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Ti ratios
Journal of Alloys and Compounds 809 (2019) 151816
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Mechanical modification and damage mechanism evolution of TiN films subjected to cyclic nano-impact by adjusting N/Ti ratios Honghong Zhang a, Zeqing Li a, Weifeng He a, b, *, Chuansheng Ma c, **, Bin Liao d, Yinghong Li a, b a
State Key Laboratory for Manufacturing Systems Engineering, School of Mechanical Engineering, Xi'an Jiaotong University, Xi'an, 710049, China Science and Technology on Plasma Dynamics Laboratory, Air Force Engineering University, Xi'an, 710038, China Jia-Lab for Interface and Atomic Structure, Xi'an Jiaotong University, Xi'an, 710049, China d Institute of Low Energy Nuclear Physics, Beijing Normal University, Beijing, 100875, China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 April 2019 Received in revised form 2 August 2019 Accepted 10 August 2019 Available online 12 August 2019
For the purpose of optimizing the anti-impact performance of TiN film, a series of TiN films with different N/Ti ratios were investigated. The phase evolution and mechanical properties of TiN films were explored. In particular, the cyclic nano-impact tests with impact energy ranging from 0.1 mJ to 0.9 mJ were conducted to evaluate the anti-impact performance of TiN films. It was found that the non-stoichiometric phases of TiN0.30, Ti2N and TiN0.61 reduced with increasing N/Ti ratio in TiN films. The hardness of film increased with increasing stoichiometric TiN phase in films, while as an indicator of toughness, the H3/E2 ratio decreased. The anti-impact performance of TiN films displayed a close relationship with both hardness and H3/E2 ratio, in which a high enough hardness was the prerequisite of outstanding impact resistance, and H3/E2 ratio was an important factor affecting the damage mechanism. As the H3/E2 ratio decreased, the damage mechanism of TiN film transformed gradually from plastic fatigue damage to brittle fracture failure. Especially, the TiN film featuring N/Ti ratio of 0.780 (TiN-16) was not only hard enough to resist penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive anti-impact performance. © 2019 Elsevier B.V. All rights reserved.
Keywords: Mechanical modification Damage mechanism Cyclic nano-impact Anti-impact performance N/Ti ratio TiN film
1. Introduction The degradation of engineering components during service process is a very common phenomenon in various industrial systems. As an important and effective surface modification method, PVD hard films are often employed to enhance the lifetime of components [1e7]. In the past, quite a lot of efforts have been devoted to produce super-hard films. Whereas, as the hardness increase, films become more and more brittle, resulting in earlier failure [8e10]. Thereby, the super hardness alone is not sufficient to provide optimal performance, especially in applications of cutting [11e13], erosion [14], wear [15] and fatigue [16] where high strain
* Corresponding author. State Key Laboratory for Manufacturing Systems Engineering, School of Mechanical Engineering, Xi'an Jiaotong University, Xi'an, 710049, China. ** Corresponding author. Jia-Lab for Interface and Atomic Structure, Xi'an Jiaotong University, Xi'an, 710049, China. E-mail addresses: [email protected] (W. He), [email protected] (C. Ma). https://doi.org/10.1016/j.jallcom.2019.151816 0925-8388/© 2019 Elsevier B.V. All rights reserved.
rate impact may occur. Thus there is growing interest in developing hard-yet-tough films to improve the anti-impact performance of components. Particularly, the H3/E2 ratio has been widely considered as a qualitative indicator of film toughness. Xu et al. noted that higher H3/E2 ratio often resulted in improved fracture resistance [9]. Even in the solid particle erosion conditions, the films with higher H3/E2 ratio showed better erosion resistance [11]. Due to the combinational properties of exceptional chemical stability, eminent wear resistance and excellent sand erosion resistance, titanium nitride (TiN) film would be a promising candidate for anti-impact protection [17e19]. In general, the properties of TiN film are strongly dependent on the film architecture [20e23] and process parameters [24e27]. For example, by adjusting the nitrogen flow rate during deposition, the N/Ti ratio in TiN film could be varied [28e31], and since the N/Ti ratio occupies a crucial position in TiN film, the variation of nitrogen flow rate would obviously affect the properties of TiN films. J. Yang et al. deposited compositionally graded and multilayered TiN films by varying the amplitude and velocity of N2 gas injection, they
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reported that the hardness and electrical resistivity of TiN films increased with increasing amplitude and velocity of N2 gas [29]. L. Lu et al. studied the effect of N2 flow rate on the resistivity of TiNx films, and the resistivity of films firstly went up, and then dropped with increasing N2 flow rate [28]. J. Huang et al. investigated the properties of TiN films with N/Ti ratios ranging from 0.4 to 1.1, and they found the preferred orientation changed from (111) to (200), then back to (111) as the nitrogen flow rate increased [32]. Although considerable efforts have been devoted to the influence of nitrogen flow rate on the microstructure, mechanical properties and even electrical properties of TiN films, there are surprisingly few studies addressed on improving the anti-impact performance of TiN films by adjusting the N/Ti ratio in films so far. To investigate the impact behavior of thin films, a dynamic loading instrument developed by Knotek et al. would be a good assistant [33]. It is essentially a high-strain rate cyclic loading technique, in which the repeated high-strain rate (~103 s1) impact load is applied to the same location of films [34]. During testing, the damage evolution of films can be in-situ monitored by recording the impact depth of indenter [35]. Attributing to the characteristics of high depth resolution (<0.1 nm) and convenient in-situ monitoring, the nano-impact technique has been applied to evaluate the impact resistance of a wide range of films, including carbon films [35e37], metal films [38e40], and ceramic films [41e45]. In this study, a series of TiN films were fabricated by adjusting the nitrogen flow rate. The evolution of phase composition with N/ Ti ratio, and the variation of mechanical properties with phase constituents were explored. In particular, the cyclic nano-impact tests with impact energy ranging from 0.1 mJ to 0.9 mJ were conducted to evaluate the anti-impact performance of TiN films. Furthermore, particular emphases have been placed on analyzing the evolution of damage mechanism and the influence of mechanical properties on the anti-impact performance of TiN films. 