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Article

Mechanical and Tribological Behavior of TiAlSiN/AlSiN Coatings Depending on the High-Temperature Treatment

by
Stefan Kolchev
1,
Lilyana Kolaklieva
1,*,
Daniela Kovacheva
2,
Genoveva Atanasova
2,
Tetiana Cholakova
1,
Vasiliy Chitanov
1,
Ekaterina Zlatareva
1,
Roumen Kakanakov
1 and
Chavdar Pashinski
1,3
1
Central Laboratory of Applied Physics, Bulgarian Academy of Sciences, 61 St. Petersburg Blvd., 4000 Plovdiv, Bulgaria
2
Institute of General and Inorganic Chemistry, Bulgarian Academy of Sciences, 1113 Sofia, Bulgaria
3
Department of Mechanics, Technical University—Sofia, Br. Plovdiv, 25 Tsanko Dyustabanov St., 4000 Plovdiv, Bulgaria
*
Author to whom correspondence should be addressed.
Coatings 2025, 15(5), 542; https://doi.org/10.3390/coatings15050542
Submission received: 2 April 2025 / Revised: 24 April 2025 / Accepted: 29 April 2025 / Published: 30 April 2025
(This article belongs to the Special Issue Tribological and Mechanical Properties of Coatings)
Browse Figures
Versions Notes

Abstract

:
TiAlSiN/AlSiN coatings, with 3 and 30 periods, were successfully deposited by cathodic-arc evaporation technology. The composition, structure, mechanical, and tribological properties were studied at thermal treatment from 700 °C to 900 °C. The SEM observation and EDS analysis verified the dense structure and stable element composition in the coating depth at increased temperatures. A limited surface oxidation was identified at 800 °C, which increased moderately at a higher temperature of 900 °C. The coating period displays a nanocomposite structure of TiAl(Si)N and AlN nanograins incorporated in an amorphous Si3N4 matrix obtained by XRD and XPS analyses. The coatings exhibit high hardness of 41.1 GPa and 36.4 GPa for the 3- and 30-period coatings, respectively. The coatings with higher modulation periods demonstrate an excellent high temperature hardness and resistance to elastic and plastic deformations up to 900 °C. The hardness of the coatings with a smaller modulation period reduces to 29.7 GPa at the same temperature, causing a decrease in the H/E and H3/E*2 ratios. The tribological tests found that the high-temperature wear resistance depends strongly on the coating composition and architecture. An oxidation wear mechanism dominates the coatings with a large modulation period, and the wear rate decreases with a temperature increase. Abrasive wear is predominant in coatings with a lower modulation period, leading to an increasing wear rate. Wear rate values of 7.27 × 10−6 mm3/N·m and 8.53 × 10−6 mm3/N·m were determined after annealing at 900 °C for the 3- and 30-period coatings, respectively.

Graphical Abstract

1. Introduction

In many modern industrial technologies, parts and mechanical components are subjected to the harsh effects of various factors such as aggressive environments, oxidation processes, high temperatures, etc. This imposes increased requirements on coatings that protect working surfaces regarding temperature-resistant hardness and resistance to deformation, oxidation, and corrosion. For instance, due to the high operating speeds of cutting and forming tools used in the metal-working industry, the resistance of coatings at temperatures of the order of 800 °C and higher is an essential factor for extending their service life [1]. Temperature resistance refers to the main properties characterizing coatings, such as mechanical and tribological properties, corrosion and oxidation resistance, and wear resistance [2,3,4,5]. The resistance of coatings to thermal stress is mainly determined by their composition, structure, and architecture [3,6,7].
TiAlSiN nanocomposites have been widely used for industrial applications [8]. Due to their specific structure of TiAlN nanocrystals embedded in a Si3N4 phase, they pose enhanced mechanical and tribological properties at elevated temperatures [9,10,11,12]. The coating hardness increases to 40.9 GPa, and the coefficient of friction decreases to 0.2 by adding 9% of silicon to the coating composition [11]. The TiAlSiN coatings exhibit improved wear resistance of 2.33 × 10−6 mm3/Nm during annealing up to temperatures as high as 800 °C [12]. However, the high-temperature applications of TiAlSiN nanocomposites are limited by processes ongoing in this system. It is known that, at temperatures exceeding 800 °C, a spinodal decomposition of the TiAlN solid solution occurs [13]. At temperatures as high as 900 °C, an oxidation process starts in TiAlSiN films and Al2O3, (TiO2 + TiAlSiN) and TiAlSiN compositions could be recognized [14].
In recent years, multilayer design has been proven to be an extremely suitable approach for overcoming the limitations of monolayer coatings [15]. Multilayer architecture possesses notable potential for achieving coatings with superior performance that is impossible in monolayer structures. The multilayers allow alternating two or more layers to combine their best characteristics. Through the suitable composition and structure of the constituent layers, the elasticity of superhard monolayer coatings can be increased and their toughness can be improved. Similarly, their temperature resistance and wear resistance can be improved as well. It is known that the ultra-high hardness of monolayer coatings is a prerequisite for good wear resistance. However, it is also a prerequisite for increased brittleness. An appropriate combination of a superhard sublayer with a more elastic one could improve the multilayer toughness [16]. In addition, the interfaces in the multilayer structure hamper the dislocation spreading and crack propagation, thus enhancing the durability of the coating [17]. Also, energy distribution at the interfaces improves the adhesion and cohesive properties of the coatings [18]. Multilayer coatings have been intensively developed and investigated based on the enhanced mechanical and tribological properties of the TiAlSiN monolayer. Various studies report improved hardness, enhanced elastic strain to failure, and improved resistance to plastic deformation compared to the monolithic TiAlSiN film due to the multilayered architecture [19,20,21,22,23,24]. The role of period thickness [19,20] and technological parameters on the mechanical properties, composition, and structure are also discussed [21]. It was found that the increase in the thickness of the sublayer with better elasticity causes an improvement in ductility [19]. The modulation period reduction in the AlTiN/AlTiSiN coatings affects the enhancement of adhesion strength and mechanical properties, whereas the friction coefficient and wear rate are reduced [20]. A few papers present the effect of high-temperature treatment on the mechanical and tribological properties of the developed TiAlSiN-based multilayers [2,25,26,27]. The temperature is ranged from 400 °C to 1000 °C in the different papers. The study of TiAlSIN/NiCr multilayer coatings with varied periods showed that an increase in the period modulation has a decreasing effect on the hardness at room temperature. It causes a reduction of the coefficient of friction and wear rate at temperatures of 600 °C [25]. The dependence of the mechanical properties on the bilayer period of TiAlSiN/VSiN coatings showed that the hardness increases with the bilayer period decrease of the as-deposited specimens. It was found that the coefficient of friction decreased at a temperature of 700 °C due to the formation of a surface V2O5 oxide, which acts as a lubricant [26]. Chang et al. reported an AlCrBN/AlTiSiN multilayered coating with high hardness stability in a vacuum and enhanced oxidation resistance in air during annealing at 900 °C [2]. High hardness of 38.3 GPa and elastic modulus of 463.7 GPa were reported for AlCrN/TiAlSiN coatings [27]. They exhibit a stable wear rate of 2.5 × 10−6 mm3/Nm at a temperature of 500 °C to 800 °C and better oxidation resistance than the corresponding monolayers. These studies have shown that the high-temperature behaviors of TiAlSiN-based multilayers strongly depend on the composition and structure at thermal treatment.
In our previous paper [28], we presented the new developed TiAlSiN/AlSiN multilayer coatings, which are characterized by the nanocomposite structure of both sublayers. The role of the modulation period on the structure, composition, and mechanical properties of the as-deposited coatings was studied. This investigation showed that by combining a TiAlSiN layer having enhanced hardness with an AlSiN one having better elasticity, it is possible to obtain a coating with superhardness of 49.5 GPa and a low elastic modulus of 430 GPa and thus to improve the elastic strain resistance to 0.11, plastic deformation resistance to 0.58 GPa, and elastic recovery to 68%. Because these coatings were developed for high-temperature industrial applications, we continued the study of their behavior at high temperatures up to 900 °C. In this paper, we present the results of the investigation of the mechanical and tribological behavior of TiAlSiN/AlSiN multilayer coatings with varied period modulation subjected to high thermal treatment.

