The Structure, Morphology, and Mechanical Properties of Ta-Hf-C Coatings Deposited by Pulsed Direct Current Reactive Magnetron Sputtering

: Ta, Hf, TaC x , HfC x , and Ta x Hf 1-x C y coatings were deposited by reactive pulsed Direct Current (DC) magnetron sputtering of Ta or Hf pure metallic targets in Ar plus CH 4 gas mixtures. The properties have been investigated as a function of the carbon content, which is tuned via the CH 4 ﬂow rate. The discharge was characterized by means of Optical Emission Spectroscopy and, in our conditions, both Ta-C and Hf-C systems seem to be weakly reactive. The structure of the as-deposited pure tantalum ﬁlm is metastable tetragonal β -Ta. The fcc-MeC x carbide phases (Me = Ta or Hf) are {111} textured at low carbon concentrations and then lose their preferred orientation for higher carbon concentrations. Transmission Electron Microscopy (TEM) analysis has highlighted the presence of an amorphous phase at higher carbon concentrations. When the carbon content increases, the coating’s morphology is ﬁrst compact-columnar and becomes glassy because of the nano-sized grains and then returns to an open columnar morphology for the higher carbon concentrations. The hardness and Young’s modulus of TaC x coatings reach 36 and 405 GPa, respectively. For HfC x coatings, these values are 29 and 318 GPa. The MeC x coating residual stresses increase with the addition of carbon (from one-hundredth of 1 MPa to 1.5 GPa approximately). Nevertheless, the columnar morphology at a high carbon content allows the residual stresses to decrease. Concerning Ta x Hf 1-x C y coatings, the structure and the microstructure analyses have revealed the creation of a nanostructured coating, with the formation of an fcc superlattice. The hardness is relatively constant independently of the chemical composition (22 GPa). The residual stress was strongly reduced compared to that of binary carbides coatings, due to the rotation of substrates.


Introduction
Nowadays, surface treatments are widely used at industrial scale to improve the lifetime of many components or tools in many applications fields [1]. Surface functionalities can be addressed by a wide coatings obtained by magnetron sputtering [27]. Due to the lack of literature on Ta x Hf 1-x C y coatings, it is necessary to control the process with binary carbides, in order to determine the experimental conditions necessary to synthesize ternary alloys. In this work, we, for the first time, synthesized Ta x Hf 1-x C y coatings by reactive magnetron sputtering with a characterization of the structure, the microstructure, the morphology, and the mechanical properties (hardness, Young's modulus, and residual stress).

Deposition
Ta, TaC x , Hf, HfC x , and Ta x Hf 1x C y thin films were deposited by pulsed DC reactive magnetron sputtering of disc targets (200 mm in diameter) in a 100 L Alcatel vacuum chamber in pure Ar or in Ar-CH 4 reactive discharges. The deposition chamber was pumped down to 10 −4 Pa before each run. Substrates (Si wafers, glass slides, and M2 steel) were degreased in acetone by an ultrasonic bath and then successively rinsed by ethanol and dried with warm air. The target-to-substrate distance was 13 cm. A substrate etching was performed in pure argon before each deposition stage with biasing at −600 V for 30 min. Ta (99.9% purity) and Hf (99.9% purity) targets were also etched in pure argon at 2 A. Solvix (50 kHz) and Advanced Energy MDX pulsed DC (50 kHz) power supplies powered the substrate holder and targets, respectively. The argon flow rate (DAr) was kept constant at 30 sccm for all experiments, and the CH 4 flow rate (DCH 4 ) increased from 1 to 6 sccm for the coating's synthesis in reactive conditions. The working pressure was 1 Pa, and the substrate holder was kept at floating potential for all experiments.
Optical Emission Spectroscopy (OES) measurements were carried out with a triple optical fiber placed at mid-distance between the target and the substrate. The Ar 0 , Ta 0 , Hf 0 , and H wavelengths used to calculate the I Ta 0 I Ar 0 , I H f 0 I Ar 0 , and I Hα I Ar 0 ratios ((I) being the peak intensity) were 750.4, 696.9, 706.4, and 656.3 nm respectively. Experimental conditions of the Ta, TaC x , Hf, and HfC x coatings synthesized by pulsed DC reactive and non-reactive magnetron sputtering are summarized in Table 1. Concerning Ta x Hf 1-x C y coatings, the discharge current on each (Ta and Hf) target was tuned between 1 and 3 A by maintaining the total current dissipated on both targets at 4 A. The CH 4 flow rate was 8 sccm (indeed, the pure TaC and HfC phases were obtained with 4 sccm of CH 4 , as shown in Section 3.2). The substrate holder rotation was set to 10 rpm. Table 2 summarizes the experimental conditions of the ternary alloy synthesis.