2. Experimental methods 2.1. Deposition of TiN films Due to the wide applications of titanium alloys (Ti-6Al-4V) in aero-engine compressor blades, kinds of anti-erosion or anticorrosion films have been developed on the titanium alloys to improve their performance in services, such as TiN and TiAlN films. Since the N/Ti ratio is a crucial parameter of TiN films, by controlling the nitrogen flow rate (8 sccm, 16 sccm, 20 sccm and 26 sccm) into the vacuum chamber, the TiN films with different N/Ti ratios were fabricated on the titanium alloy (Ti-6Al-4V) substrates using a filtered cathodic vacuum arc (FCVA) system as described previously [2,46]. Before deposition, the substrates were ultrasonically cleaned with acetone and anhydrous alcohol for 20 min, respectively. During deposition, the substrates were placed on a holder with a rotation speed of 10 rpm, and the applied bias voltage of holder was 200 ± 1 V. The base pressure in vacuum chamber was ~1.0 103 Pa, and the arc current applied to Ti target was 100 ± 1 A. 2.2. Characterization techniques The relative atomic content of N and Ti atoms in films was determined by an X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific ESCALAB Xiþ). Before analysis, all films were etched by Arþ ion for 120 s to remove the contaminant on surface. The surface microstructure and roughness were analyzed by an atomic force microscope (AFM), and the cross-sectional and impact damage morphologies were investigated by a field emission scanning electron microscopy (FESEM, FEI Quanta 250).
The phase constituent of films was measured by a X-ray diffractometer (Bruker D8 Advance), in which the grazing incident mode with Cu Ka radiation at 40 kV and 40 mA was applied. A grazing incident angle of 2 and a scanning range of 30 e90 were conducted. To acquire stronger diffraction peaks, the step increment of 0.01 with a step scanning time of 0.5 s was applied. A nano-indenter (MTS Nano-indenter G200) was utilized to measure the nano-hardness (H) and Young's modulus (E) of films. The indentation depth was less than 10% of the film thickness to eliminate the influence from substrate. The values of H and E presented in this paper are the average of 12 individual indentations. To qualitatively evaluate the toughness of film, the parameter of H3/ E2 ratio was calculated. 2.3. Cyclic nano-impact test The cyclic nano-impact tests were conducted in the NanoTest system (Micro Materials Ltd., Wrexham, UK) with a multiple impulse impact module. Fig. 1(a) illustrates the schematic configuration of the nano-impact system. It is a pendulum-based depth sensing system, which can repeatedly impact the same position on sample surface with a regular frequency [38,47]. The solenoid connected to a timing relay and the periodic step electrical signal are responsible for generating the cyclic impacts on sample surface. Specifically speaking, at the beginning of test, the indenter indented into the sample quasistatically with a set load. When a step electrical signal is applied to the solenoid, through the connection of pendulum, the indenter is attracted by the solenoid to a set accelerating distance from the sample surface. While when the electrical signal turns off, the attraction between solenoid and pendulum disappears, and under the action of the set impact load, the indenter accelerates to impact the sample surface. Ignoring the work done by gravity, the impact energy per cycle is the product of the accelerating distance and the set impact load [41]. As shown in Fig. 1(b), the impact cycle in this study consisted of 2 s load on and 2 s load off. The total impact time for each test was 800 s, corresponding to 200 cyclic impacts at the same site. It is worth noting that during the impact test, the evolution of impact depth was in-situ tracked by recording the impact depth versus time. The accelerating distance was 20 mm and the impact load varied from 5 mN, 15 mN, 25 mN, 35 mN and 45 mN, thus the corresponding impact energy per cycle was 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ. Since cyclic impact is a fatigue process [8], repeated tests are requisite to acquire the statistical evaluation of the impact resistance of films. Thereby, at least five repeated impact tests were carried out at different sites for each load. 3. Results and discussions 3.1. Microstructure and composition To distinguish, the TiN films fabricated at different nitrogen flow rates would be hereafter referred to as TiN-8, TiN-16, TiN-20 and TiN-26, respectively. The relative atomic ratio of N/Ti in TiN film with respect to the nitrogen flow rate was illustrated in Fig. 2. Obviously, the N/Ti ratio in TiN films was sensitive to the nitrogen flow rate. For TiN-8 film, the N/Ti ratio was only 0.471, indicating the severe nitrogen deficiency in the film. As nitrogen flow rate rose to 16 sccm and even 20 sccm, more and more nitrogen atoms reacted with titanium atoms, leading to the steep increase of N/Ti ratio to 0.780 and 0.951, respectively. Whereas, with continued increasing nitrogen flow rate, more frequent collisions between atoms consumed more and more plasma energy, leaving less energy for nitrogen molecules to dissociate and react with titanium atoms [28,31], thus resulting in the slowly growth of N/Ti ratio from
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Fig. 1. Schematic configuration of the nano-impact system (a); nano-impact mode and typical data of impact depth versus time (b).