2. Materials and Methods

The TiAlSiN/AlSiN coatings were deposited onto two types of substrate. Square stainless steel (10 mm × 10 mm) DIN 1.4541 plates were used for specimen preparation planned for SEM/EDS, XRD, and XPS analyses. Mechanical and tribological properties were studied on coupons of high-speed stainless steel (HSS DIN 1.3343) with a diameter of 20 mm and 5 mm in thickness. Two multilayer TiAlSiN/AlSiN coating architectures were investigated, with 3 and 30 periods, named ML-3 and ML-30, respectively. Before the deposition, the substrates were cleaned in an ultrasonic bath with a standard alkaline solution. After that, they were rinsed in de-ionized water and dried at 130 °C. In addition, the surface cleaning was performed in the vacuum chamber in an Ar discharge at a bias voltage of 1000 V. The deposition was performed from Ti (99.99 wt.%) and AlSi (82 at. % Al, 18 at. % Si) LARC® cathodes in a vacuum system Platit π80+. Nitrogen (99.9999%) at a pressure ranging from 9 × 10−1 Pa to 4 Pa was used as a reactive gas. The coating architecture contains a Ti adhesion layer, transition TiN and gradient TiAlSiN layers, and a main coating film. The letter is formed by alternating periods of TiAlSiN and AlSiN sub-layers. Finally, a TiAlSiN film was deposited as a top layer. The cathode currents and bias voltage were changed depending on the grown layer [28]. The deposition temperature was 500 °C. Immediately after deposition, the coatings were annealed for two hours at 525 °C in nitrogen to complete the formation of the nanocomposite structure. The thermal treatment was performed at temperatures ranging from 700 °C to 900 °C for 2 h at each temperature in argon using a resistance furnace (CLAP, Plovdiv, Bulgaria). After heating, the specimens were cooled down in the quartz tube of the furnace in an argon atmosphere.
The morphology and element composition of the surface and cross-section were obtained on a Hitachi SU 5000 Schottky field emission scanning electron microscope (Hitachi, Tokyo, Japan) equipped with an energy-dispersive X-ray spectroscopy (EDS) system (Thermo Scientific, Waltham, MA, USA). The cross-section was made using ion milling equipment IM 4000 Plus (Hitachi, Japan). The surface composition and electronic structure of materials were investigated by X-ray photoelectron spectroscopy (XPS). The measurements were performed in a VG ESCALAB II (VG Scientific, Manchester, UK) system using AlKα radiation with an energy of 1486.6 eV. The binding energies (BEs) were determined with an accuracy of ±0.1 eV, utilizing the C1s line at 285.0 eV (from adventitious carbon) as a reference. The composition and chemical bonding of the samples were investigated based on the areas and binding energies of the photoelectron peaks and Scofield’s photoionization cross-sections. XPSPEAK 4.1 software was used to deconvolute the XPS data. To determine the coatings’ phase composition and obtain unit cell parameters and mean coherent domain size (crystallite size) of the presented phases, all samples were analyzed by Powder X-ray diffraction (PXRD) using a Bruker D8 Advance diffractometer (Bruker AXS, Karlsruhe, Germany) with a Cu tube operating at 40 kV/40 mA. The patterns were recorded within the 5–80 degrees 2θ by the LynxEye detector (Bruker AXS, Karlsruhe, Germany). The phase identification was performed using EVA software (version V4) and the ICDD PDF-2 (2021) database, and the program Topas 4.2 was used for the Rietveld analyses.
Mechanical properties were investigated by a Compact Platform CPX (MHT/NHT) system (CSM Instruments-Anton Paar GmbH, Graz, Austria) equipped with a diamond Berkovich indenter. The measurements were carried out at a loading force from 20 mN to 200 mN. The nanohardness and module of elasticity were determined by the Oliver & Pharr method [29].
Tribological tests were performed on a Bruker UMT TriboLab tribometer (Bruker Nano Surfaces Division, San Jose, CA, USA) using the ball-on-disk method. A sapphire sphere (Al2O3) with a diameter of ¼” (6.35 mm) and a hardness of 19.2 GPa was used as a counter body. The normal load on the counter body was 5 N, the rotation speed was fixed at 150 rpm, and the offset radius was 2 mm. The wear duration of each test was 4000 s (≈67 min), and the total sliding distance was 1256 mm. The number of cycles performed was set to 10,000. The tests were performed under dry sliding conditions in the air at room temperature (24 ± 2 °C) with a relative humidity of 30 ± 2%. The formed wear tracks were observed on the microscope of the equipment Compact Platform CPX (MHT/NHT) (CSM Instruments, Anton Paar, Austria). Their dimensions (width and depth) were measured in four sections located at 90° to each other using a Rockwell diamond indenter with a tip radius of 200 μm. The profiles were recorded by traversing the formed tracks and equal lengths of the adjacent surrounding surfaces on both sides of the tracks. The depth of the wear track was determined using a prescan mode of the scratch test program of the equipment Compact Platform CPX (MHT/NHT) at a constant load of 0.03 N and a movement speed of 0.1 mm/min of the sample relative to the indenter over a distance of 1.6 mm. The worn volume of the coating was calculated using the equation V = P d / 6 W ) ∗ ( 3 P d 2 + 4 W 2 ∗ 2 π r  [30]. In this formula, Pd is the maximum penetration depth of the indenter in the track, W is the track width, and r is the track radius. The Archard equation K = V / ( F n ∗ S ) was used to calculate the wear rate [31,32], where Fn is the load of 5 N applied to the sapphire sphere, and S is the total distance traveled of 1256 mm.

