms that hardness enhancement effectively mitigates material removal. Overall, Si doping exerts only a marginal influence on the tribological performance under ambient conditions.

Optical-microscope images of the wear tracks produced on a-C:H:Si films with varying Si contents under ambient conditions are shown in Figure 8. All films display well-defined wear scars, together with clearly visible transfer films and accumulated wear debris. The 8.3 at.% Si film exhibits the narrowest wear track, approximately at 132 µm in width, whereas the 4.6 at.% and 14.7 at.% Si films show somewhat wider tracks of 143 µm and 145 µm, respectively. The undoped a-C:H film endures the greatest damage, with a maximum track width of 153 µm. At low or zero Si content, insufficient formation of a lubricious C-Si-O transfer layer leads to more severe material removal; conversely, at high Si contents, in-situ generation of hard SiC particles during sliding accelerates abrasive wear [17].

High-temperature friction tests on a-C:H:Si films with different Si contents are shown in Figure 9. The tribological properties of films tested at high temperatures show more obvious differences compared with those tested at room temperature. The friction coefficient at 300 °C is shown in Figure 9a. Each film reached a relatively stable value after a period of the running-in stage. The friction coefficient of the film with 8.3 at.% Si began to stabilize at about 680 s and maintained an ultra-low friction coefficient of 0.01 all the time. In contrast, the friction coefficients of the undoped a-C:H film and other films fluctuated around 0.1 after stabilizing. Moreover, with the increase of Si content, the time required for the running-in stage also increases, from about 300 s at 4.6 at.% to about 900 s at 14.8 at.%. The running-in stage time of a-C:H film is comparable to the doping content of 8.3 at.% Si, which indicates that low-doping Si content is conducive to rapid running-in. The wear rates of each film at 300 °C are shown in Figure 9b. The a-C:H:Si film with 4.6 at.% Si content has the highest wear rate, which is 5.3 × 10⁻⁶ mm³ N⁻¹·m⁻¹, and the wear rate with 14.8 at.% Si content is 2.03 × 10⁻⁶ mm³ N⁻¹·m⁻¹. The a-C:H film has the lowest wear rate of 0.96 × 10⁻⁶ mm³ N⁻¹·m⁻¹. Notably, the a-C:H:Si film with a Si content of 8.3 at.% exhibits the lowest coefficient of friction while having a relatively low wear rate of 1.2 × 10⁻⁶ mm³ N⁻¹·m⁻¹.

The friction coefficient of films at 400 °C is shown in Figure 9c. The a-C:H films have the highest initial friction coefficient and the longest running-in period. After 230 s of running-in, the friction coefficients of a-C:H:Si films with Si contents of 4.6 at. % and 14.8 at.% remained at 0.01, which is the same as the a-C:H films during the stable stage. Meanwhile, the friction coefficient of a-C:H:Si films with 8.3 at.% Si gradually increased at the beginning. It reached a maximum of 0.068 at about 220 s, then gradually decreased, dropped below 0.01 at 870 s, and continued to decrease slowly, reducing to about 0.0034 at the end of the friction test. During the high-temperature friction process, the friction coefficient may occasionally show a sharp peak shape, and this phenomenon can be regarded as caused by the instability of the transfer film [29]. The wear rate at 400 °C is shown in Figure 9d. The wear rate of the a-C:H film is 4.6 × 10⁻⁶ mm³ N⁻¹·m⁻¹, and the wear rate of 4.6 at.% Si film is 3.32 × 10⁻⁶ mm³ N⁻¹·m⁻¹. The film wear rates of the films containing 8.3 at.% and 14.8 at.% Si are comparable, which are 2.52 × 10⁻⁶ mm³ N⁻¹·m⁻¹ and 2.35 × 10⁻⁶ mm³ N⁻¹·m⁻¹, respectively. It can thus be concluded that at 400 °C, the incorporation of silicon effectively reduces the wear rate of the amorphous carbon films, with the anti-wear performance becoming increasingly pronounced as the Si doping concentration rises.

Compared with the room temperature environment, the friction coefficient has been significantly reduced in the high-temperature environments of 300 °C and 400 °C, but the wear rate has increased. The reduction in friction coefficient is primarily driven by thermally induced graphitization and the formation of highly lubricious transfer films, which dramatically lower shear strength at the sliding interface. Meanwhile, the structural transformation of the a-C:H:Si thin film caused by high temperature leads to a decrease in mechanical properties, thereby resulting in an increase in the wear rate.

The wear track morphologies of the films after friction testing at 300 °C and 400 °C are shown in Figure 10. The depths of all wear tracks remained below the thickness of the deposited films, indicating that the films were not worn through during the tests. All films exhibited relatively narrow wear tracks. At 300 °C, the edges of the wear scars on both the undoped a-C:H film and the a-C:H:Si film with 4.6 at.% Si was smooth, whereas other samples displayed pronounced plastic deformation and raised edges. As the temperature increased to 400 °C, noticeable changes in wear morphology were observed. The wear track edges became flatter, but distinct ploughing grooves appeared within the wear tracks. This suggests a transition in the dominant wear mechanism from adhesive wear at 300 °C to abrasive wear at 400 °C [30].

