The effect of steel substrate pre-hardening on structural, mechanical, and tribological properties of magnetron sputtered TiN and TiAlN coatings

Wear 352-353 (2016) 92–101

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The effect of steel substrate pre-hardening on structural, mechanical, and tribological properties of magnetron sputtered TiN and TiAlN coatings F.F. Komarov a,n, V.M. Konstantinov b, A.V. Kovalchuk b, S.V. Konstantinov a, H.A. Tkachenko b a b

Institute of Applied Physics Problems of the Belarusian State University, 7 Kurchatov Str., Minsk, Belarus Belarusian National Technical University, Nezavisimosty Avenue 65, Minsk, Belarus

art ic l e i nf o

a b s t r a c t

Article history: Received 12 November 2015 Received in revised form 2 February 2016 Accepted 7 February 2016 Available online 12 February 2016

The effect of pre-hardening of steel substrate on mechanical and tribological properties of TiN and TiAlN coating–steel substrate composites was studied for various treatment conditions. We have found that prior nitrocarburization of the steel substrate increases the microhardness by 7 times, the wear resistance of the working surface of the composite up to 2.3 times, the resilience of the composite by 4.5 times. It is thus feasible to produce thinner coatings in composite layered systems with no drawbacks regarding performance. We also suggest a method to determine wear resistance of the deposited coatings, as well as to estimate the influence of the hardness gradient and contribution of all layers to overall durability of such layered systems. & 2016 Elsevier B.V. All rights reserved.

Keywords: Nanostructured coatings TiN TiAlN Magnetron sputtering Substrate pre-hardening Hardness and mechanical properties

1. Introduction Physical vapor-deposited (PVD) and chemical vapor-deposited (CVD) coatings have passed the scientific laboratory stage of study and are now widely used in various technical fields. Despite this, much of the research has been focused on the coating part of the coating/substrate materials system and less so on the substrate. There is a number of different ways to improve properties of these coatings. These include the creation of multi-component coatings based on high-entropy compounds [1–3]. Nanostructuring of coatings by technical methods and corresponding doping are also very effective [4–12]. Traditionally, much research and practical development are devoted to multilayer coatings [13]. These trends can definitely improve performance of the above-discussed coatings. Nevertheless, common for them is the tendency to enhance properties of only the coating. It should be noted, however, that performance of the working surface of a coated steel substrate or coated detail is not completely determined by properties of the coating. Furthermore, coating properties in their turn are not only determined by the n

Corresponding author. E-mail addresses: [email protected] (F.F. Komarov), [email protected] (V.M. Konstantinov), [email protected] (A.V. Kovalchuk), [email protected] (S.V. Konstantinov). http://dx.doi.org/10.1016/j.wear.2016.02.007 0043-1648/& 2016 Elsevier B.V. All rights reserved.

thickness, chemical composition and structure. The coating and steel substrate is a layered system for which these values do not represent independent parameters determining properties of the whole structure. Properties of the working surface of layered system will be determined by mechanical properties of the substrate and by the value of properties gradient between the coating and substrate [4,14]. It is also known, that in the case of formation of hard thin PVD and CVD coatings there is a sharp transition in hardness values at the substrate interface. If the surface to be coated has high plasticity and insufficient toughness, despite the high hardness, the coating will be sagged and collapsed under high specific loads during the friction with a counterbody. This destruction may occur due to reduction in bearing capacity of the layered system, i.e., at loads less than needed for the appearance of the plastic deformation of the substrate material [13,15–17]. As a hard coating material, TiN coatings exhibit high hardness and good wear resistance [4]. It is nowadays widely used to improve the wear resistance and durability of machining tools such as drills, cutting inserts and mills. However, TiN coatings suffer from severe disadvantages in high-temperature applications stemming from the fact that it oxidizes very rapidly at temperatures above 500 °C to form a layer that is composed of TiO2 which is a more brittle and less tough material [18]. Therefore, TiN will be considered here as a model system.