Fig. 2. Variation of Ra roughness and N/Ti ratio with nitrogen flow rate.
0.951 to 0.995. Fig. 2 also showed a nonlinear evolution of the surface roughness Ra with nitrogen flow rate. The Ra values of all films were less than 10 nm, and as the nitrogen flow rate increased, the Ra value decreased dramatically from 7.29 nm (TiN-8) to 4.57 nm (TiN-16) and 5.05 nm (TiN-20), but then it increased to 6.17 nm (TiN-26). In addition, the surface morphology of TiN films was displayed in Fig. 3. Unsurprisingly, with increasing nitrogen flow rate, a good agreement was found between the variation of surface morphology
and Ra roughness. The TiN-8 film (Fig. 3 (a)) showed a rough and nonuniform mounds-like structure. This should be attributed to the severe nitrogen deficiency and the coexistence of several TixNy (x > y) phases [2]. As nitrogen flow rate increased to 16 sccm and 20 sccm, there were appropriate nitrogen atoms to react with titanium atoms, leading to the efficient synthesis of TiN phase, so the mounds-like structure became more smooth and uniform (Fig. 3 (b)e(c)). Nevertheless, with the continuous increase of nitrogen flow rate, the considerable energy consumed by the frequent collisions between atoms reduced the deposition rate and bombardment energy, which was unfavorable to form a densely packed crystalline structure [31,32,48,49]. Accordingly, the surface morphology displayed a tendency of roughing (Fig. 3 (d)). Fig. 4 presented the FESEM cross-sectional micrographs of the as-deposited films. It can be seen that the thickness of TiN-8, TiN16, TiN-20 and TiN-26 films were 5.478 mm, 5.676 mm, 5.304 mm and 5.238 mm, respectively. Thus the average thickness of all films was 5.424 mm with a deviation of just 0.196 mm, indicating that all films have a similar thickness due to the same deposition time of an hour.
3.2. Phase compositions and mechanical properties The GIXRD patterns and phase constituent analysis were presented in Fig. 5. It is obvious that the phase constituents of TiN-8 film were distinctive to those of other films. Lacking of TiN phase, the TiN-8 film was mainly composed of TiN0.30 phase and a small quantity of Ti2N phase. This should be ascribed to the severe nitrogen deficiency in TiN-8 film [2]. While for TiN-16 film, due to the sharply increased N/Ti ratio in the film, the severe non-
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Fig. 3. AFM graphs for surface morphology of (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN-26 films.
Fig. 4. Cross- sectional micrographs of (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN-26 films.
stoichiometric phases of TiN0.30 and Ti2N phases disappeared. The typical peaks of (111), (200), (220), (311) and (222) indicated the formation of TiN phase. However, ascribing to the coexistence of
TiN0.61 and TiN phases in the film, the diffraction patterns overlapped to some extent, thus the diffraction peaks of TiN-16 film were broad and asymmetric. For those diffraction peaks, we firstly
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Fig. 5. GIXRD patterns and phase constituents of as-deposited films.
carried out the process of peak-resolving and fitting, then the ICDD/ JCPDS PDF retrievals were conducted on the resolved peaks. With the disappearance of TiN0.61 phase, there was only TiN phase in the TiN-20 and TiN-26 films, and the diffraction peaks became symmetric and sharp. It's worth noting that when compared with the standard card of TiN, the TiN-20 film exhibited a preferred growth orientation of TiN (111). Fig. 6 illustrated the hardness, elastic modulus and H3/E2 ratio of as-deposited films. It was observed that attributing to the absence of stoichiometric TiN phase, the TiN-8 film showed the absolutely lower H and E of 16.99 GPa and 185.41 GPa (Fig. 6 (a)) than other films. As the stoichiometric TiN phase gradually generated in the TiN-16 film, both H and E showed a significant increase to 25.53 GPa and 345.51 GPa, respectively. For the TiN-20 and TiN-26 films, because of the disappearance of non-stoichiometric TixNy (x > y) phases, the hardness further grew to 29.98 GPa and 28.32 GPa, respectively. To qualitatively evaluate the toughness of as-deposited films, the H3/E2 ratios of as-deposited films were obtained and presented in Fig. 6 (b). Unlike the variation trend of H and E, the H3/E2 ratio showed a monotone tendency of decrease with increasing N/Ti ratio. Due to the existence of large amount of non-stoichiometric
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TixNy (x > y) phases, the TiN-8 and TiN-16 films possessed relatively higher H3/E2 ratios of 0.143 GPa and 0.140 GPa, respectively. As the non-stoichiometric phases disappeared, the H3/E2 ratios of TiN-20 and TiN-26 films obviously decreased to 0.128 GPa and 0.114 GPa, respectively. It can be inferred that the mechanical properties of TiN films could be modified by adjusting the phase constituents in films. At ambient temperature and normal pressure, TiN crystallizes in a face-centered cubic (fcc) structure, in which Ti atoms occupy the angular tops and face centers of the lattice, while N atoms fill the edge centers of the cube [50e53]. In the stoichiometric TiN phase, the titanium atoms and nitrogen atoms are bonded by covalent and ionic bonds [50]. However, in the non-stoichiometric TixNy (x > y) phases, due to the severe deficiency of nitrogen atoms, a certain number of N-vacancies will form in the lattice. Guemmaz et al. reported that the metallic bond of Ti-Ti, absent in the stoichiometric TiN phase, would established between the neighboring Ti atoms through N vacancies [51,52]. While when