3. Results and Discussion

3.1. Morphology and Composition

The surface and cross-section morphology of the TiAlSiN/AlSiN coatings were evaluated by SEM observation. Figure 1 presents the surface SEM images of the as-deposited and annealed at 900 °C ML-3 and ML-30 coatings. The coatings exhibit a dense surface due to the Si content in the composition [33]. Silicon promotes the grain size refinement and formation of nanocomposite structure surrounding the grains by an amorphous phase. Randomly distributed micro-particles and pits can be distinguished within the surface matrix. These features are typical for cathodic arc technology. Micro-particles are created when drops of molten target metal reach the surface and adhere to it. Some of them are re-spattered during deposition, originating pits on the surface. It should be noted that due to the LARC® technology, macro-particles and pits are reduced. The number of micro-particles in the as-deposited ML-30 coating increases compared to that in the ML-3 coating. A similar increase in the number of micro-particles with an increase in the number of periods is also observed for multilayers AlTiN/AlTiSiN [20]. No noticeable difference is observed between the surface topology of the as-deposited and annealed at 900 °C coatings. The sizes of similar features vary within the same ranges and do not depend on the modulation period and thermal treatment up to 900 °C. Despite the diffusion of Al toward the surface at higher temperatures, agglomerations on the surface were not observed. In addition, no micro-cracks are visible after annealing, implying that the coating morphology remains stable even at high temperatures.
Table 1 presents the surface composition as determined by the EDS analysis in an area mode. It is seen that with an increase in the annealing temperature up to 900 °C, the nitrogen amount decreases gradually with 2.7 at. % in the ML-3 coating and 4.3 at. % in the ML-30 coating. This decrease can be referred to as the release of weakly bound nitrogen ions that are adsorbed at the grain boundaries [32]. The increased Al content in the ML-3 could be caused by releasing the Al atoms incorporated in the TiN lattice at high temperatures. The free atoms out-diffuse toward the surface through the grain boundaries. The results from the EDS analysis revealed that there is no oxygen on the surface up to 800 °C. A small amount was detected after annealing at 800 °C. At 900 °C, the oxygen content increases, but its amount cannot affect the volumetric characteristics of the coatings. The oxygen presence on the coating’s surface after annealing at higher temperatures is most probably caused by the residual oxygen in the quartz tube of the furnace. The presence of carbon is due to the exposure of the samples to the external environment.
The SEM cross-section images of as-deposited and annealed at 900 °C coatings reveal the dense structure without any features attributed to the columnar growth (Figure 2). Besides the effect of Si, the existence of interfaces in the coating architecture hamper the columnar grain growth. It is seen that, during the growth, the sub-layers repeat the topology of the substrate. The planarity of the interfaces is also disturbed by the incorporated micro-particles originating from the molten droplets embedded in the growing layer. The inserts show an enlarged part of the multilayer architecture. It is seen that the coatings are characterized by a compact structure without pinholes and cracks. No cracks or peeling of the interfaces were observed, which indicates good cohesive strength up to temperatures as high as 900 °C. The element composition was determined by EDS analysis in a linear mode. Measurements were performed in several places, and the average values were calculated. The dependence of the element concentration on the thermal treatment is present in Figure 3. A small decrease in the Al content was determined after annealing at 800 °C. This result is consistent with the Al increase on the coating surface at the same temperature. The dependences of other elements display that the concentration does not change noticeably in the coating depth. This implies that both types of TiAlSiN/AlSiN coatings, ML-3 and ML-30, have a stable composition up to 900 °C. Indeed, the results from the element distribution in the coatings determine a Ti0.24Al0.64Si0.12N composition for the as-deposited ML-3 multilayer. The annealing at a temperature of 900 °C caused a Ti0.24Al0.63Si0.13N composition of the coating. Similarly, for the ML-30 coating, Ti0.28Al0.58Si0.14N and Ti0.28Al0.57Si0.15N compositions were determined for the as-deposited and annealed at 900 °C specimens, respectively. However, the results indicate that the composition of the as-deposited coatings with different modulation periods differs slightly in terms of the element amounts. This discrepancy could be due to technological factors such as changes in the reactive gas flow or the base voltage change during sublayers deposition, rather than the effect of the period modulation.
The compound composition of the multilayer TiAlSiN/AlSiN coatings before and after annealing at 800 °C and 900 °C was identified by XPS analysis. The measurements were performed on the coating surfaces after 10 min of etching by Ar+ ions. Figure 4 displays the Ti2p, Al2p, Si2p, and N1s spectra obtained with ML-3 and ML-30 coatings. The Ti2p peak is expressed by spin doublet peaks, Ti2p3/2 and Ti2p1/2, with a spin–orbit splitting of 5.9 eV (Figure 4). The Ti2p peak was deconvoluted into three double peaks. The Ti2p3/2 spectra consist of peaks, which correspond to (Ti, Al)N (455.6 ± 0.2 eV), TiOXNY (457.6 ± 0.2 eV), and TiO2 (459.7 ± 0.2 eV) [34,35]. The peak positions are the same in both types of coatings, ML-3 and ML-30, implying the same Ti-based compounds. The photoelectron peak intensity of the TiAlN phase is slightly lower in the coating with 3 periods than in the coating with 30 ones (Figure 4a). Hence, the amount of the TiAlN compound in the ML-3 could be less, which agrees with the lower Ti content determined by EDS analysis. The annealing at 800 °C did not change the peak positions, i.e., no transformations of the compounds were detected. Nevertheless, the TiAlN content slightly decreases, and the intensity of TiO2 and TiOxNy peaks increases. These changes are more pronounced in the ML-30 coating (Figure 4b). After the thermal treatment at 900 °C, no alterations in the compound composition were observed in the ML-30 coating. However, in the ML-3 coating, the amounts of TiAlN and TiOXNY decreased, while the quantity of TiO2 increased significantly (Figure 4a). The Al2p peak of the as-deposited coatings was decomposed into two peaks. The most intensive peak, corresponding to the binding energy of 74.4 ± 0.2 eV was recognized as an Al-N bond in AlN and TiAlN [35]. The peak positioned at BE 75.8 ± 0.2 eV relates to the Al-O bond in Al2O3. In the as-deposited coatings, the peak intensity is too small, supposing that the oxidation is caused by exposition on air rather than the coating oxidation because of the annealing. After annealing at 800 °C, that peak increased slightly, implying that the surface oxidation was due to the thermal treatment. Similar to the Ti2p peak, a further increase in the temperature did not affect the Al2p spectrum and, consequently, the composition of the ML-30 coating. The measurement of the ML-3 coating determined only a peak at 75.0 eV, which is attributed to Al2O3 [36]. This result indicates that, after annealing at 900 °C, the oxidized coating surface is dominated by Al2O3. The deconvolution of the N1s spectrum revealed that it is composed of two peaks centered at 396.8 eV and 398.5 eV, determining two nitride compounds (Figure 4). The more intensive peak was found at BE 396.8 eV, whose position is associated with nitrogen bonds in the compounds TiN, AlN, and TiAlN [34,35]. The peak located at BE 398.5 eV corresponds to Si-N bonds in Si3N4 [37]. There is an overlap in the Si2p region with an Al KLL Auger peak, which makes deconvolution difficult or even impossible (Figure 4). Nevertheless, peaks can be observed at 101.9 eV and 102.8 eV. The Si2p core level is positioned at a binding energy of 101.9 eV, associated with Si-N bonds in Si3N4 [37,38]. The peak at 102.8 eV is attributed to Si-O bonds.