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Ti1
xAlxN coatings exhibit greatly enhanced high-temperature oxidation resistance due to the formation of a stable passive bilayer oxide [18]. The hardness of Ti1
xAlxN films increases with increasing Al content up to xmax of 0.65 and then a rapid decrease is observed [19–21]. Ti1
xAlxN ternary nitrides crystallize in metastable NaCl face-centered cubic (fcc) for the low Al contents and in ZnS hexagonal wurtzite structure (hcp) for the high Al content (xZ 0.65) [20–24]. These coatings with x Z0.5 are widely used in many applications including metal machining such as drilling, cutting, milling, die casting, extrusion etc. [25–28]. TiAlN coatings deposited on the AISI 304 steel and Armco-iron substrates with the preliminary carburizing of the surface layer have obviously revealed enhanced performance characteristics [17]. Among them, the purely plastic deformation of the coating occurred rather than its purely plastic fracture. The aim of this research is to shed more light on the effect of steel substrate pre-hardening by prior nitrocarburization on the properties of magnetron deposited TiN and TiAlN coatings and to evaluate the effect of coating-steel strengthening depending on the combination of previous complex treatments. Moreover, the present research is aimed at optimizing the coating–substrate system by studying the way in which the substrate can affect and improve the tribological characteristics of the entire coating/substrate system.

2. Experimental details The experiments were carried out using samples of TiN and TiAlN coatings on hardened and unhardened Armco-iron and AISI 304 steel substrates. The coatings were manufactured by means of magnetron sputtering using ArþN2 plasma discharges. The initial microstructure of steel substrates was in equilibrium due to the preannealing. Prior to deposition, grinding and polishing of substrates were performed. Then, the substrates were cleaned in boiling toluene. The surface roughness of substrates does not exceed Ra E0.1 μm. To harden the samples, thermochemical treatment (TCT) was used, namely, nitrocarburization in a mixture of potassium ferricyanide K4Fe(CN)6 (60%)þcharcoal C (30%)þchemical reaction activator BaCO3 (10%) powder. During the TCT processing that lasted 7 hours, the substrate temperature was kept at 550– 600 °C. The saturating environment in a working container was isolated from the ambient air by a sealing fusible gate. The inclined microsection with an angle of 45° was made to measure microhardness distribution in the depth of diffusion layer. The above-mentioned coatings were deposited using an URM 327 vacuum setup for magnetron sputtering equipped with an automatic system controlling supply of argon and nitrogen based on the S100 portable spectrometer (SOLAR Laser Systems, Belarus). We used 99.99% pure Ti and Ti–Al cathodes formed as disks with a diameter of 110 mm and up to 6 mm height. The base of the composite targets was made from a fine titanium powder and an extra-pure aluminum powder. The produced disks were compacted using explosion compression. Prior to deposition, all targets were sputter-cleaned for 10 min in the pure Ar plasma discharge, while the substrate was shielded by a shutter. To control the production of TiAl–TiAlNx–TiAlN coatings, we recorded time variation of spectral intensities of the atomic line of titanium, the bands of molecular and atomic nitrogen, as well as the atomic line of argon in discharge plasma. The ratio of intensities IN2/ITi of 0.33 corresponds to stoichiometric composition of TiAlN. Gas pressure of Arþ N2 was maintained at 0.7 Pa. No more than one target was used simultaneously in deposition process. The distance between cathode and substrate was 0.06 m. During deposition of TiN and TiAlN coatings, the substrate temperature was kept constant at 300 °C. Other conditions such as

