**3. Results**

### *3.1. Elemental and Phase Composition*

The elemental composition of the coatings obtained by the EDS method is presented in Table 2. Both coatings had an almost stoichiometric structure. The amount of Si was low but proved to be su fficient for the goal of this study. Both coatings exhibited a preferred growth orientation after (111) plan (Figure 1). The TiCN peaks were located between those of TiC and TiN, indicating that the TiCN is a solid solution and consisted of a mixture of TiC and TiN, with a face-centered cubic (FCC) structure (like NaCl). The same results were also reported by other authors regarding the TiCN coatings [21]. The identification of TiCN was performed according to JCPDS no. 042-1488 (red lines in Figure 1). It can be seen that our TiCN is shifted towards a lower Bragg angle compared with the TiCN standard, which is related to the formation of more TiC phase (Figure 1). This finding is in accordance with the results reported by Wang et al. [34] regarding the TiSiCN coatings. The preferred orientation of both coatings was also confirmed by the texture coe fficient T(*hkl*), a strong texture intensity at (111) plane being found in the case of both coatings. Karlsson et al. reported that the plane with a preferred orientation parallel to the investigated surface has a texture coe fficient value higher than 1 [21].

**Table 2.** Elemental composition measured by EDS, texture coe fficient T(*hkl*), crystallite size *d* and strain ε determined by the Williamson–Hall plot method.


**Figure 1.** X-ray diffraction patterns of the investigated coatings (S-substrate) as well as of the TiCN standard (JCPDS 042-1488).

The crystallite size was determined by the Debye–Scherrer equation of the peak (111) and is presented in Table 2. The crystallite size decreased from 16.4 nm for TiCN coatings to 14.6 nm for TiSiCN. The possible explanation is the penetration of impinging Si ions into the lattice of the TiCN coating, and the decreased generation of defects, led to a decreased number of preferential nucleation sites causing reduced grain [35,36]. The strain ε was determined by the Williamson–Hall plot method and is presented in Table 2. In both cases, the strain was low, but after the addition of Si, the strain was significantly reduced. It is reasonable to believe that the observed difference in strain of the TiSiCN coatings is related to the amorphous phases in which Si can be found (Si, SiNx, SiCN) or the C=C phase at grain boundaries, which is detrimental for crystallite development. These effects will be goal of another paper, in which other complex analyses, such as TEM and XPS, will be carried out.

The lattice parameter obtained from 2θ values of the (111), (200), (220) and (311) of TiSiCN was 4.2771, while for TiCN it was 4.2725. According to JCPDS no. 042-1488, TiCN has a lattice parameter equal to 4.2644. When the diffraction peak is shifted towards lower Bragg angle, an increase of lattice parameter is observed. This result is in good agreemen<sup>t</sup> with those of Constantin et al. related to TiCN coatings [37]. Both coatings exceeded the size of the standard unit cell parameters. This finding can be a sign that the hardness can be low and, on the other hand, the toughness will be high [38].

Grieveson et al. stated that the TiN and TiC compounds do not combine perfectly; the mixture is far from the ideal Raoultian behavior [39,40]. The same conclusion was also considered by Levi et al., who revealed that the occupation of all sites by Ti, C and N is not random, it is a TiN-based structure with a replacement of the N site by C, forming an FCC or tetragonal structures on the low concentration of vacancies that might be present [41]. All these aspects had an important effect on the characteristics of the TiCN coatings.

### *3.2. Morphology and Roughness (Before Corrosion)*

The surface morphologies of the uncoated CoCr substrate and the TiCN and TiSiCN coatings can be observed in Figure 2. The appearance of the coatings is characterized by a continuous coverage of the substrate with visible individual microparticles distributed on the surface. These droplets are considered defects and are characteristic of the energetic ejection of the particles during the arc deposition process [42,43].

 **Figure 2.** SEM images of the (**a**) uncoated CoCr substrate and (**b**) TiCN and (**c**) TiSiCN coatings.

### *3.3. Roughness of Surface*

The main surface texture parameters of the investigated surfaces are: *Ra*: the arithmetic average of the roughness profile, *<sup>R</sup>*q: the root mean square average of the roughness profile, *Sk*: skewness (according to the ISO 2517 standard [44]), and are presented in Table 3. The uncoated substates had an *R*a roughness around 50 μm, as intended. The roughness of the coated surfaces was much higher and it was more evident in the TiCN coated surfaces. Similar values of both *Ra* and *Rq* were found for the TiCN and TiSiCN coatings.