compared with the ionic and covalent bonds, the metallic bond usually possesses lower bonding strength, but better ductility. So the presence of nonstoichiometric TixNy (x > y) phases decreased the hardness, while enhanced the toughness of TiN films. This is in accordance with the researches by Qiao et al. [53]. In addition to the presence of non-stoichiometric TixNy (x > y) phases, the preferred growth orientation also plays an important role in the hardness of TiN films. Since the preferred growth orientation of TiN (111) is the close-packed plane of TiN phase, in which the atoms are closely packed with strong bonding force [54]. Thus the TiN-20 film exhibited denser microstructure (Fig. 3) and slightly higher hardness (Fig. 6 (a)), ultimately resulting in a higher H3/E2 ratio in comparison with TiN-26 film. 3.3. Damage mechanism of films in cyclic nano-impact tests In the cyclic nano-impact tests, the evolution of impact depth vs. impact times was in-situ monitored. Fig. 7 presented some depth evolution curves of the as-deposited films, which were continuously impacted by 200 times with single impact energy ranging from 0.1 mJ to 0.9 mJ. It was observed that during the initial few impacts, all films showed a similar behavior of a rapid increase in depth. This may be attributed to the fact that the initial contact stress was greater than the yield strength of films, causing the instantaneous plastic deformation or crack initiation in films [55]. As the impact continued, obvious differences occurred in the depth evolution. In order to reveal the damage mechanism of TiN films with different N/Ti ratios under the repeated nano-impact
Fig. 6. Mechanical properties of TiN films: (a) nano-hardness and elastic modulus, (b) H3/E2 ratio.
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Fig. 7. Evolution of impact depth vs. impact times in the cyclic nano-impact test with single impact energy ranging from 0.1 mJ to 0.9 mJ: (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN26 film.
condition, the morphology of corresponding impact craters was displayed in Fig. 8. Obviously, except for the initial impact stage, all the depth curves of TiN-8 and TiN-16 films increased gradually with increasing impact cycles, and the final impact depth increased with increasing impact energy from 0.1 mJ to 0.9 mJ. Meanwhile, the impact craters in Fig. 8 became larger and larger as the impact load increased. Moreover, the higher magnification image in Fig. 9 (a) showed plastic fatigue damage with lots of furrows in the craters and plenty of strip-like removals distributing around the craters [56]. Whereas, the TiN-20 film exhibited two types of depth evolution curves at different impact loads. For the impact energy ranging from 0.1 mJ to 0.5 mJ, the impact depth grew gradually and the corresponding craters displayed plastic fatigue damage. However, as the impact energy increased to 0.7 mJ and even 0.9 mJ, there was a sudden mutation in depth occurring at the 155th and 120th impact, respectively. Correspondingly, the higher magnification morphology in Fig. 9(b) illustrated typical brittle fracture failure with multiple brittle spallations and severe fragmentations distributing on the periphery of craters [56]. Similarly, the TiN-26 film also presented two types of damage mechanism. However, the impact energy threshold leading to brittle fracture failure was only 0.3 mJ, and during the tests with impact energies of 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ, the impact depth increased abruptly at the 135th, 52nd, 30th and 15th impact, respectively. As expected, the corresponding morphology of craters showed severe brittle fracture failure with quite extensive spallations and fragments. Intriguingly, Fig. 6 showed that the TiN-20 and TiN-26 films exhibited higher hardness than the TiN-16 film. However, at the
impact energy of 0.9 mJ, the impact depths of TiN-20 and TiN-26 films before fracture were much higher than that of TiN-16 film, as shown in Fig. 7. This should be explained by the fact that the sudden mutation in depth curve is an indication of brittle fracture, which means that the film has already experienced crack initiation and propagation processes. Accordingly, due to the severe brittleness of TiN-20 and TiN-26 film, cracks may have initiated during the initial few impacts, thus resulting in the higher initial impact depth than TiN-16 film. It can be informed that attributing to the unpredictability of crack initiation, propagation and even fracture, once brittle fracture occurs, both of the initial impact depth and the final impact depth would probably no longer show regular relationship with the hardness of films. In addition, by calculating the product of single impact energy and the impact cycles before fracture (Eq. (1)), the ability of absorbing impact energy before fracture, which refers to the toughness of films was obtained.
Ea ¼ Es ,N
(1)
Where, Ea , Es and N refers to the absorbed energy before fracture, the single impact energy and the impact duration before fracture, respectively. Herein, the work done by gravity and the heat loss during impact were ignored. Since the TiN-8 and TiN-16 films never showed any fracture phenomenon during the 200 times impact, even at the impact energy of 0.9 mJ, the TiN-8 and TiN-16 films could absorb impact energy of at least 0.18 mJ before fracture. The TiN-20 film showed depth mutations at the 155th and 120th impact
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Fig. 8. Morphologies of the as-deposited films after impacted by 200 cycles with single impact energy ranging from 0.1 mJ to 0.9 mJ.