Characterizations
The coatings' thickness was measured by an Altisurf 500 profilometer, and it remained between 2 and 3 µm. The coatings' cross-sectional morphology was observed with an FEG (Field Emission Gun), Hitachi SU 8030 SEM.
Young's modulus and hardness were obtained by means of a Nano Indenter XP, MTS Systems Corporation, with a Continuous Stiffness Measurement (CSM) option. The indenter is a three-side pyramidal diamond tip (Berkovitch). The values of hardness and Young's modulus were measured from an average of 10 indents. The depth of penetration was 200 nm.
The coating's composition was estimated via XPS measurement by calculating (H f 4 f +C1s) area peak ratios. Ar + etching was performed to eliminate the surface's contamination layer (60 s for TaC x coatings and 2 min for HfC x coatings). A resolution of 0.1 eV was used with K (Thermo) fitted with a monochromatic Al K α X-ray source (spot size: 400 µm) and a flood gun for static charge compensation.
The transmission electron microscope (TEM) was a JEM -ARM 200F Cold FEG TEM/STEM operating at 200 kV and equipped with a spherical aberration (Cs) probe and image correctors (a point resolution of 0.12 nm in TEM mode and 0.078 nm in STEM mode).
The residual stresses are measured from an iron-substrate curvature before and after deposition, using an Altisurf profilometer. The stress was then calculated using Stoney's equation [28], as in [29].

Reactive Discharge Characterization
The Ta-Ar and Hf-Ar discharges seem to have a similar behavior when the CH 4 flow rate increases.
A continuous decrease of ITa 0 IAr 0 ( Figure 1a) and IAr 0 (Figure 1c) ratios was observed with an increasing C 2 H 2 flow rate. These evolutions are due to the decrease in the metallic species density, which can first be explained by the target's poisoning associated with the compound's sputtering yield, which is lower than that of the metal. The decrease in the argon partial pressure when CH 4 is introduced at a constant total pressure also leads to a decrease in the sputtering rate. The electronic energy was also shared between the Ar excitation/ionization and the CH 4 dissociation, leading to an Ar + rarefaction.
The IH α IAr 0 (Figure 1b,d corresponding to Ta-Ar-CH 4 and Hf-Ar-CH 4 discharges) ratio increases when CH 4 is introduced into the chamber. This behavior is due to the increase of the H radicals concentration following the different dissociation reactions of CH 4 molecules.

Figure 1.
Evolution of (a) the 0 0 ratio and (b) the 0 ratio according to the CH4 flow rate (from 1 to 6 sccm) within Ta-Ar discharges and the evolution of (c) the 0 0 ratio and (d) the 0 ratio according to the CH4 flow rate (from 1 to 6 sccm) within Hf-Ar discharges. For both Ta-Ar and Hf-Ar discharges, the total pressure and discharge current are kept constant at 1 Pa and 2 A, respectively Figure 2 shows the evolution of the TaCx ( Figure 2a) and HfCx ( Figure 2b) coatings' growth rate when the inlet CH4 flow rate increases. The growth rate drops as DCH4 increases for both TaCx and HfCx coatings (from 4 to 1.5 µm/h for TaCx and from 5 to 2 µm/h for HfCx). This is a common phenomenon due to the sputtering yield of the compound lower than that of the metal. This also confirms the target's poisoning as seen in Figure 1. Moreover, the Ar partial pressure declines when CH4 is introduced because the total pressure is kept constant at 1 Pa. Hence, the sputtering rate decreases, which is consistent with the previous observations ( Figure 1).