The obtained XPS spectra revealed that the TiAlSIN/AlSIN coatings are composed of nitrides TiN, AlN, TiAlN, and Si3N4. After annealing at 800 °C and 900 °C, Ti-O-N, TiO2, Al2O3, and Si-O oxides were observed on the surfaces. These changes in the coating composition lead to a decrease in the nitrides’ amount on the surface.
Powder diffraction patterns of the as-deposited and annealed at 800 °C and 900 °C coatings with different modulation periods can be seen in Figure 5a. Except for the peaks from the substrate, two phases can be distinguished for ML-3 and ML-30 coatings, namely, cubic TiN with fcc-lattice and hexagonal hcp-AlN. The PXRD patterns do not show any peaks of individual crystalline silicon nitride phases, so it is assumed that part of the silicon is incorporated into the TiN lattice and the rest is presented in the coatings in an amorphous silicon nitride state.
A strong preferred orientation is observed for the TiN phase with the (200) peak being the most intensive, while other phase peaks (111) and (220) show very low intensity. Hence, the dominant (200) peak determines the main growth plane of this phase. The preferred growth plane of the film depends on the competition between the surface and the strain energy [39]. In the fcc-TiN crystal phase, the (111) plane is characterized by the lowest strain energy, while the (200) plane has the smallest surface energy [39,40]. In the case of a multilayer structure, the strain energy resulting from the lattice deformation can be converted into interface energy. Accordingly, the strain energy of the (200) plane becomes lower than that of the (111) plane and thus facilitates the grain growth [40]. Several technological parameters, such as bias voltage [41,42], nitrogen flow [21,43], ion current density [44], and modulation period [19,20] were reported to influence the preferential growth of coatings in the (200) plane. Following these studies, it can be concluded that the technological conditions under which the TiAlSiN/AlSiN multilayers were deposited contribute to the preferable growth of the (200) plane of the fcc-TiAl(Si)N phase.
In the diffraction spectra of the ML-3 and ML-30 coatings, the (200) peak is slightly shifted to the higher 2θ values compared to the standard peak for TiN (PDF 00-038-1420) (Figure 5a). This shift can be associated with the existence of internal stress due to the contraction of the crystal lattice of the fcc-TiN phase as a result of the replacement of a part of Ti atoms (at radius 47.867 Å) by Al and/or Si atoms (both with smaller radii than Ti), which leads to the formation of the fcc-TiAl(Si)N phase [19,32]. The shift is more pronounced in the ML-3 coatings, as the offset at 2θ is 0.2, while, in the ML-30 coatings, it is barely 0.07. Incorporating Al and/or Si atoms into the TiN lattice causes a decrease in the lattice parameter. Indeed, the calculated unit cell parameter of the as-deposited coatings, 4.2249 Å for ML-3 and 4.2386 Å for ML-30, is lower than the referent, reported in PDF 00-038-1420 (Table 2), corroborating that part of the Ti atoms is substituted by smaller Al and/or Si ones in the lattice, thus forming a solid solution. The TiN lattice contraction caused by incorporating Al(Si) atoms was evaluated by the formula ( α T i N − α T i A l S i N / A l S i N ) / α T i A l S i N / A l S i N [45]. Values of 0.40% and 0.07% for the ML-3 and ML-30 coatings, respectively, were found. This result implies that TiN dominates the properties of the ML-30 coating.
In the ML-3 coatings annealed at 800 °C, a shift of the (200) peak to the lower value of 2θ is observed, bringing it closer to the standard for TiN (Figure 5b). This shift can be attributed to two processes at high temperatures. One of them is the spinodal decomposition of TiAlN, which is observed at 800 °C [46], and the relaxation of the internal stress. Accordingly, the unit cell parameter of the ML-3 coating slightly increases. The position of the (200) peak and the size of the unit cell parameter indicate that a part of the TiAlN phase is preserved. Consequently, two phases, TiAlN and TiN, exist in the ML-3 coating after annealing at 800 °C. Further, an increase in temperature to 900 °C does not change the position and intensity of the peak, which suggests the stability of the crystal structure. The annealing of the ML-30 coating at 800 °C and 900 °C does not cause noticeable changes in the position of the (200) peak compared to that of the as-deposited coatings. However, the FWHM decreased from 0.702 ⁰2θ for the as-deposited coating to 0.658 ⁰2θ and 0.621 ⁰2θ after annealing at 800 °C and 900 °C. This feature is associated with rearranging towards lower-energy sites of the structural defects (point and line defects) generated during the coating growth [13]. In the TiN coatings, this feature was observed even at temperatures as low as 500 °C [47]. The unit cell parameter of the TiN phase remains relatively close to the referent one and does not change significantly with heating.
The mean coherent domain size (crystallite size) obtained for this phase is around 20 nm and does not change with the annealing.
The second crystalline phase identified in the coatings is hexagonal AlN. Because, according to the EDS analyses, the Al content exceeds the solubility limit of x ≥ 0.7, AlN grows in a wurtzite structure. The peaks related to (100) and (110) planes are slightly shifted to the lower 2θ values, implying the expansion of the crystal lattice. Indeed, the calculated cell parameters for AlN are larger than those presented in the referent PDF 00-025-1133 card (Table 2). The possible explanation for this observation is related to the constraint between the lattices of the AlN and TiN. The peak’s position and unit cell parameters of the ML-3 coating do not change noticeably with an increase in the annealing temperature up to 900 °C. It is worth mentioning that for the 30-period coating, the unit cell parameters of the AlN phase tend to shorten upon annealing, indicating its lattice relaxation. Contrary to the TiN phase, the crystallite size of AlN is small (3 ηm), which is seen from the large peak width and low maximal intensity. Despite the low peak intensity, the evaluation of the phase quantity based on integral intensity shows almost equal mass content as the TiN phase.