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supply voltage, discharge current and bias on substrate were 300–320 V, 1.4–1.75 A,
90 V, respectively. The deposition rate was about 1.5 nm/s. The thickness of TiN and TiAlN coatings was selected between 1 and 5 μm with precision of 0.01 μm. To determine elemental composition of coatings, including nitrogen atoms depth distributions, the deposition of films on the substrates of commercially pure graphite was also performed. The composition of the coatings was measured by Rutherford backscattering spectrometry (RBS) using an AN-2500 accelerator (High Voltage Engineering Europe B.V.) by using 1.5 MeV He þ ions and detector with energy resolution of 15 keV. The obtained spectra were fitted using HEAD-6 and SIMNRA-6 software to get depth distributions of elements in the coatings. Furthermore, X-Ray Diffraction (XRD) and Transmission Electron Microscopy (TEM) investigations were employed for structural investigations by using a DRON-3 diffractometer operating in Bragg–Brentano configuration (λ ¼1.79021 Å) as well as a Hitachi-H800 electron microscope operating at 200 kV, respectively. The thickness of resulting coatings was measured by means of scanning electron microscope (SEM) «Mira» by «Tescan» together with an energy-dispersive X-ray spectrometer «INCA Energy 350» by «Oxford Instruments Analytical» (Great Britain). The relative error of the method did not exceed 5%. Mechanical properties of the coatings were evaluated by nanoindentation using a Nanoindentor G200 system (MES Systems, USA) and the Oliver–Pharr method implemented in this system [29]. A Berkovich diamond tip with 20 nm radius of curvature at the vertex was used. Ten indentations were made for each sample. The microhardness tests were performed by indentation with a recovered print of Vickers indenter using a DuraScan 20 microhardness tester (Switzerland). A comparative evaluation of coating adhesion was performed, using the test method of stretching out holes with special spherical punches by the Eriksen’s device (Eriksen's stamp) [30]. This method relates to technological tests and is used in order to evaluate material’s capability to withstand certain amount of residual strain under conditions corresponding to the industrial manufacturing (cold forming operations). As a measure of material’s ability to be stretched out, the elongated hole depth was used. Mechanical exercising on samples, bending radius and loading rate were kept constant. Adhesion quality was evaluated taking into account cracks and cleavings on the spherical part of the tested samples. With respect to a proper statistic, all the microhardness, nanoindentation and adhesion tests were performed at least 10 times. To measure wear resistance of the fabricated layered composites, a method of tribological testing [31] with a shaft-block scheme (see Fig. 1) was developed. The setup was mounted on a frame equipped with vibration dampers (see Fig. 2). A disc counterbody made of hard alloy HG20 (6 mass. % Co, 94 mass. % WC, hardness of 12 GPa) is rotated at a constant speed of 20 revolutions per minute (r. p. m.) by the engine through a gearbox. The sample, pre-balanced with a movable counterweight, is fixed to the disk by the calibration weights placed under the rotation axis. The following parameters and test conditions were employed: the diameter of counterbody was 0.06 m, the rotation speed was 0.063 m/s, the normal load was 0.1 N and the testing time ttest was 5400 s. The accuracy of linear wear control was 710 nm, an uncertainty of the sliding distance did not exceed 0.01 and the axial load was varied from 0.05 to 1.0070.01 N. With respect to a proper statistic, all the tribological tests were performed at least 3 times. Besides, the following assumptions were made in the subsequent data determination: wear rate of the coating and substrate is constant in the total duration of testing, thickness of the pre-hardened

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Fig. 1. Testing setup employed to measure wear resistance of layered composites [32]. Here R is the radius of disk; hi is the depth of wear track, and Ci is the value of linear wear equal to 2 AB.

Fig. 2. Schematic illustration of the setup employed for tribological test. 1 – testing sample, 2 – disk, 3 – weights, 4 – console, 5 – oscillation axis, 6 – engine, 7 – gearbox, 8 – rotation axis, 9 – lubricant, 10 – frame, 11 – vibration dampers.

sub-layer is much larger than the coating thickness and depth of the wear track. The discussed above testing conditions were selected allowing us to compare the present data with previously published results concerning properties of other coatings, as well as diffusion layers, obtained by thermochemical treatment. Values of linear wear of the employed substrates and these substrates with coatings, obtained under identical test conditions, were used in determining the wear resistance. The value of coating wear rate was obtained by comparing the wear rate of the coating on pre-hardened or not pre-hardened substrate with the wear rate of the uncoated substrates [32]. Wear tracks formed in dry-sliding tests on the composite surfaces were measured with an optical microscope. The following equations were used to calculate wear rates: V wear ¼ hwear = t test ;