Composition, Structure, and Microstructure
The measured atomic composition of TaCx and HfCx coatings according to the CH4 flow rates are summarized in Table 3. The carbon content increases as the CH4 flow rate increases for both TaCx and HfCx coatings. Nevertheless, the composition of TaCx coatings seems to remain constant from 4 to 6 sccm of CH4 while keeping a stoichiometric compound without forming an over-stoichiometric coating. IAr 0 ratio and (b) the IH α IAr 0 ratio according to the CH 4 flow rate (from 1 to 6 sccm) within Ta-Ar discharges and the evolution of (c) the IH f 0 IAr 0 ratio and (d) the IH α IAr 0 ratio according to the CH 4 flow rate (from 1 to 6 sccm) within Hf-Ar discharges. For both Ta-Ar and Hf-Ar discharges, the total pressure and discharge current are kept constant at 1 Pa and 2 A, respectively Figure 2 shows the evolution of the TaC x ( Figure 2a) and HfC x (Figure 2b) coatings' growth rate when the inlet CH 4 flow rate increases. The growth rate drops as DCH 4 increases for both TaC x and HfC x coatings (from 4 to 1.5 µm/h for TaC x and from 5 to 2 µm/h for HfC x ). This is a common phenomenon due to the sputtering yield of the compound lower than that of the metal. This also confirms the target's poisoning as seen in Figure 1. Moreover, the Ar partial pressure declines when CH 4 is introduced because the total pressure is kept constant at 1 Pa. Hence, the sputtering rate decreases, which is consistent with the previous observations ( Figure 1).  (Figure 2b) coatings' growth rate when the inlet CH4 flow rate increases. The growth rate drops as DCH4 increases for both TaCx and HfCx coatings (from 4 to 1.5 µ m/h for TaCx and from 5 to 2 µ m/h for HfCx). This is a common phenomenon due to the sputtering yield of the compound lower than that of the metal. This also confirms the target's poisoning as seen in Figure 1. Moreover, the Ar partial pressure declines when CH4 is introduced because the total pressure is kept constant at 1 Pa. Hence, the sputtering rate decreases, which is consistent with the previous observations ( Figure 1).

Composition, Structure, and Microstructure
The measured atomic composition of TaCx and HfCx coatings according to the CH4 flow rates are summarized in Table 3. The carbon content increases as the CH4 flow rate increases for both TaCx and HfCx coatings. Nevertheless, the composition of TaCx coatings seems to remain constant from 4 to 6 sccm of CH4 while keeping a stoichiometric compound without forming an over-stoichiometric coating.