**Table 3.** Main roughness parameters of the investigated specimens.

A negative value of *Sk* shows that the surface consists of many valleys, while a surface with a positive *Sk* contains mainly peaks and asperities. Therefore, a surface with positive *Sk* has a good tribological performance in lubrication conditions. Taking into account this observation, both investigated coatings exhibited a positive *Sk*, which is suitable for load bearing implants, which work in a corrosive environment. The uncoated substrate had an *Sk* close to 0, meaning that the surface was very flat. The TiCN surfaces had a high *Sk*, indicating many peaks and asperities, whereas Si-containing coatings had a smaller value, indicating a decrease in peaks and asperities. The dependence of corrosion resistance on these parameters will be discussed below.

### *3.4. Hardness and Elastic Modulus*

Nanoindentation is a widely used technique, and it is a powerful tool to investigate the nanomechanical properties of materials, such as hardness and elasticity [45]. In order to overcome the substrate and roughness effects on nanoindentation tests, many studies have pointed out the importance of two basic rules. They refer to the maximum penetration depth, which should not exceed 1/10 of the layer thickness, and should be at least 20 times higher than the average roughness of the indented surfaces [46,47]. Another important limitation that needs to be considered for a reliable nanoindentation test is represented by the tip geometry and radius used, which limits the good testing contact depths, hence, in our case, was necessary to obtain contact depths higher than 40 nm [48].

In view of these testing conditions, the hardness and reduced modulus of the investigated sample coatings were determined according to the Oliver–Pharr method using similar force–displacement curves to the ones presented in Figure 3a. As it can be seen, these curves were obtained in both elastoplastic (CoCr substrate) and nearly elastic regimes (TiCN and TiSiCN coatings) [49], which were used to obtain the mechanical parameters of the samples. Note that all experimental force–displacement curves were qualitatively similar for the TiCN and TiSiCN coatings, providing indentation depth values smaller than 130 nm and exhibiting well-defined edges of the residual indent impressions (inset in Figure 3a). Moreover, the relatively low surface roughness in the investigated regions of the samples (~5 nm) was confirmed by performing several scans on 4 μm<sup>2</sup> areas using the same Berkovich tip. It has to be noted that this roughness corresponds to the flat areas between the microparticles found on the surface and it is therefore much smaller than the values reported for 4000 μm length surface scans. The resulting H = 7 ± 0.2 GPa and E = 97 ± 1.1 GPa values for the CoCr substrate are much smaller than the ones obtained for the investigated coatings. The hardness of the coatings significantly increased from 36.6 ± 2.9 GPa for the TiCN coating to 47.4 ± 1 GPa for TiSiCN, despite the small amounts of Si added into the TiCN layer structure (Figure 3b), while the e ffective modulus changed from 277 ± 8 GPa to 310 ± 2 GPa. The hardness enhancement originates from the grain size e ffect (H~d−1/2, according to the Hall–Petch strengthening mechanism), since the TiSiCN coatings have smaller grain sizes. The hardness and modulus values were consistent with the previous results [50,51]. This implies that the condition H/E > 0.1 [52] was fulfilled for both coatings, as the calculated values for TiCN and TiSiCN were about 0.13 and 0.15, respectively. As the TiCN structure was modified, the sharp increase in the H<sup>3</sup>/E<sup>2</sup> ratio (resistance to plastic deformation) from 0.63 to 1.1, testifies to the superior toughness of the TiSiCN layer [53], and it may be ascribed to the beneficial e ffect of the Si content. The H<sup>3</sup>/E<sup>2</sup> parameter was considered to be a parameter sensitive to the tribological and corrosion properties of the materials [54].

**Figure 3.** (**a**) Experimental force–displacement curves of the CoCr substrate, TiCN and TiSiCN coatings and (**b**) averaged hardness and reduced Young's modulus values.

### *3.5. In Vitro Corrosion Investigations*