ð1Þ

t test ¼ t coat: wear þ t sub: wear ¼ hcoat =V coat: wear þ ðhwear – hcoat Þ=V sub: wear ; ð2Þ were Vwear is the wear rate, hwear is the depth of wear tracks, tsub. wear is the time of substrate wear, tcoat. wear is the time of coating wear, hcoat is the thickness of coating, Vsub. wear is the substrate wear rate, and Vcoat. wear is the coating wear rate. The time of the coating wear was determined as time from the beginning of the wear test until the coating was worn out. The

Fig. 3. Rutherford backscattering spectra of carbon substrate with the deposited stoichiometric TiAlN coating. The analysis was performed with 1.5 MeV He ions.

time of the substrate wear was determined as time from the coating was worn out until the end of the wear test.

3. Results and discussion 3.1. Elemental compositions and microstructure Based on the collected RBS data, depth distributions of element composition in TiAlN and TiN coatings was calculated. No detectable traces of oxygen has been found in the coatings (see Fig. 3) deposited on both type of substrates, which indicates high quality of the obtained structures. Particularly, this means that no oxides are formed on the surface which could adversely influence the performance of the produced nano-scaled composite structures. Titanium and aluminum are present in the TiAlN coatings in approximately equal atomic concentrations (0.5 70.02) and their depth distributions are homogeneous in the coating layer. The content of titanium in the TiN coating is also uniform in the coating layer (the corresponding RBS spectrum is not presented). As for nitrogen, its depth distribution in TiN and TiAlN is also uniform apart from low gradients both at the surface and at the interface to the substrate. According to our previous results [17], it has been observed, that TiN and TiAlN coatings, deposited at the stoichiometric concentration of nitrogen, have revealed the best performance. An interface sublayer to TiAlN coatings represents an

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Fig. 4. X-ray diffraction patterns taken from the surfaces of (a) TiN and (b) TiAlN coatings deposited on the virgin AISI 304 stainless steel.

Fig. 5. TEM results for TiAlN coating deposited on the virgin AISI 304 stainless steel: (a) microphotograph, and (b) electron diffraction pattern.

adhesion sublayer of a 200 nm thick TiAl film. In the case of TiN coating, the adhesion sublayer is a 200 nm thick Ti film. Fig. 4 shows X-ray diffraction patterns taken from the surface of TiN and TiAlN coatings deposited on the virgin AISI 304 stainless steel. In the case of TiN coating deposited on the AISI 304 steel substrate, γ-Fe-peaks (from the substrate) and TiN-peaks are detected (Fig. 4a). The TiN crystalline structure matches that of the NaCl-B1 (fcc) type structure with a strong preferential orientation where the (110) planes are parallel to the substrate surface. In the case of TiAlN coatings deposited on the same substrate, the peaks of a complex nitride (Ti, Al)N are observed (Fig. 4b). It is known that both titanium nitride and aluminum nitride have isomorphic fcc lattices, which only differ in the values of their parameters. Therefore, these lattices form continuous series of solid solutions [4]. The atomic radius of aluminum is less than that of titanium [4]. If it is assumed that in the nitride phase aluminum atoms substitute titanium atoms to form a substitutional solid solution, the XRD-peaks of titanium nitrides will shift towards larger reflection angles. This situation is observed in the XRD-patterns taken from the TiAlN coatings (see Fig. 4b). The fact that the peaks of (TiAl)N-phase are highly diffuse and make up a band in the form of halo in the diffraction pattern is indicative of nanocrystalline structure of the TiAlN coatings. We estimated the crystallite sizes for TiN and TiAlN coatings deposited on the AISI 304 steel substrate by using the Scherrer’s equation. That is, we ignored in the first approximation the contribution of microstrains, but took into account the instrumental broadening.