Composition, Structure, and Microstructure
The measured atomic composition of TaC x and HfC x coatings according to the CH 4 flow rates are summarized in Table 3. The carbon content increases as the CH 4 flow rate increases for both TaC x and HfC x coatings. Nevertheless, the composition of TaC x coatings seems to remain constant from 4 to 6 sccm of CH 4 while keeping a stoichiometric compound without forming an over-stoichiometric coating. The evolution of TaC x and HfC x coatings structure as a function of the CH 4 flow rate is shown in Figures 3 and 4, respectively. When pure Ta is sputtered (Figure 3a), a β-Ta tetragonal structure textured {002}, which is not referenced in the pressure-temperature diagram of Ta [30,31], is obtained. The stable phase α-Ta has a body-centered cubic (bcc) structure. The β-Ta metastable structure is commonly obtained in sputtering processes [32][33][34][35][36]. Collin et al. have proposed a theory explaining the precipitation of this phase [36] by thermodynamic calculations. TEM imagery has revealed the presence of an amorphous tantalum phase (a-TaSi) at the substrate/β-Ta interface. Thermodynamic calculations have shown that growth on the amorphous a-TaSi phase of the β-Ta phase is more favorable than that of the α-Ta phase. A growth with planes (002) parallel to the substrate's surface, characterized by a strong {002} texture in the XRD pattern (Figure 3a), is an expected result because the film grows according to the planes with the lowest surface energy. However, Bernoulli et al. have proposed another hypothesis based on the observed results in the literature [37]. They assume that the chemical composition of the substrate's surface (the presence of a native oxide, for example) has a strong influence on the Ta germination and growth.
When 1 sccm of CH 4 is introduced (Figure 3b), the α-Ta bcc phase with a {110} texture is stabilized. This could be due to the insertion of carbon in the lattice, which creates a local compressive stress, allowing a tetragonal to bcc transition via lattice contraction. Furthermore, the shortest Ta-Ta distance in the β-Ta tetragonal cell is smaller than that of the α-Ta bcc cell (i.e. 0.265 nm against 0.286 nm, respectively) [38]. It should be easier for carbon to be inserted in the bcc structure, and this could be a reason why tetragonal β Ta is transformed into bcc α-Ta. A fcc nanocristalined TaC structure with a {111} texture is obtained for 2 sccm of CH 4 . The texture is lost when DCH 4 increases until 6 sccm. The loss of preferential orientation can be explained by a reduced adatom mobility associated with alower bombardment of the growing film.
The evolution of TaC 0.76 microstructure across the film by TEM is shown in Figure 5. The Selected Area Electron Diffraction (SAED) patterns confirm the presence of a TaC fcc crystalline phase. However, both the microstructure and crystalline orientation evolve following the selected zone within the film. Near the substrate/film's interface (Figure 5a,b), a polycrystalline film growth with equiaxed grain oriented following the <111> direction is observed. When the coating becomes thicker (Figure 5c-f), a columnar microstructure appears. The <111> preferential orientation is also reinforced. This is a common phenomenon of thin film growth [39], and this is consistent with the XRD pattern in Figure 3d. Twin grains can also be observed (Figure 5c).   The evolution of TaC0.76 microstructure across the film by TEM is shown in Figure 5 Area Electron Diffraction (SAED) patterns confirm the presence of a TaC fcc cryst However, both the microstructure and crystalline orientation evolve following the s within the film. Near the substrate/film's interface (Figure 5a,b), a polycrystalline film equiaxed grain oriented following the <111> direction is observed. When the coating bec (Figure 5c-f), a columnar microstructure appears. The <111> preferential orient  The evolution of the HfCx coatings structure is similar to that of the TaCx coatings. Nevertheless, the stable hcp α-Hf metallic structure with a {002} texture is first obtained when the Hf target is sputtered in pure argon (Figure 4a). When 1 sccm of CH4 is introduced (Figure 4b), the α-Hf phase remains hcp, but the insertion of carbon leads to the peak shift toward smaller angles because of the lattice's distortion associated with the carbon atom's insertion. Figure 6a confirms the coating's hcp structure. Indeed, the FFT carried out with an HRTEM image, coupled to the theoretical pattern with the [100] zone axis, calculated using jems software (Figure 6b), helps to highlight this structure.  The evolution of the HfC x coatings structure is similar to that of the TaC x coatings. Nevertheless, the stable hcp α-Hf metallic structure with a {002} texture is first obtained when the Hf target is sputtered in pure argon (Figure 4a). When 1 sccm of CH 4 is introduced (Figure 4b), the α-Hf phase remains hcp, but the insertion of carbon leads to the peak shift toward smaller angles because of the lattice's distortion associated with the carbon atom's insertion. Figure 6a confirms the coating's hcp structure. Indeed, the FFT carried out with an HRTEM image, coupled to the theoretical pattern with the [100] zone axis, calculated using jems software (Figure 6b), helps to highlight this structure.
Coatings 2020, 10, x FOR PEER REVIEW 9 of 17 The evolution of the HfCx coatings structure is similar to that of the TaCx coatings. Nevertheless, the stable hcp α-Hf metallic structure with a {002} texture is first obtained when the Hf target is sputtered in pure argon (Figure 4a). When 1 sccm of CH4 is introduced (Figure 4b), the α-Hf phase remains hcp, but the insertion of carbon leads to the peak shift toward smaller angles because of the lattice's distortion associated with the carbon atom's insertion. Figure 6a confirms the coating's hcp structure. Indeed, the FFT carried out with an HRTEM image, coupled to the theoretical pattern with the [100] zone axis, calculated using jems software (Figure 6b), helps to highlight this structure. From 2 to 6 sccm of CH4, a NaCl fcc structure of the carbide phase (HfC) is obtained where the {111} texture is also progressively lost. TEM microscopy (Figure 7) also shows an fcc structure. From 2 to 6 sccm of CH 4 , a NaCl fcc structure of the carbide phase (HfC) is obtained where the {111} texture is also progressively lost. TEM microscopy (Figure 7) also shows an fcc structure. Nevertheless, a halo on the FFT, Figure 7c, reveals the presence of an amorphous phase. This is probably an a-C:H phase commonly encountered in transition metal carbide coatings [26]. Concerning TaxHf1-xCy coatings, Figure 8 shows the microstructure of a coating where 3A and 1A were applied to the Ta and Hf targets, respectively. It reveals a nanostructured coating where the period is 2.2 nm. If the period is predicted with Equation (1), the experimental period is consistent with the calculated period: where (nm) is the period, en (nm) is the coating's thickness, n (rpm) is the rotation speed, and t (min) is the deposition time. The calculated period is about 2 nm in our conditions. The XRD patterns ( Figure 9) show a crystalline monophasic fcc material, which is a result coming from the superlattice TaC/HfC coating. The chemical composition does not influence the coatings' crystallographic structure. Concerning Ta x Hf 1-x C y coatings, Figure 8 shows the microstructure of a coating where 3A and 1A were applied to the Ta and Hf targets, respectively. It reveals a nanostructured coating where the period is 2.2 nm. If the period is predicted with Equation (1), the experimental period is consistent with the calculated period: Concerning TaxHf1-xCy coatings, Figure 8 shows the microstructure of a coating where 3A and 1A were applied to the Ta and Hf targets, respectively. It reveals a nanostructured coating where the period is 2.2 nm. If the period is predicted with Equation (1), the experimental period is consistent with the calculated period: where (nm) is the period, en (nm) is the coating's thickness, n (rpm) is the rotation speed, and t (min) is the deposition time. The calculated period is about 2 nm in our conditions. The XRD patterns ( Figure 9) show a crystalline monophasic fcc material, which is a result coming from the superlattice TaC/HfC coating. The chemical composition does not influence the coatings' crystallographic structure. The XRD patterns ( Figure 9) show a crystalline monophasic fcc material, which is a result coming from the superlattice TaC/HfC coating. The chemical composition does not influence the coatings' crystallographic structure.