for the tests with impact energy of 0.7 mJ and 0.9 mJ. Correspondingly, the impact energy absorbed by TiN-20 before fracture was about 0.1085 mJ and 0.108 mJ, respectively. Meanwhile, as the single impact energy increased, the impact energy absorbed by TiN-26 film before fracture was about 0.04 mJ, 0.026 mJ, 0.021 mJ and 0.014 mJ, respectively. It indicates that for the films with severe brittleness, the impact energy absorbed by film before fracture will decrease with increasing impact load. Furthermore, a good agreement was found between the ability of absorbing impact energy before fracture and the H3/E2 ratio of films. Thus it was intuitively demonstrated that the H3/E2 ratio was an evaluation indicator of toughness [9]. Hereafter, the toughness of films would be qualitatively compared by the values of H3/E2 ratio. Accordingly, it was informed that in the case of repeated nanoimpact, the damage mechanism of TiN films transformed gradually from plastic fatigue damage to brittle fracture failure with the increase of N/Ti ratio in films. The evolution of plastic fatigue damage can be described as follows: (1) during the initial few impacts, due to the small contact area between the indenter and film, the contact
stress was much greater than the yield strength of the film, so the film underwent significant plastic deformation and the impact depth increased dramatically; (2) With the increase of impact times, the contact area increased gradually, thus the contact stress decreased, which slowed down the growth rate of impact depth; (3) When the contact stress decreased to lower than the yield stress of film, the impact depth was approximately constant. Meanwhile, there were large amounts of dislocations and slides in the plastic deformation zone, and the high density of dislocations would inevitably produce blockings, leading to the work hardening [57,58]. (4) After a fatigue period, the local material would gradually remove in the forms of small sheets or strips, as shown in Fig. 9 (a). For the brittle fracture failure, since some of the films had very limited toughness, cracks were prone to initiate during initial few impacts. As the impact continued, the cracks would propagate and coalescence, and finally resulting in the brittle fracture reflected by the sudden increase in depth. Furthermore, our previous researches demonstrated that during impact, the extremely high tensile stresses, which could easily lead to the cracking of ceramic
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Fig. 9. Two typical damage mechanisms of TiN films subjecting to cyclic impact load.
materials, always distributing on the periphery of contact region [46], thus the brittle spallation mainly occurred around the impact craters, as shown in Fig. 9 (b). 3.4. Influence of mechanical properties on the anti-impact performance It can be seen from the impact crater morphologies in Fig. 8 that even under the same impact condition, both the damage mechanism and the size of impact craters were different among films. In order to provide a basis for the mechanical optimization of impact resistant TiN films in the future, the influence of mechanical properties on the anti-impact performance of TiN films was investigated. Firstly, dividing the number of brittle fracture events by the total number of tests under the same loading condition, the fracture percentage of films were calculated and graphed in Fig. 10. Possessing the relatively higher toughness with H3/E2 ratios of 0.143 GPa and 0.140 GPa, respectively, the fracture percentage of both TiN-8 and TiN-16 films throughout the impact energy ranging from 0.1 mJ to 0.9 mJ, was always 0%, indicating no brittle fracture event happened. Whereas, owning the lower toughness with H3/E2 ratio of 0.128 GPa, the TiN-20 film showed a critical impact energy leading to brittle fracture of 0.7 mJ. As expected, the TiN-26 film featuring the worst toughness with H3/E2 ratio of 0.114 GPa performed a so severe brittle failure tendency that the critical impact energy of brittle fracture was only 0.3 mJ, and even all of the five repeated tests at the impact energy of 0.9 mJ showed brittle fracture. It was further demonstrated that in the case of repeated nanoimpact, the film with higher toughness, the greater the critical impact energy causing brittle fracture, the lower the possibility of brittle fracture, and the more impact cycles it could withstand.
Fig. 10. Fracture percentage of the as-deposited TiN films at the single impact energy of 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ.
Furthermore, Fig. 11 displays the average final impact depth of the as-deposited TiN films under different impact loads. It was obvious that as the impact load increased, all films showed an increase in the final impact depth. Herein, as shown in Fig. 11, the full range of impact load was divided into the low load (0.1 mJ), medium load (0.3 mJe0.5 mJ) and high load (0.7 mJe0.9 mJ). In the impact condition of low load, since no brittle fracture event occurred in any films, the final impact depths of all films were small. Since hardness is known to be the ability of materials to prevent them from being pressed in by a hard object, the final
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of mechanical properties on the anti-impact performance of TiN films were evaluated by cyclic nano-impact tests.
Fig. 11. Average final impact depth of the as- deposited TiN films from the five repeated tests under the single impact energy of 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ.