The estimated average sizes of TiN and TiAlN crystallites are about 45–50 nm and 9.5–12.5 nm, respectively. Plan-view TEM revealed that crystallites of the obtained coatings are finely dispersed and form a homogeneous nanostructure. The observed strong broadening of the diffraction peaks in electron diffraction patterns indicates that the produced coatings are nanostructured (Fig. 5). As measured by TEM, crystallite size of the TiAlN coatings amounts to 10–15 nm (Fig. 5). The results, obtained by TEM, confirm the results of XRD analysis. Also we can state that residual stress in the substrate decreases after TCT, as this treatment was conducted at 550–600 °C thus leading to defect annealing [4]. It was not detected any difference in X-ray diffraction patterns and TEM images for coatings deposited on the virgin or pre-hardened steel substrates. As an example, Fig. 6 shows a cross-section structure of TiAlN coatings, deposited on the pre-hardened AISI 304 steel substrate. Thin transition and diffusion layers of 0.3–0.4 μm in total thickness between the TiAlN coatings and the steel substrate, formed during the reactive magnetron deposition by an incident flux of Ti þ and Al þ ions at the first stage of deposition, and Ti þ , Al þ and N þ ions after that, are observed. The thermal and radiationenhanced diffusion of the components of the coating and substrate, including the upward diffusion of nitrogen from the prehardened by TCT thick layer, may occurs in the process of coating deposition (temperature of the substrate exceeds 300 °C). No doubt, this layer improves the adhesion of coatings and mechanical properties of TiAlN coating–substrate composites such as microhardness, wear resistance and resilience. The corresponding

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Fig. 6. Cross-section image of TiAlN coating as obtained by means of scanning electron microscopy. Table 1 The hardness of steel substrates before and after nitrocarburization. Material

Armco-iron Steel AISI 304

Property Microhardness, MPa Before nitrocarburization

After nitrocarburization

1250 1520

1520 1840

Fig. 7. Depth distribution of microhardness in the diffusion layer.

SEM cross-section image for TiN system (not presented) behaves like that one. Moreover, a globular structure of the deposited film is obviously observed in this image. It is a very favorable characteristic to fabricate hard coatings with high functional properties [4]. 3.2. Microhardness tests of composite structures According to the results of microhardness tests, it was found that low-temperature nitrocarburization leads to higher microhardness of substrates up to 20% (Table 1) due to the formation of the inner zone of saturation (Fig. 7). The thickness of the hardened layer amounts to 80 μm and 120 μm for Armco-iron and AISI304 steel, respectively. A sharp decrease in microhardness (Fig. 7) is