Morphology
The evolution of TaCx and HfCx coatings' morphology is shown in Figures 10 and 11, respectively. They exhibit a very similar evolution when DCH4 increases. Pure metals (Ta and Hf) have a columnar morphology, probably due to the high working pressure (1 Pa), which limits the adatom's mobility to create a compact morphology. When CH4 is introduced, the morphology loses its columnar character to become glassy. This effect can be observed in Figure 10c and Figure 11c when 2 sccm of CH4 are used. This effect must be related to the size of the grain, which might be very fine. This is consistent with XRD patterns (Figure 3c and Figure 4c) where the peaks are wide, traducing a nanocrystalline structure. The columnar growth is observed when DCH4 continues to increase due to the carbide phase crystallization and the growth of grains.

Morphology
The evolution of TaC x and HfC x coatings' morphology is shown in Figures 10 and 11, respectively. They exhibit a very similar evolution when DCH 4 increases. Pure metals (Ta and Hf) have a columnar morphology, probably due to the high working pressure (1 Pa), which limits the adatom's mobility to create a compact morphology. When CH 4 is introduced, the morphology loses its columnar character to become glassy. This effect can be observed in Figures 10c and 11c when 2 sccm of CH 4 are used. This effect must be related to the size of the grain, which might be very fine. This is consistent with XRD patterns (Figures 3c and 4c) where the peaks are wide, traducing a nanocrystalline structure. The columnar growth is observed when DCH 4 continues to increase due to the carbide phase crystallization and the growth of grains.

Morphology
The evolution of TaCx and HfCx coatings' morphology is shown in Figures 10 and 11, respectively. They exhibit a very similar evolution when DCH4 increases. Pure metals (Ta and Hf) have a columnar morphology, probably due to the high working pressure (1 Pa), which limits the adatom's mobility to create a compact morphology. When CH4 is introduced, the morphology loses its columnar character to become glassy. This effect can be observed in Figure 10c and Figure 11c when 2 sccm of CH4 are used. This effect must be related to the size of the grain, which might be very fine. This is consistent with XRD patterns (Figure 3c and Figure 4c) where the peaks are wide, traducing a nanocrystalline structure. The columnar growth is observed when DCH4 continues to increase due to the carbide phase crystallization and the growth of grains.   The evolution of the morphology of TaxHf1-xCy coatings as a function of the applied current target is shown in Figure 12. The obtained morphologies are relatively compact; nevertheless, when 2A is applied on both targets, the appearance of columns can be highlighted.