impact depth of films showed a negative correlation with their hardness that the TiN-20 film featuring the highest hardness showed the smallest final impact depth, followed by TiN 26 and TiN 16 films in turn, while the TiN-8 film exhibited the largest final impact depth. When impacted by medium load, owing to the occurrence of brittle fracture event, the final impact depth of TiN26 film jumped to the maximum value among all films. Meanwhile, the final impact depth rank of other films was the same as that at low load. As the impact energy increased to 0.7 mJ and even higher, both the TiN-20 and TiN-26 films suffered from brittle fracture with different probabilities, thus the final impact depth of them increased obviously to be higher than that of TiN-16 film. It is worth noting that at high load, although the TiN-8 film never underwent brittle fracture, its final impact depth increased evidently to be far greater than that of other three films. This is probably attributed to that the TiN-8 film possessed much lower hardness, which severely limited the ability to prevent it from being pressed in by the indenter, thus resulting in the excessive impact depth. It can be reasonably believed that the anti-impact performance of TiN films is complicated related to the mechanical properties of hardness and toughness. As an indicator of toughness, H3/E2 ratio is an important factor affecting the damage mechanism of TiN film, and the higher H3/E2 ratio, the lower probability of sudden brittle failure. While when no brittle failure occurs, the final impact depth of TiN films was negatively correlated with their hardness, namely, the greater the hardness, the smaller the impact depth. Therefore, sufficient hardness is the prerequisite to improve the impact resistance of TiN films. While for the film with enough high hardness, too low H3/E2 ratio is one of the main factors weakening the impact resistance of TiN films. Above all, the TiN-16 film featuring an appropriate hardness of 25.53 GPa and a higher H3/E2 ratio of 0.140 GPa was not only hard enough to resist the indenter's penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive impact resistance. 4. Conclusions In the current study, the TiN films with different N/Ti ratios have been prepared by controlling nitrogen flow rate. The evolution of phase composition with N/Ti ratio was investigated, and the variation of mechanical properties with phase constituent was further revealed. Furthermore, the damage mechanism and the influence
1) By adjusting the phase constituents in TiN films, the variation of N/Ti ratios could modify the hardness and toughness of TiN films. The hardness increased with increasing TiN phase in films, while as an indicator of toughness, the H3/E2 ratio decreased. 2) In the cyclic nano-impact tests, the damage mechanism of TiN films transformed gradually from plastic fatigue damage to brittle fracture failure with decreasing H3/E2 ratio. 3) The anti-impact performance of TiN films was closely related to both hardness and toughness. High enough hardness was the prerequisite of outstanding impact resistance, and H3/E2 ratio was an important factor affecting the damage mechanism of TiN films. For the films with sufficient hardness, the higher toughness was conducive to lowering the probability of sudden brittle failure, thus resulting in better impact resistance. 4) The TiN-16 film featuring an appropriate hardness of 25.53 GPa and a higher H3/E2 ratio of 0.140 GPa was not only hard enough to resist the indenter's penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive impact resistance. Acknowledgements This work was financially supported by the Key Research and Development Program of Shaanxi Province (Grant No. 2017ZDXMGY-048) and the National Science and Technology Major Project (Grant No. 2017-Ⅶ-0012-0107). We appreciate Mr Huang at instrument Analysis Center of Xi'an Jiaotong University for his assistance with XRD analysis. Also, we thanks Professor Chen Jian and Dr. Shi Xiangru for their help in cyclic nano-impact tests. References [1] V. Merie, M. Pustan, G. Negrea, C.B. Rleanu, Research on titanium nitride thin films deposited by reactive magnetron sputtering for MEMS applications, Appl. Surf. Sci. (2015) 525e532. [2] H. Zhang, Z. Li, C. Ma, W. He, X. Cao, Y. Li, The anti-sand erosion performance of TiN films fabricated by filtered cathodic vacuum arc technique at different nitrogen flow rates, Ceram. Int. 45 (2019) 10819e10825. [3] A.D. Pogrebnjak, V.M. Beresnev, O.V. Bondar, B.O. Postolnyi, K. Zaleski, E. Coy, et al., Superhard CrN/Mon coatings with multilayer architecture, Mater. Des. 153 (2018) 47e59. [4] Q. Yang, R. McKellar, Nanolayered CrAlTiN and multilayered CrAlTiNeAlTiN coatings for solid particle erosion protection, Tribol. Int. 83 (2015) 12e20. [5] D.E. Wolfe, B.M. Gabriel, M.W. Reedy, Nanolayer (Ti,Cr)N coatings for hard particle erosion resistance, Surf. Coat. Technol. 205 (2011) 4569e4576. [6] E. Bemporad, M. Sebastiani, C. Pecchio, S. De Rossi, High thickness Ti/TiN multilayer thin coatings for wear resistant applications, Surf. Coat. Technol. 201 (2006) 2155e2165. [7] L. Swad Ba, A. Maciejny, B. Formanek, P. Liberski, P. Podolski, B. Mendala, et al., Influence of coatings obtained by PVD on the properties of aircraft compressor blades, Surf. Coat. Technol. 78 (1996) 137e143. [8] J. Chen, H. Li, B.D. Beake, Load sensitivity in repetitive nano-impact testing of TiN and AlTiN coatings, Surf. Coat. Technol. 308 (2016) 289e297. [9] Y.X. Xu, H. Riedl, D. Holec, L. Chen, Y. Du, P.H. Mayrhofer, Thermal stability and oxidation resistance of sputtered Ti Al Cr N hard coatings, Surf. Coat. Technol. 324 (2017) 48e56. [10] M. Bartosik, C. Rumeau, R. Hahn, Z.L. Zhang, P.H. Mayrhofer, Fracture toughness and structural evolution in the TiAlN system upon annealing, Sci. Rep.UK 7 (2017). [11] S. Hassani, M. Bielawski, W. Beres, L. Martinu, M. Balazinski, J.E. KlembergSapieha, Predictive tools for the design of erosion resistant coatings, Surf. Coat. Technol. 203 (2008) 204e210. [12] K.D. Bouzakis, G. Skordaris, E. Bouzakis, P. Charalampous, T. Kotsanis, D. Tasoulas, et al., Effect of the interface fatigue strength of NCD coated hardmetal inserts on their cutting performance in milling, Diam. Relat. Mater. 59 (2015) 80e89. [13] K.D. Bouzakis, N. Michailidis, E. Bouzakis, G. Katirtzoglou, S. Makrimallakis, S. Gerardis, et al., Cutting performance of coated tools with various adhesion strength quantified by inclined impact tests, CIRP Annals 60 (2011) 105e108. [14] E. Bousser, L. Martinu, J.E. Klemberg-Sapieha, Solid particle erosion mechanisms of protective coatings for aerospace applications, Surf. Coat. Technol.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Mechanical modification and damage mechanism evolution of TiN films subjected to cyclic nano-impact by adjusting N/Ti ratios Honghong Zhang a, Zeqing Li a, Weifeng He a, b, *, Chuansheng Ma c, **, Bin Liao d, Yinghong Li a, b a