attributed to an interface between the hardened layer and the substrate bulk. It should be noted that the substrate pre-hardening causes a substantial increase of microhardness for the substrate–coating systems and more strong increasing of microhardness with load decreasing (Figs. 8 and 9). This fact is indicative of the inevitable effect of the substrate and that the deposited coating works better and the effect of the carrying capacity reduction is not observed. It is obviously confirmed by the hardness values measured under the loads of 0.25 and 0.1 N which are not typical for these materials (Figs. 8 and 9). Therefore, the coating begins to make a more significant contribution to the effective hardness of the substrate– coating system. Moreover, such coatings are not depressed on tougher pre-hardened steel substrates. It was found that the effect of substrate hardening on the total hardness of the layered system is not additive one. So, the hardness reported in Figs. 8 and 9 is the composite hardness caused by deposited coating and steel substrate. The difference in the microhardness of coatings on Armco-iron and AISI 304 steel, on the average, amounts to 300 MPa. It increases to an average of 6000 MPa after hardening of the steel substrate, but the difference in substrate microhardness remains at nearly the same level as compared with that for microhardness of coatings. The observed phenomenon is apparently due to the fact, that the hardening of the substrate, resulting from the thermochemical treatment (TCT), eliminates the presence of a sharp boundary in hardness between the substrate and the coating, with this effect thereby dampening the hardness gradient of dissimilar materials, and significantly increases the carrying capacity of coating and the overall hardness of this layered system. Confirmation of these results is reflected by the loaddisplacement (loads 0–50 mN) curves for TiAlN coatings (not presented). The observed steepness of the unloading curve for the TiAlN coating, deposited on AISI 304 steel, indicates high elasticity of this coating, but the hardness value is relatively modest, and the coating is pressed under the indenter. According to these measurements, the hardness and elastic modulus were 28.07 GPa and 210 GPa, respectively, and the stiffness amounted to 2.47 N/m. It should be noted that the pre-hardening of steel substrate significantly changes the picture. An observed small angle of unloading curve in this case means high elasticity of the coating and mainly plastic type of destruction for higher indenter loads with no cracking. As measured in these experiments, the hardness and elastic modulus were 52.88 GPa and 350 GPa, respectively, and the stiffness amounted to 1.48 N/m. As reviewed by Leyland and Matthews [33], and Musil [34], the wear behavior is usually determined by H/E* ratio (here H and E* are the microhardness and the effective Young’s modulus, respectively; E* ¼E/(1
ν2), where E is the Young's modulus and ν is the Poisson's ratio). In this respect, a high H/E* ratio is desirable, as it characterizes the value of elastic recovery during unloading for nano- or micro-contact interaction. Hard coatings satisfying the ratio H/E* 4 0.1 exhibit enhanced resistance to plastic deformation and distribute the load applied to the coating over wider area that results in increase of the resistance of coating to cracking [34]. Such coatings are simultaneously hard and tough. The TiAlN coating deposited on virgin soft versus pre-hardened AISI304 substrate yields H/E* ratio of 0.122 and 0.138, respectively. The Poisson's ratio, ν was taken equal to 0.3 for the calculation [4]. This indicates that the coating on a harder substrate becomes more resilient and more resistant to plastic deformation. Therefore, it increases the load bearing capacity, crashworthiness, resilience and crack resistance. According to our estimates, the resilience of the substrate–TiAlN coating system increases by 4.55 times if the coating is deposited onto the substrate pre-hardened by TCT.

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Fig. 8. Microhardness of TiN coatings on various substrates versus applied load: (a) Armco iron; (b) Armco iron with TCT; (c) AISI 304; (d) AISI 304 with TCT. Table 2 Results of tribological tests for samples of Armco-iron.

Fig. 9. The increase of microhardness depending on the type of treatment: (a) without treatment; (b) TCT; (c) TiN coating deposition; (d) TCT and TiN coating deposition.

The hardening of steel substrate causes an increase by about 2 times in elastic recovery after the removal of indenter load, as it is observed from the reduced integrated area under the loaddisplacement curve measured for the coating deposited on the pre-hardened substrate as compared with the virgin AISI 304 steel substrate. 3.3. Wear behavior, coating adhesion and cracking resistance It should be noted that the wear of coatings on pre-hardened and virgin steel substrates exhibits very different nature. The coatings deposited on unhardened substrates were plastically deformed already at an initial stage of wear testing. While coatings on the pre-hardened substrate were worn uniformly through the whole duration of testing and at a stage of plastic deformation, they inhibited the development of wear tracks. The calculated results of wear rate indicate that the wear resistance of TiN coating deposited on pre-hardened Armco-iron substrate is 23.5% higher than that on unhardened Armco-iron.

Sample

Without treatment

After TCT With TiN coating

After TCT with TiN coating

Depth of wear track, mm Linear wear Ci, mm Volumetric wear, mm3,  103 Wear rate, Substrate mm/min, Coating

4.80

4.26

2.03

1.13

1073.3 1718.0

1010.7 1435.0

698.2 474.2

520.5 194.9

0.0533 –

0.0473 –

0.0533 0.0142

0.0473 0.0115

Table 3 Results of tribological tests for samples of AISI304 steel. Sample

Without treatment

After TCT With TiN coating

After TCT with TiN coating

Depth of wear track, mm Linear wear Ci, mm Volumetric wear, mm3,  103 Wear rate, Substrate mm/min, Coating