At a test temperature of 400 °C, the a-C:H:Si film with 8.3 at.% Si exhibited outstanding tribological performance, characterized by both a low coefficient of friction and a low wear rate. As illustrated in Figure 11, the wear scar of the friction ball and wear track of the film were analyzed by SEM, EDS, and Raman spectroscopy to investigate the underlying mechanisms. As shown in Figure 11a, debris accumulation was observed near the wear scar, where localized enrichment of both carbon and silicon occurred. Similarly, Figure 11c reveals the presence of wear debris around the wear track, with comparable elemental distributions. The wear track itself appears smooth, suggesting shallow wear depth, consistent with results obtained from optical microscopy and white light interferometry. Elemental mapping further confirmed that the worn regions (both scar and track) were rich in Si, whereas the surrounding wear debris was predominantly composed of carbon. The superior performance of this film can be attributed to the protective role of silicon during high-temperature sliding. Silicon has a lower work function (4.85 eV) compared to carbon (5.0 eV), indicating a higher tendency to lose electrons and undergo oxidation. Consequently, Si atoms preferentially react with environmental oxygen or moisture, forming a passivating layer that protects the underlying carbon network from degradation. This contributes to the film’s enhanced lubricity and stability in oxidative atmospheres [25].

Raman spectra of the unworn surface, wear track, and wear scar (Figure 11b) indicate the formation of a tribolayer during sliding. Specifically, the wear scar region shows a pronounced amorphous carbon signal, suggesting transfer film formation in the contact zone. Compared with the unworn area, both the D peak near 1350 cm⁻¹ becomes sharper and more intense, and the G peak near 1560 cm⁻¹ increases in intensity and shifts to a higher wavenumber. These spectral changes suggest a structural transformation from sp³-C to sp²-C during the friction process, indicative of thermally induced graphitization. The growth of sp² clusters under high temperature and stress may further promote the formation of graphene-like structures, enhancing the lubricating behavior of the tribolayer [31].

The above results demonstrate that silicon doping can effectively reduce the high-temperature friction coefficient of amorphous carbon (a-C:H) films and even exhibit unexpectedly outstanding performance when optimally doped. As illustrated in Figure 12, the Si-doped a-C:H:Si film possesses a coupled structure in which silicon atoms are incorporated into the sp²/sp³ amorphous carbon matrix. Notably, Si atoms preferentially bond with sp³-hybridized carbon atoms to form Si-C bonds, while the sp²-C network remains largely unaffected.

During sliding contact between the a-C:H:Si film and the counterface ball, the interfacial region undergoes complex processes including carbon dissociation, surface reactions, and interfacial interactions. These processes result in the adsorption and attachment of oxygen species, water molecules, and free amorphous carbon fragments to the counterface, ultimately forming a carbon-rich transfer layer. In contrast to conventional hard carbide-forming dopants that tend to disrupt the carbon matrix and elevate internal stresses—promoting delamination or acting as abrasive particles [32]—Si plays a more beneficial role.

Due to its lower work function, Si is more readily oxidized than carbon. During friction, this preferential oxidation of Si helps suppress the oxidation of the graphitic transfer layer, preserving its lubricating function. Furthermore, the generated oxides are often encapsulated within the graphitic matrix, preventing their direct participation in tribological interactions, thereby enhancing the stability and durability of the lubricating film. At the tribological interface, which is subject to both high temperature and mechanical stress, Si atoms can also react with water molecules in air, forming bonds with -OH or -H groups [33,34]. These terminations serve to passivate dangling bonds in a manner similar to hydrogen atoms already present in the film. The resulting H-H interactions across the sliding interface are associated with weak van der Waals forces, which contribute significantly to the low-friction behavior.

Additionally, under tribo-thermal activation, a structural reorganization occurs involving the transformation from sp³-C to sp²-C configurations. This graphitization process may lead to the formation of graphene-like domains [35,36,37], which are characterized by interlayer weak π*-π* interactions, another critical factor enabling superlubricity. Therefore, the low-friction behavior of Si-doped a-C:H:Si films at elevated temperatures can be attributed to the synergistic effect of interfacial H-H interactions and π*-π* stacking between graphene-like lamellae. Together, these mechanisms contribute to the formation of a stable, low-shear tribological interface that sustains excellent lubricating performance under harsh thermal–mechanical conditions [38].

4. Conclusions

This study addresses the temperature sensitivity limitation of hydrogenated amorphous carbon (a-C:H) films by exploring the effects of silicon doping on their high-temperature tribological performance. A series of a-C:H:Si films with varying Si content were deposited via direct current magnetron sputtering, and their microstructure, mechanical properties, and tribological behavior were systematically investigated at elevated temperatures (300 °C and 400 °C). The findings demonstrate that moderate Si doping significantly enhances the wear resistance and tribological performance of a-C:H films. Notably, the film with 8.3 at.% Si exhibited an ultralow friction coefficient of 0.0034 at 400 °C, which is attributed to the synergistic effects of a carbon-rich transfer layer formation, preferential Si oxidation, and tribo-induced phase transformation to graphitic structures at the sliding interface. The results also reveal that Si content influences film hardness, with the maximum hardness observed at 8.3 at.% Si, followed by a decrease at higher doping levels. These findings emphasize the critical role of controlled Si doping in tailoring the structural and tribological properties of a-C:H:Si films, providing new insights for the development of high-performance films in high-temperature environments.