State Key Laboratory for Manufacturing Systems Engineering, School of Mechanical Engineering, Xi'an Jiaotong University, Xi'an, 710049, China Science and Technology on Plasma Dynamics Laboratory, Air Force Engineering University, Xi'an, 710038, China Jia-Lab for Interface and Atomic Structure, Xi'an Jiaotong University, Xi'an, 710049, China d Institute of Low Energy Nuclear Physics, Beijing Normal University, Beijing, 100875, China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 April 2019 Received in revised form 2 August 2019 Accepted 10 August 2019 Available online 12 August 2019
For the purpose of optimizing the anti-impact performance of TiN film, a series of TiN films with different N/Ti ratios were investigated. The phase evolution and mechanical properties of TiN films were explored. In particular, the cyclic nano-impact tests with impact energy ranging from 0.1 mJ to 0.9 mJ were conducted to evaluate the anti-impact performance of TiN films. It was found that the non-stoichiometric phases of TiN0.30, Ti2N and TiN0.61 reduced with increasing N/Ti ratio in TiN films. The hardness of film increased with increasing stoichiometric TiN phase in films, while as an indicator of toughness, the H3/E2 ratio decreased. The anti-impact performance of TiN films displayed a close relationship with both hardness and H3/E2 ratio, in which a high enough hardness was the prerequisite of outstanding impact resistance, and H3/E2 ratio was an important factor affecting the damage mechanism. As the H3/E2 ratio decreased, the damage mechanism of TiN film transformed gradually from plastic fatigue damage to brittle fracture failure. Especially, the TiN film featuring N/Ti ratio of 0.780 (TiN-16) was not only hard enough to resist penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive anti-impact performance. © 2019 Elsevier B.V. All rights reserved.
Keywords: Mechanical modification Damage mechanism Cyclic nano-impact Anti-impact performance N/Ti ratio TiN film
1. Introduction The degradation of engineering components during service process is a very common phenomenon in various industrial systems. As an important and effective surface modification method, PVD hard films are often employed to enhance the lifetime of components [1e7]. In the past, quite a lot of efforts have been devoted to produce super-hard films. Whereas, as the hardness increase, films become more and more brittle, resulting in earlier failure [8e10]. Thereby, the super hardness alone is not sufficient to provide optimal performance, especially in applications of cutting [11e13], erosion [14], wear [15] and fatigue [16] where high strain
* Corresponding author. State Key Laboratory for Manufacturing Systems Engineering, School of Mechanical Engineering, Xi'an Jiaotong University, Xi'an, 710049, China. ** Corresponding author. Jia-Lab for Interface and Atomic Structure, Xi'an Jiaotong University, Xi'an, 710049, China. E-mail addresses: [email protected] (W. He), [email protected] (C. Ma). https://doi.org/10.1016/j.jallcom.2019.151816 0925-8388/© 2019 Elsevier B.V. All rights reserved.
rate impact may occur. Thus there is growing interest in developing hard-yet-tough films to improve the anti-impact performance of components. Particularly, the H3/E2 ratio has been widely considered as a qualitative indicator of film toughness. Xu et al. noted that higher H3/E2 ratio often resulted in improved fracture resistance [9]. Even in the solid particle erosion conditions, the films with higher H3/E2 ratio showed better erosion resistance [11]. Due to the combinational properties of exceptional chemical stability, eminent wear resistance and excellent sand erosion resistance, titanium nitride (TiN) film would be a promising candidate for anti-impact protection [17e19]. In general, the properties of TiN film are strongly dependent on the film architecture [20e23] and process parameters [24e27]. For example, by adjusting the nitrogen flow rate during deposition, the N/Ti ratio in TiN film could be varied [28e31], and since the N/Ti ratio occupies a crucial position in TiN film, the variation of nitrogen flow rate would obviously affect the properties of TiN films. J. Yang et al. deposited compositionally graded and multilayered TiN films by varying the amplitude and velocity of N2 gas injection, they
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reported that the hardness and electrical resistivity of TiN films increased with increasing amplitude and velocity of N2 gas [29]. L. Lu et al. studied the effect of N2 flow rate on the resistivity of TiNx films, and the resistivity of films firstly went up, and then dropped with increasing N2 flow rate [28]. J. Huang et al. investigated the properties of TiN films with N/Ti ratios ranging from 0.4 to 1.1, and they found the preferred orientation changed from (111) to (200), then back to (111) as the nitrogen flow rate increased [32]. Although considerable efforts have been devoted to the influence of nitrogen flow rate on the microstructure, mechanical properties and even electrical properties of TiN films, there are surprisingly few studies addressed on improving the anti-impact performance of TiN films by adjusting the N/Ti ratio in films so far. To investigate the impact behavior of thin films, a dynamic loading instrument developed by Knotek et al. would be a good assistant [33]. It is essentially a high-strain rate cyclic loading technique, in which the repeated high-strain rate (~103 s1) impact load is applied to the same location of