3.2. Mechanical Properties

The hardness of the as-deposited and thermally treated TiAlSiN/AlSiN coatings was measured in the interval of loading forces from 20 mN to 200 mN. The results obtained at a loading of 30 mN were used as representative because they were obtained at a maximum penetration depth lower than 330 nm. This penetration depth meets the requirement of being less than one-tenth of the coating’s thickness, which was 3.81 µm for ML-3 and 3.30 µm for ML-30. Therefore, it could only represent the coating hardness. The coating hardness of 41.1 GPa and 36.4 GPa of ML-3 and ML-30, respectively, was determined from the load-displacement curves. The acquired hardness values depend on the coating composition and structure. EDS analyses obtained that the silicon content is between 6.9% and 7.9%. It is known that a Si amount between 4% and 10% causes grain size refinement and nanocomposite structure formation [48,49,50]. In the Ti-Al-based coatings, this Si quantity also contributes to the solid solution formation. Accordingly, the Si content in the TiAlN/AlSiN coatings facilitates several hardening mechanisms, solution hardening, Hall–Petch strengthening, and a nanocomposite structure formation [50,51]. The XPS and XDR studies revealed that both coating types have a nanocomposite structure consisting of TiAlN, TiN, and AlN nanocrystals incorporated in the Si3N4 amorphous phase. The formation of a nanocomposite structure leads to enhanced coating hardness. In addition, the multilayer structure supports the hardness increase, hampering the dislocation and crack propagation. Our results show that the ML-30 coating with a lower modulation period of 123 nm exhibits lower hardness than the ML-3 with a modulation period of 1248 nm. Usually, it is accepted that a decrease in the modulation period causes an increase in the coating hardness. However, in multilayers with a very small modulation period, an existing shielding effect may be overshadowed by different factors, leading to a hardness decrease [52]. At small modulation periods (λ < 50 nm), atomic mixing occurs at the interface and may become a factor for the dislocation slip at the interface [53]. The sublayer thickness could also be a factor in the crack propagation. It is reported that for the c-Ti0.4Al0.6N/h-Cr2N coatings, the hardness decreases for the modulation period of 10 nm and the softer sublayer thickness is lower than 75% of the total coating thickness [54]. The increased droplet amount in the coating could additionally affect the hardness at λ < 60 nm [55]. In the ML-30 coating, the modulation period is greater than that where factors limiting the shielding effect of the interface have been observed. Therefore, it could be supposed that the lower hardness of the ML-30 coating is not due to overshadowing the shielding effect. According to the position of the (200) XRD peak, the properties of the ML-3 coating are dominated by the fcc-TiAlN phase, i.e., the solid solution hardening occurs, whereas the behavior of ML-30 coating is mainly determined by the fcc-TiN phase. We suggest the slightly higher hardness in the ML-3 coating is due to the dominant solid solution TiAlN phase. The hardness of both types of coatings did not change after annealing at 700 °C (Figure 6a). The increase in the temperature to 800 °C caused a small decrease in hardness to 39.3 GPa for the ML-3 multilayer. A more pronounced reduction to 29.7 GPa was obtained for the ML-30 multilayer. This decrease in hardness due to the surface oxidation was verified by the EDS, XPS, and XRD analyses. Other factors, such as the spinodal decomposition of a part of TiAlN, which occurs at a temperature of 800 °C, could partially reduce the effect of solid solution hardening. The strain stress relaxation existing at interfaces can also contribute to a hardness decrease. The further increase in temperature enhances the oxidation process, resulting in a reduction of the hardness. Values of 34.4 GPa for ML-3 and 27.4 GPa for ML-30 were determined after annealing at 900 °C. The module of elasticity follows a similar dependence on the temperature. The XRD and EDS studies did not identify any significant changes in the structure and composition in the coating’s depth. Indeed, small changes of 55 GPa and 54 GPa in the modulus of elasticity were found for both coatings after the thermal treatment at 900 °C. The behavior of the hardness and module of elasticity during the thermal treatment up to 900 °C affects the H/E and H3/E*2 ratios (Figure 6b). According to the dependence of the H/E ratio on the temperature, the ML-3 multilayer exhibits better elastic strain to failure. It is attributed to the more stable composition and structure expressed by the better stability of the hardness and modulus of elasticity. Despite the good stability of the elastic modulus during the heating in the ML-30 multilayer, its resistance to elastic deformations decreases significantly from 0.1 to 0.072. This result is due to the greater decrease in the hardness of this coating. The same as-deposited coating shows better resistance to plastic deformations. However, it decreases sharply from 0.448 GPa to 0.126 GPa during the annealing. The ML-3 coating has enhanced stability against plastic deformations. Its H3/E*2 ratio changes less from 0.370 GPa to 0.291 GPa, suggesting better wear resistance. These results determine the ML-3 TiAlSiN/AlSiN coating to have better mechanical properties.