4.44

3.83

1.82

0.80

1032.2 1528.0

958.8 1228.1

660.2 296.1

437.0 117.4

0.0493 –

0.0426 –

0.0493 0.0136

0.0426 0.0089

The substrate's hardness increasing by 21.6% causes the wear resistance increase by 23.5%. The ratio of wear rate for the TiN coating on pre-hardened substrate to that on unhardened substrate amounts to 1.24 (Table 2). Similarly, the ratio of wear rate for TiN coating deposited on the pre-hardened AISI304 steel compared to that on unhardened AISI304 steel equals 1.53 (see Table 3). The data of calculating wear rate showed that the wear resistance of the TiN coating deposited

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on the hardened AISI 304 steel increased by 52.8% (Table 3). The increase in the substrate's hardness by 21.1% causes the wear resistance increasing by 52.8%. It means a not additive dependence of the formed system’s wear resistance on a type of treatment (Fig. 10). Thus, the summary effect of wear resistance improving from the pre-hardening of substrate by TCT and deposition of coating is much less, than the observed effect of wear resistance improving by the combination of TCT for substrate material and further deposition of vacuum coating Table 4. The observed phenomenon of non-additive effect of the substrate pre-hardening on the wear resistance of studied layered system can be explained as follows. Increasing the hardness of substrate by means of nitrocarburization allows to diminish wear until the appearance of permanent deformation in the substrate material. Moreover, this new quality permits to suppress significantly the residual compressive stress generated in the substrate, which prolongs a wearing time until the formation of cracks. A hardened substrate inhibits the development of the destruction zone after the beginning of sequential destruction. Then, after full destruction of coating under counterbody, the prehardened substrate works longer than unhardened one [35–39]. Also, the effect of changing wear mechanism takes place after the substrate nitrocarburization. In the case of the TiAlN coating deposited on the virgin AISI 304 steel substrate, the coating

Fig. 10. Increase of wear resistance dependent on the type of treatment: (a) no treatment; (b) TCT; (c) TiN coating; (d) TCT and TiN coating.

Table 4 Total length (μm) of surface cracks for TiAlN coatings deposited on virgin and prehardened AISI 304 steel substrates. Substrate/thickness of coating, mm Steel AISI 304 without pre-hardening Steel AISI 304 after nitrocarburization

3 258 10

5 110 35

crashes under the load of the counterbody during the wear test and the abrasive wear mechanism takes place due to the presence of wear products of TiAlN coating in the contact area. In the case of the TiAlN coating deposited on the pre-hardened AISI 304 steel substrate, there is no crushing or spalling of the coatings, thus, there is no abrasive wear. The fact of the non-additive effect of substrate nitrocarburization and coating deposition can be explained by eliminating the punching of the coating (effect of reducing the carrying capacity of the layered system [40,16]) as well as by pre-hardened substrate contribution to the elastic recovery response of the coating. Thus, as a result of a complex treatment, a new microcomposite material (topokomposite) [41,42] is formed on the surface, whose properties are unattainable for the substrate material and the coating separately. Obviously, the obtained layered composites should be characterized taking into account the adhesion properties of deposited coatings. Moreover, the required robustness of such layered systems cannot be achieved in practice without proper adhesion between the coating and the substrate. The coating adhesion was measured for both systems including the TiN coatings deposited on the pre-hardened and non-modified AISI 304 substrates. To set an upper value for punch penetration in stretching tests, at first, the sample of non-modified AISI 304 stainless steel was examined. The measured value of test bending by the Eriksen’s device for virgin AISI 304 steel amounted to 8.5 mm. The deposited coatings on non-modified substrate showed a complete delamination under examinations, i.e., the worst adhesion. This effect is observed on the spherical hole as well as nearby (Fig. 11). The plate with TiN coating, deposited onto the pre-hardened steel, kept its plasticity. Obviously, the adhesion strength and plasticity of the coating increased essentially. Under the test bending of 5.5 mm, the coating keeps completely the shape of hole (Fig. 12a), though the coating’s surface is covered with a net of cracks (see Fig. 12b), still it is well preserved on the substrate. There are no rings of delaminated coating around the spherical hole (Fig. 12b). This effect implies rather high adhesion between the coating and pre-hardened substrate. Based on the optical images of coatings, their cracking resistance was determined. The estimate was made by calculating the total length of cracks on the surface of coating, which appeared after the application of a load of 1.96 N (Fig. 13). This method was also discussed in Refs. [43,44]. The maximum total length of the cracks was observed for the specimen with TiAlN coating deposited on the AISI 304 steel (Fig. 13a). This coating has the lowest cracking resistance and thus is very susceptible regarding brittle fracture. The minimum total length of cracks was found for the sample with the TiAlN coating deposited on the AISI 304 steel pre-