Nanoindentation
The evolution of the coating's hardness and Young's modulus is shown in Figure 13.  The evolution of the morphology of Ta x Hf 1-x C y coatings as a function of the applied current target is shown in Figure 12. The obtained morphologies are relatively compact; nevertheless, when 2A is applied on both targets, the appearance of columns can be highlighted. The evolution of the morphology of TaxHf1-xCy coatings as a function of the applied current target is shown in Figure 12. The obtained morphologies are relatively compact; nevertheless, when 2A is applied on both targets, the appearance of columns can be highlighted.

Nanoindentation
The evolution of the coating's hardness and Young's modulus is shown in Figure 13.

Nanoindentation
The evolution of the coating's hardness and Young's modulus is shown in Figure 13. For both TaC x and HfC x coatings, mechanical properties are enhanced when DCH 4 increases from 1 to 4 sccm. This should be due to carbide precipitation (TaC and HfC fcc-structured) with a higher covalent bond degree (between Ta or Hf and C) and to the film's morphology (corresponding to the observed morphologies in Section 3.4). Nevertheless, when DCH 4 = 5 and 6 sccm, both the coating's hardness and Young's modulus drop. This is a well known effect due to the presence of a carbon-based amorphous phase (a-C:H) (as can be seen in Figure 7), which is harmful to the hardness and Young's modulus, as shown in [20]. Moreover, the morphology is open-columnar for the higher CH 4 flow rates. Pure Ta and Hf reach a hardness and Young's modulus of about 17 and 213 GPa for Ta, respectively, and 8 and 156 GPa for Hf, respectively. The hardness and Young's modulus of TaC x are between 18 and 36 GPa and between 236 and 405 GPa, respectively. The hardness and Young's modulus of HfC x are between 11 and 29 GPa and between 175 and 318 GPa, respectively. This is consistent with values found in the literature and reported in Table 4.

Nanoindentation
The evolution of the coating's hardness and Young's modulus is shown in Figure 13.  The measured hardness and Young's modulus of Ta x Hf 1-x C y coatings are reported in Table 5. The measured values are consistent with other values found in the literature on bulk materials [40][41][42][43]-i.e., 23 GPa for hardness and 240 GPa for Young's modulus. Nevertheless, the composition does not influence strongly either the hardness or Young's modulus.  Figure 14 shows the evolution of the residual stresses of TaC x (Figure 14a) and HfC x (Figure 14b) coatings according to the CH 4 flow rate. The compressive stresses in both coatings rise significantly with the increase in carbon content; i.e., it increases from -350 MPa to -1750 MPa for TaC x and from -200 MPa to -1600 MPa for HfC x . This is linked to the lattice deformation due to the carbon incorporation. This is consistent with Windishmann's model, based on volumetric distortions caused by the atom's displacement [44]. Moreover, the coating compaction due to the insertion of carbon (Figures 10b and 11b) participates in this phenomenon. Indeed, the interatomic repulsive forces lead to a compressive macro-state stress in the film [45]. At a high carbon content, the compressive stresses become weaker. This is consistent with the coatings' columnar morphology (Figures 10e and  11e). According to Hoffman's model, the interaction of the columns between them generates tensile stresses [46,47]. Hence, the diminution of residual stresses within TaC x and HfC x coatings at high carbon content results from a competition between the tensile and compressive stresses. A similar effect is shown in Figure 15 for Ta x Hf 1-x C y coatings, where the residual stresses are reduced when 2A is applied on both targets, where a columnar morphology is found. The residual stresses of ternary coatings are strongly reduced compared to the binary coatings, due to the rotation of substrates, which limits their bombardment.
Coatings 2020, 10, x FOR PEER REVIEW 14 of 17 lead to a compressive macro-state stress in the film [45]. At a high carbon content, the compressive stresses become weaker. This is consistent with the coatings' columnar morphology (Figure 10e and Figure 11e). According to Hoffman's model, the interaction of the columns between them generates tensile stresses [46,47]. Hence, the diminution of residual stresses within TaCx and HfCx coatings at high carbon content results from a competition between the tensile and compressive stresses. A similar effect is shown in Figure 15 for TaxHf1-xCy coatings, where the residual stresses are reduced when 2A is applied on both targets, where a columnar morphology is found. The residual stresses of ternary coatings are strongly reduced compared to the binary coatings, due to the rotation of substrates, which limits their bombardment.