3.3. Tribological Properties

The adhesion strength between the coating and substrate was evaluated by the scratch test performed at a distance of 3 mm and a normal force progressively increasing from 2 N to 100 N by 10 N/min. The determined critical loads are presented in Table 3. According to the ISO 20502:2005(E) standard [56], three failure events were classified: cracking (Lc1), spallation (Lc2), and penetration of the coating to the substrate at the center of the track (Lc3). The obtained results show better adhesion strain of the M-3 multilayers. Increasing the treatment temperature from 700 °C to 900 °C initiated the appearance of spallation (Lc2) at loads between 77.8 N and 55.5 N, respectively. Penetration of the coating to the substrate at the center of the track (Lc3) was observed only at the high temperatures of 800 °C and 900 °C at 84 N and 85 N, respectively. A similar trend of failure events was also found for the 30-period coating. This test revealed that both coatings had excellent adhesion to the substrate at normal forces up to 50 N for the M-3 multilayer and 35 N for the M-30 multilayer.
Figure 7 presents a schematic diagram of the tribo-mechanical method “ball/pin-on-disk” used for determination of the coefficient of friction.
The coefficient of friction (CoF) of ML-3 and ML-30 TiAlSiN/AlSiN coatings as it was measured during the wear tests of the as-deposited and thermally treated specimens is presented in Figure 8. The trend of CoF differs for both coatings and shows a well-defined dependence on annealing temperature. The unsteady state period, equal to 480 s (1200 revolutions), is the same for the freshly deposited and annealed at 700 °C and 800 °C ML-3 coatings (Figure 8a). After that, the coefficient of friction remains relatively constant. The observed fluctuations in the progression of the curves suggest the formation, periodic accumulation in the contact zone, and periodic clearance of wear debris [57]. No noticeable changes in CoF were found for these coatings. The average values of 0.4 ± 0.043, 0.41 ± 0.031, and 0.35 ± 0.062 were determined for the as-deposited and annealed at 700 °C and 800 °C multilayers. The curve corresponding to the ML-3 coating annealed at 900 °C shows an oscillation in the CoF. After the unsteady state period reaches 700 s (1750 revolutions), a region with oscillations and peaks in the curve is visible up to 3000 s (7500 revolutions). The observed variation is caused by the increased surface oxidation after treatment at a temperature of 900 °C [24]. XPS analysis showed the predominant presence of titanium and aluminum oxide on the TiAlSiN surface of the ML-3 coating at this temperature. After 3200 s (8000 revolutions) the curve becomes smooth, showing a decrease in CoF to 0.27 ± 0.054. CoF measurements showed that the unsteady state periods of the annealed ML-30 coatings coincide. Moreover, it decreases almost twofold to 250 s (625 revolutions) compared to the multilayer ML-3. After that, the curves show relatively constant values (Figure 8b). A slight increase in the friction coefficient was recorded from 0.36 ± 0.039 at 700 °C to 0.40 ± 0.028 at 900 °C. This result suggests a relative stability of CoF for these coatings up to temperatures of 900 °C. An exception was found for the deposited ML-30 coating, in which the CoF demonstrated a transient period of up to 2700 s (6750 revolutions). From the beginning of the test to 1400 s (3500 revolutions), a plateau with significant oscillations in the CoF curve with an average value of 0.2 ± 0.135 is observed. The characteristic of the interval suggests the formation of abrasive particles, formed as a result of the crushing of the micro-protrusions from the surface of the coating. In the interval from 1400 s to 2700 s, an increase in the average value of CoF to 0.26 ± 0.023 is recorded. The newly formed plateau suggests a generation of a sufficient amount of debris/abrasive particles falling into the contact wear zones in the tribo pair and scraping the surface of the coating. After 2700 s, a decrease in the value of CoF to 0.23 ± 0.036 is observed, and a steady state is maintained until the end of the test. The behavior of the curve suggests grinding of the abrasive particles under continuous friction and a decrease in the roughness of the contact surfaces. These variations are caused by the increased internal stress existing in the coatings with many interfaces [32]. As a result of the thermal effect, this stress partially (or completely) relaxes. A low average CoF of 0.23 ± 0.116 was determined with the as-deposited ML-30 coating. A similar behavior of the CoF is observed in the TiAlSiN/NiCr coatings with 32 periods [25].
The dependence of the coefficient of friction of ML-3 and ML-30 multilayers is presented in Figure 9a. Both coatings exhibit an opposite tendency in the dependence on the thermal treatment. The CoF of TiAlSiN/AlSiN coatings with three periods is not affected by heating at 700 °C. Further increase in the annealing temperature causes a decrease in the coefficient of friction of those coatings. The obtained result suggests the existence of high adhesive strength [58]. An opposite relationship is observed for the coating with 30 periods. It exhibits a significant increase in the CoF even after annealing at 700 °C. Increasing the treatment temperature does not significantly affect the value of the CoF. This difference in the dependence of the CoF of the ML-3 and ML-30 coatings could be explained by the different wear mechanisms that dominated in them. The wear rate has the same dependence on the annealing temperature as the coefficient of friction. In the ML-3 multilayer, the wear rate smoothly decreased from 1.08 ± 0.09 × 10−5 mm3/N·m at the room temperature, 9.63 ± 0.82 × 10−6 mm3/N·m at 700 °C, 7.99 ± 0.91 × 10−6 mm3/N·m at 800 °C, to 7.27 ± 0.53 × 10−6 mm3/N·m at 900 °C, implying an improvement in the wear resistance (Figure 9b). The wear rate of the ML-30 coating showed an increase from 7.53 ± 0.68 × 10−6 mm3/N·m at the room temperature, 8.68 ± 0.76 × 10−6 mm3/N·m at 700 °C, to 9.70 ± 0.91 × 10−6 mm3/N·m at 800 °C. It exhibits a decrease to 8.53 ± 0.58 × 10−6 mm3/N·m at 900 °C.
Figure 10 presents the morphology and element composition of the wear tracks for the ML-3 and ML-30 TiAlSiN/AlSIN coatings. The depth profile of each track is also present. It is seen that in the ML-3 coating, the penetration depth (Pd) of the tracks decreases with increasing temperature (Table 4). The dependence is similar to the tracks’ width (W), except the one obtained at 900 °C. The micrographs show the presence of adhesives in the tracks, as well as wear debris located at both ends of the tracks. This morphology determines adhesive and abrasive wear mechanisms, which explains the observed oscillations in the CoF curves. The results of the EDS analysis in the worn track show an increase in the atomic concentration of oxygen with increasing temperature compared to the unworn coating surface. This result suggests a third oxidation wear mechanism, which is dominant at high temperatures. The increased Ti and oxygen concentrations at 800 °C, accompanied by a decrease in the nitrogen amount, suggest the formation of TiO2 simultaneously with SiO2 and Al2O3 [12]. The formed oxides act as a “solid lubricant” in the contact areas, reducing the wear rate [59,60,61,62]. The lowest wear rate at 900 °C is characterized by the wider track and smaller penetration depth compared to ones at 700 °C and 800 °C. The penetration depth and the width of the wear tracks on the ML-30 coatings increased with the temperature increase (Table 4). Similar to the ML-3 coating, after annealing at 900 °C, the track width does not follow the tendency observed at 700 °C and 800 °C. The micrograph of the as-deposited ML-30 coating shows micro-grooves along the sliding direction, determining the abrasive mechanism as dominant [63]. As the temperature increased to 700 °C and 800 °C, the amount and size of the adhesives inside the tracks increased, verifying the main adhesive wear mechanism at these temperatures. At a temperature of 900 °C, the track became narrower and smoother, and the adhesives were reduced. The EDS analysis verified that the oxygen content reached the maximum value, suggesting the increasing role of the oxidation mechanism of wear towards the abrasive and adhesive mechanisms. The observed results are similar to those obtained by Weicai and colleagues [64]. It is assumed that a lower coefficient of friction increases wear resistance. Despite the decrease in CoF at 900 °C on ML-30, the increased wear rate can be explained by the formation of a uniform oxide film, but the lower hardness of the coating worsens its resistance. And, due to the stresses generated by the wear process, a fourth wear mechanism is formed—“material fatigue”, leading to the appearance of cracks and separation of parts of the tribo-oxide layer. These cycles of oxide layer formation (oxidation wear); cracking, destruction, and removal of debris (fatigue wear); and subsequent re-formation of oxides lead to a significant increase in the wear rate.
The results from the investigation of TiAlSiN/AlSiN with different modulation periods indicate that several factors affect the tribological properties. Combining a harder sublayer with a more elastic one in the coating period benefits the coating’s resistance to elastic deformations (H3/E*2 ratio). In the thicker sublayer with higher elasticity, the stress is distributed over a larger volume, and the coating is more resistant to elastic and plastic deformations [19]. Hence, the hardness and the resistance to elastic deformation depending on the modulation period are the main factors determining the tribological behavior. Furthermore, composition and architecture are very important factors for good tribological properties. Thus, in the TiAlSIN/AlSiN coatings with three periods, the coefficient of friction and the wear rate decrease due to the oxides’ formation in a wear track, which leads to better tribological behavior at elevated temperatures. In the 30-period coating, the modulation period is thin and cannot oxidize enough to provoke an oxidative wear mechanism. However, the sublayers are destroyed, creating abrasive particles in the contact zone. As a result, an abrasive wear is predominant, and the friction coefficient and wear rate increase.