Fig. 11. Eriksen's test results for a sample with TiN coating deposited on virgin AISI 304 stainless steel.

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Fig. 12. Eriksen's test results for a sample with TiN coating deposited on the pre-hardened AISI 304 stainless steel.

Fig. 13. Surface cracks for TiAlN coating that appeared after unloading indenter: (a) 3 μm TiAlN coating deposited on AISI 304 stainless steel; (b) 5 μm TiAlN coating deposited on AISI 304 steel; (c) 3 μm TiAIN coating deposited on pre-hardened by TCT AISI 304 steel; (d) 5 μm TiAlN coating deposited on pre-hardened by TCT AISI 304 steel.

hardened by means of TCT (Fig. 13c). This coating has the highest cracking resistance and is not prone to brittle fracture. A similar behavior of the cracking resistance is observed under different conditions of the loading of indenter (2.94 and 4.9 N).

Further, we have found that TiAlN coatings deposited on various types of substrates are characterised by different modes of fracture. On the AISI 304 stainless steel, the brittle or partly brittle fracture of the coating occurs, as a rule, as well as the spiral and

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circular cracks appear around the indent (Fig. 13a and b). On the pre-hardened substrate, the coating demonstrates more plastic fracture and the appearance of very short straight cracks, which propagate from the vertices of the indent (Fig. 13c and d) or no crack formation (at lower loadings). Thus, the coating on the softer substrate (AISI 304 stainless steel) is more prone to brittle fracture and has the lower cracking resistance than the coating on the harder substrate (TCT modified AISI 304 steel), which demonstrates predominantly plastic fracture. It should be noted that this conclusion is in accordance with a result in discussion of this problem in Ch. 12 of Ref. [4].

4. Conclusions

1. A scientific hypothesis on the influence of hardening of the steel substrate in the system «steel-PVD coating» on hardness and tribological properties of the coating is confirmed by the results of this study. The pre-hardening of ferritic and austenitic steel substrates by nitrocarburization improves both microhardness and wear resistance of vacuum coatings based on TiN and TiAlN. 2. We have found that nitrocarburization of the steel substrate enables to increase the microhardness up to 7 times, the wear resistance for working surfaces of composites by 2.3 times, and the resilience of the studied composites by more than 4.5 times. This allows to predict a science-based reduction in the thickness of the coatings in layered systems without loss of performance characteristics. 3. It was also observed that the effect of pre-hardening of the substrate in the systems «steel-PVD coating» by TCT and PVD coating deposition on the performance characteristics of working surfaces is not additive. 4. Substrate nitrocarburization substantially increases the toughness of coatings and their adhesion to the substrate, which leads to the complex enhancement of their performance characteristics including the wear and cracking resistances of the coatings. 5. With regard to the modes of coatings fracture under the loading, the coating deposited onto the virgin steel substrates showed brittle or partially brittle fracture with the formation of spiral and circular cracks around the indent. On the substrate pre-hardened by nitrocarburization, the coating demonstrates more plastic fracture and the appearance of the very small straight cracks, which propagate from the vertices of the indent, or no crack formation. This effect can be attributed to the formation of the optimal diffusion-controlled transition layer at the substrate–coating interface. 6. The discussed system of PVD coating þpre-hardened steel thick layer þsteel base is a new composite material, which is characterized by the performance properties unattainable by the coating or substrate individually. The conclusions of this work indicate the value of treating the coating and substrate as a material system in order to improve wear performance. This approach may also be applied to more complex layered material systems.

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