Conclusion
The properties (structure, morphology, and mechanical properties) of TaxHf1-xCy thin films synthesized by reactive magnetron sputtering (with CH4 as the reactive gas) have been evaluated.
Concerning binary coatings, the deposition rate progressively decreases as the CH4 flow rate increases due to the target's poisoning. The HfCx or TaCx coating structure can be controlled by tuning the CH4 flow rate. A β-Ta (with a tetragonal structure {002} textured) phase is firstly obtained, and the α-Ta phase (bcc structure) is stabilized when DCH4 = 1 sccm. The TaC phase crystalizes from lead to a compressive macro-state stress in the film [45]. At a high carbon content, the compressive stresses become weaker. This is consistent with the coatings' columnar morphology (Figure 10e and Figure 11e). According to Hoffman's model, the interaction of the columns between them generates tensile stresses [46,47]. Hence, the diminution of residual stresses within TaCx and HfCx coatings at high carbon content results from a competition between the tensile and compressive stresses. A similar effect is shown in Figure 15 for TaxHf1-xCy coatings, where the residual stresses are reduced when 2A is applied on both targets, where a columnar morphology is found. The residual stresses of ternary coatings are strongly reduced compared to the binary coatings, due to the rotation of substrates, which limits their bombardment.

Conclusion
The properties (structure, morphology, and mechanical properties) of TaxHf1-xCy thin films synthesized by reactive magnetron sputtering (with CH4 as the reactive gas) have been evaluated.
Concerning binary coatings, the deposition rate progressively decreases as the CH4 flow rate

Conclusions
The properties (structure, morphology, and mechanical properties) of Ta x Hf 1-x C y thin films synthesized by reactive magnetron sputtering (with CH 4 as the reactive gas) have been evaluated.
Concerning binary coatings, the deposition rate progressively decreases as the CH 4 flow rate increases due to the target's poisoning. The HfC x or TaC x coating structure can be controlled by tuning the CH 4 flow rate. A β-Ta (with a tetragonal structure {002} textured) phase is firstly obtained, and the α-Ta phase (bcc structure) is stabilized when DCH 4 = 1 sccm. The TaC phase crystalizes from 2 sccm of CH 4 with a {111} texture, which is lost when DCH 4 increases until 6 sccm. Concerning the Hf-C system, the α-Hf (the hcp structure with a {002} texture) is primarily sputtered and keeps this structure when 1 sccm of CH 4 is used. The HfC phase fcc {111} textured is synthesized at 2 sccm of CH 4 . The texture is also progressively lost as DCH 4 increases. Both Ta and Hf coatings exhibit a compact-columnar morphology. When CH 4 is introduced in the gas mixture, the coating's morphology becomes glassy due to the nano-grain sizes. However, when more reactive gas is used, an open-columnar growth is observed. TEM microscopy has allowed one to highlight the amorphous a-C:H phase when 6 sccm of CH 4 is added to the Ar-Hf discharge. The hardness and Young's modulus of TaC x coatings reach 36 and 405 GPa, respectively. Concerning HfC x coatings, the hardness and Young's modulus reach 29 and 318 GPa. Nevertheless, the mechanical properties, with a high carbon content, drop. This is assumed to be related to the presence of an amorphous a-C:H phase.
Concerning ternary coatings, nanostructured Ta x Hf 1-x C y coatings were synthesized by co-sputtering in reactive conditions. The period between TaC/HfC nanolayers was measured at 2 nm, while forming an fcc superlattice. The hardness and Young's modulus are not influenced by the coating's composition and were found to be 23 and 240 GPa, respectively. The compressive stresses around 10 MPa are relatively weak, due to the rotation of substrates.