4. Conclusions

Composite coatings TiAlSiN/AlSiN with two modulation periods of 1248 nm and 123 nm, 3- and 30-period, respectively, were successfully deposited and studied regarding their composition and structure, mechanical, and tribological properties during heating at temperatures ranging from 700 °C to 900 °C. The SEM observation and EDS analysis verified the dense structure and stable element composition at high treatment temperatures. Limited surface oxidation of 1.9 at. % and 1.5 at. % oxygen was identified at 900 °C on the surface of the coatings with 3 and 30 periods, respectively. The surface and phase examination by XPS and XRD analyses revealed that both types of coatings exhibit a thermally stable nanocomposite structure of TiAl(Si)N and AlN nanograins incorporated in an amorphous Si3N4 matrix, since the thermal treatment up to 900 °C does not destroy the coating structure and architecture. This fact can be regarded as a valuable advantage of the obtained coatings, since structural stability during high-temperature treatment is a prerequisite for stable mechanical and tribological properties. The structure and composition, providing a grain refinement, solid solution, and nanocomposite hardening, cause increased hardness of 41.1 GPa and 36.4 GPa for the 3- and 30-period coatings, respectively. The coatings with higher modulation periods demonstrate a high temperature hardness of 34.4 GPa after heating at 900 °C, implying better hardness stability. The high temperature hardness contributes to the enhanced resistance to elastic deformations and improved wear resistance at high working temperatures. The three-period coatings exhibit thermally stable high resistance to elastic deformation, H/E, which decreases slightly from 0.1 at room temperature to 0.092 at 900 °C. Similarly, it shows good stability against plastic deformation, pronouncing with a small change from 0.37 GPa to 0.29 GPa observed in the H3/E*2 ratio at annealing from room temperature and 900 °C. The decrease in the resistance to the plastic deformations is consistent with the decreased wear resistance from 1.08 × 10−5 mm/N·m to 7.27 × 10−6 mm/N·m at elevated temperatures, implying good wear resistance. In the coatings with a smaller modulation period, a greater decrease in the hardness from 36.4 GPa to 29.7 GPa at the same thermal treatment was found, causing the relaxation of the greater internal stress due to the increased number of interfaces. That hardness behavior affects the mechanical and tribological properties of the coatings with a smaller modulation period. They exhibit less resistance to elastic and plastic deformations, which decrease significantly with a temperature increase. The tribological tests showed that the modulation period influences the coating’s behavior at increased temperatures differently due to the difference in the dominant wear mechanisms. An oxidation wear mechanism dominates the coatings with a large modulation period, and the wear rate decreases with a temperature increase, while, in coatings with a lower modulation period, abrasive wear is predominant, leading to an increasing wear rate.
The results from the investigation of the mechanical and tribological properties of TiAlSiN/AlSiN coatings with different modulation periods in dependence on the high treatment temperature determined the three-period TiAlSiN/AlSiN coating as being more suitable for high-temperature applications.
This study showed that the nanocomposite structure of the sublayers in the period benefits the thermal stability of the multilayer structure and composition. In addition, by combining a sublayer having a high hardness with a thicker layer with more elasticity, an improvement in the resistance to elastic and plastic deformations can be realized, i.e., the coating toughness can be enhanced. Furthermore, it was found that the thickness of the modulation period affects the wear mechanism. The increasing number of interfaces in the multilayer with a small modulation period promotes an increase in internal stress, causing instability of the mechanical properties expressed in a faster decrease in the hardness and resistance to elastic and plastic deformations. The smaller modulation period hampers the formation of oxides of sufficient thickness on the entire surface of the worn track, which affects the wear mechanism and, respectively, the wear rate.
The results of this study revealed that a beneficial approach in creating temperature-stable coatings is the combination of a multilayer architecture with nanocomposite structured sublayers of the period. In those coatings, enhanced hardness can be achieved, which results in improved mechanical and wear resistance at high temperatures.

Author Contributions

Conceptualization, L.K. and R.K.; Formal analysis, S.K., V.C. and T.C.; Investigation, L.K., S.K., D.K., G.A., V.C., T.C., E.Z. and C.P.; Methodology, L.K. and R.K.; Project administration, L.K. and R.K.; Resources, R.K.; Supervision, L.K.; Validation, L.K., R.K. and D.K.; Visualization, L.K. and E.Z.; Writing—original draft, S.K., L.K., D.K. and G.A.; Writing—review and editing, L.K., R.K. and D.K. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by the European Regional Development Fund within the OP “Research, Innovation and Digitalization Programme for Intelligent Transformation 2021–2027”, Project No. BG16RFPR002-1.014-0005 Center of competence “Smart Mechatronics, Eco- and Energy Saving Systems and Technologies”.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The data are not publicly available due to issues related to proprietary rights.

Acknowledgments

Research equipment of the Distributed Research Infrastructure INFRAMAT, part of the Bulgarian National Roadmap for Research Infrastructures, supported by the Bulgarian Ministry of Education and Science, was used in this investigation.

Conflicts of Interest

The authors declare no conflicts of interest.

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Coatings 15 00542 g001
Figure 1. Surface morphology of TiAlSiN/AlSiN coatings: (a) ML-3—as-deposited; (b) ML-3—annealed at 900 °C; (c) ML-30—as-deposited; (d) ML-30—annealed at 900 °C.
Figure 1. Surface morphology of TiAlSiN/AlSiN coatings: (a) ML-3—as-deposited; (b) ML-3—annealed at 900 °C; (c) ML-30—as-deposited; (d) ML-30—annealed at 900 °C.
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Figure 2. SEM images of a cross-section of TiAlSiN/AlSiN coatings: (a) ML-3—as-deposited; (b) ML-3—annealed at 900 °C; (c) ML-30—as-deposited; (d) ML-30—annealed at 900 °C.
Figure 2. SEM images of a cross-section of TiAlSiN/AlSiN coatings: (a) ML-3—as-deposited; (b) ML-3—annealed at 900 °C; (c) ML-30—as-deposited; (d) ML-30—annealed at 900 °C.
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Figure 3. Dependence of the element concentration of 3-period (a) and 30-period (b) TiAlSiN/AlSiN coatings on the annealing temperature.
Figure 3. Dependence of the element concentration of 3-period (a) and 30-period (b) TiAlSiN/AlSiN coatings on the annealing temperature.
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Figure 4. XPS spectra of the 3-periods (a) and 30-periods (b) TiAlSiN/AlSiN coatings.
Figure 4. XPS spectra of the 3-periods (a) and 30-periods (b) TiAlSiN/AlSiN coatings.
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Figure 5. XRD pattern of 3-period and 30-period multilayer TiAlSiN/AlSiN coatings (a) and extended pattern of the (200) peak (b).
Figure 5. XRD pattern of 3-period and 30-period multilayer TiAlSiN/AlSiN coatings (a) and extended pattern of the (200) peak (b).
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Figure 6. Dependence of the mechanical properties of TiAlSiN/AlSiN coatings on the annealing temperature: (a) hardness and elastic modulus; (b) resistance to elastic and plastic deformation.
Figure 6. Dependence of the mechanical properties of TiAlSiN/AlSiN coatings on the annealing temperature: (a) hardness and elastic modulus; (b) resistance to elastic and plastic deformation.
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Figure 7. A schematic diagram of the “ball/pin-on-disk” method.
Figure 7. A schematic diagram of the “ball/pin-on-disk” method.
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Figure 8. Graphs of the coefficient of friction of TiAlSiN/AlSIN coatings with 3 periods (a) and 30 periods (b) at different annealing temperatures.
Figure 8. Graphs of the coefficient of friction of TiAlSiN/AlSIN coatings with 3 periods (a) and 30 periods (b) at different annealing temperatures.
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Figure 9. Dependence of the coefficient of friction (a) and wear rate (b) on the thermal treatment of the TiAlSiN/AlSiN coatings with 3 and 30 modulation periods.
Figure 9. Dependence of the coefficient of friction (a) and wear rate (b) on the thermal treatment of the TiAlSiN/AlSiN coatings with 3 and 30 modulation periods.
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Coatings 15 00542 g010aCoatings 15 00542 g010b
Figure 10. Morphology, wear profiles, and element composition of the tribological tracks of TiAlSiN/AlSiN coatings annealed at different temperatures: (a) 3-period multilayers, (b) 30-period multilayers.
Figure 10. Morphology, wear profiles, and element composition of the tribological tracks of TiAlSiN/AlSiN coatings annealed at different temperatures: (a) 3-period multilayers, (b) 30-period multilayers.
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Table 1. Element composition of the 3-period and 30-period TiAlSiN/AlSiN coating surfaces.
Table 1. Element composition of the 3-period and 30-period TiAlSiN/AlSiN coating surfaces.
CoatingELEMENT COMPOSITION, at. %
-TiAlSiNCO
TiAlSiN/AlSiN—3 periods------
AD19.8024.905.6046.802.900.00
700 °C19.4025.106.6045.003.900.00
800 °C19.0026.006.2044.503.900.40
900 °C13.9029.406.9044.103.801.90
TiAlSiN/AlSiN—30 periods------
AD12.9030.807.3045.803.200.00
700 °C13.4031.407.3043.604.300.00
800 °C13.5032.507.5042.003.501.00
900 °C13.4031.707.4041.504.501.50
Table 2. Unit cell parameters of phases determined in TiAlSiN/AlSiN coatings.
Table 2. Unit cell parameters of phases determined in TiAlSiN/AlSiN coatings.
Sample/PhaseTiN
a = 4.24173 Å
PDF 00-038-1420		AlN
a = 3.114 Å c = 4.9792 Å
PDF 00-025-1133
TiAlSiN/AlSiN—3 periods--
AD4.2248 (6)a = 3.141 (3) c = 5.19 (1)
800 °C4.2283 (5)a = 3.154 (2) c = 5.20 (3)
900 °C4.2337 (4)a = 3.145 (2) c = 5.20 (3)
TiAlSiN/AlSiN—30 periods--
AD4.2386 (6)a = 3.154 (4) c = 5.19 (1)
800 °C4.2336 (8)a = 3.150 (6) c = 5.16 (2)
900 °C4.2359 (5)a = 3.145 (6) c = 5.13 (2)
Table 3. Critical loads of the as-deposited and annealed TiAlSiN/AlSiN coatings.
Table 3. Critical loads of the as-deposited and annealed TiAlSiN/AlSiN coatings.
Coatings LC1LC2LC3
F1, [N]l1, [mm]F2, [N]l2, [mm]F3, [N]l3, [mm]
TiAlSiN/AlSiN—3 periods------
AD56.91.68----
700 °C55.21.6377.82.32--
800 °C58.91.7966.01.9784.02.50
900 °C52.81.5655.51.6485.02.54
TiAlSiN/AlSiN—30 periods------
AD40.01.1854.41.63--
700 °C35.01.0057.51.7083.02.5
800 °C44.31.3060.01.7780.02.4
900 °C38.01.1056.01.6580.52.4
Table 4. Dependence of the width and depth of the wear tracks in TiAlSiN/AlSiN coatings on the annealing temperature.
Table 4. Dependence of the width and depth of the wear tracks in TiAlSiN/AlSiN coatings on the annealing temperature.
CoatingWear Track Width
(W), [µm]		Penetration Depth
(Pd), [µm]
TiAlSiN/AlSiN—3 bilayers
AD		3402.39
700 °C2802.58
800 °C2752.18
900 °C3101.76
TiAlSiN/AlSiN—30 bilayers
AD		2652.13
700 °C2852284
800 °C3052385
900 °C2602.46
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Kolchev, S.; Kolaklieva, L.; Kovacheva, D.; Atanasova, G.; Cholakova, T.; Chitanov, V.; Zlatareva, E.; Kakanakov, R.; Pashinski, C. Mechanical and Tribological Behavior of TiAlSiN/AlSiN Coatings Depending on the High-Temperature Treatment. Coatings 2025, 15, 542. https://doi.org/10.3390/coatings15050542

AMA Style

Kolchev S, Kolaklieva L, Kovacheva D, Atanasova G, Cholakova T, Chitanov V, Zlatareva E, Kakanakov R, Pashinski C. Mechanical and Tribological Behavior of TiAlSiN/AlSiN Coatings Depending on the High-Temperature Treatment. Coatings. 2025; 15(5):542. https://doi.org/10.3390/coatings15050542

Chicago/Turabian Style

Kolchev, Stefan, Lilyana Kolaklieva, Daniela Kovacheva, Genoveva Atanasova, Tetiana Cholakova, Vasiliy Chitanov, Ekaterina Zlatareva, Roumen Kakanakov, and Chavdar Pashinski. 2025. "Mechanical and Tribological Behavior of TiAlSiN/AlSiN Coatings Depending on the High-Temperature Treatment" Coatings 15, no. 5: 542. https://doi.org/10.3390/coatings15050542

APA Style

Kolchev, S., Kolaklieva, L., Kovacheva, D., Atanasova, G., Cholakova, T., Chitanov, V., Zlatareva, E., Kakanakov, R., & Pashinski, C. (2025). Mechanical and Tribological Behavior of TiAlSiN/AlSiN Coatings Depending on the High-Temperature Treatment. Coatings, 15(5), 542. https://doi.org/10.3390/coatings15050542

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