films [34]. During testing, the damage evolution of films can be in-situ monitored by recording the impact depth of indenter [35]. Attributing to the characteristics of high depth resolution (<0.1 nm) and convenient in-situ monitoring, the nano-impact technique has been applied to evaluate the impact resistance of a wide range of films, including carbon films [35e37], metal films [38e40], and ceramic films [41e45]. In this study, a series of TiN films were fabricated by adjusting the nitrogen flow rate. The evolution of phase composition with N/ Ti ratio, and the variation of mechanical properties with phase constituents were explored. In particular, the cyclic nano-impact tests with impact energy ranging from 0.1 mJ to 0.9 mJ were conducted to evaluate the anti-impact performance of TiN films. Furthermore, particular emphases have been placed on analyzing the evolution of damage mechanism and the influence of mechanical properties on the anti-impact performance of TiN films. 2. Experimental methods 2.1. Deposition of TiN films Due to the wide applications of titanium alloys (Ti-6Al-4V) in aero-engine compressor blades, kinds of anti-erosion or anticorrosion films have been developed on the titanium alloys to improve their performance in services, such as TiN and TiAlN films. Since the N/Ti ratio is a crucial parameter of TiN films, by controlling the nitrogen flow rate (8 sccm, 16 sccm, 20 sccm and 26 sccm) into the vacuum chamber, the TiN films with different N/Ti ratios were fabricated on the titanium alloy (Ti-6Al-4V) substrates using a filtered cathodic vacuum arc (FCVA) system as described previously [2,46]. Before deposition, the substrates were ultrasonically cleaned with acetone and anhydrous alcohol for 20 min, respectively. During deposition, the substrates were placed on a holder with a rotation speed of 10 rpm, and the applied bias voltage of holder was 200 ± 1 V. The base pressure in vacuum chamber was ~1.0 103 Pa, and the arc current applied to Ti target was 100 ± 1 A. 2.2. Characterization techniques The relative atomic content of N and Ti atoms in films was determined by an X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific ESCALAB Xiþ). Before analysis, all films were etched by Arþ ion for 120 s to remove the contaminant on surface. The surface microstructure and roughness were analyzed by an atomic force microscope (AFM), and the cross-sectional and impact damage morphologies were investigated by a field emission scanning electron microscopy (FESEM, FEI Quanta 250).
The phase constituent of films was measured by a X-ray diffractometer (Bruker D8 Advance), in which the grazing incident mode with Cu Ka radiation at 40 kV and 40 mA was applied. A grazing incident angle of 2 and a scanning range of 30 e90 were conducted. To acquire stronger diffraction peaks, the step increment of 0.01 with a step scanning time of 0.5 s was applied. A nano-indenter (MTS Nano-indenter G200) was utilized to measure the nano-hardness (H) and Young's modulus (E) of films. The indentation depth was less than 10% of the film thickness to eliminate the influence from substrate. The values of H and E presented in this paper are the average of 12 individual indentations. To qualitatively evaluate the toughness of film, the parameter of H3/ E2 ratio was calculated. 2.3. Cyclic nano-impact test The cyclic nano-impact tests were conducted in the NanoTest system (Micro Materials Ltd., Wrexham, UK) with a multiple impulse impact module. Fig. 1(a) illustrates the schematic configuration of the nano-impact system. It is a pendulum-based depth sensing system, which can repeatedly impact the same position on sample surface with a regular frequency [38,47]. The solenoid connected to a timing relay and the periodic step electrical signal are responsible for generating the cyclic impacts on sample surface. Specifically speaking, at the beginning of test, the indenter indented into the sample quasistatically with a set load. When a step electrical signal is applied to the solenoid, through the connection of pendulum, the indenter is attracted by the solenoid to a set accelerating distance from the sample surface. While when the electrical signal turns off, the attraction between solenoid and pendulum disappears, and under the action of the set impact load, the indenter accelerates to impact the sample surface. Ignoring the work done by gravity, the impact energy per cycle is the product of the accelerating distance and the set impact load [41]. As shown in Fig. 1(b), the impact cycle in this study consisted of 2 s load on and 2 s load off. The total impact time for each test was 800 s, corresponding to 200 cyclic impacts at the same site. It is worth noting that during the impact test, the evolution of impact depth was in-situ tracked by recording the impact depth versus time. The accelerating distance was 20 mm and the impact load varied from 5 mN, 15 mN, 25 mN, 35 mN and 45 mN, thus the corresponding impact energy per cycle was 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ. Since cyclic impact is a fatigue process [8], repeated tests are requisite to acquire the statistical evaluation of the impact resistance of films. Thereby, at least five repeated impact tests were carried out at different sites for each load. 3. Results and discussions 3.1. Microstructure and composition To distinguish, the TiN films fabricated at different nitrogen flow rates would be hereafter referred to as TiN-8, TiN-16, TiN-20 and TiN-26, respectively. The relative atomic ratio of N/Ti in TiN film with respect to the nitrogen flow rate was illustrated in Fig. 2. Obviously, the N/Ti ratio in TiN films was sensitive to the nitrogen flow rate. For TiN-8 film, the N/Ti ratio was only 0.471, indicating the severe nitrogen deficiency in the film. As nitrogen flow rate rose to 16 sccm and even 20 sccm, more and more nitrogen atoms reacted with titanium atoms, leading to the steep increase of N/Ti ratio to 0.780 and 0.951, respectively. Whereas, with continued increasing nitrogen flow rate, more frequent collisions between atoms consumed more and more plasma energy, leaving less energy for nitrogen molecules to dissociate and react with titanium atoms [28,31], thus resulting in the slowly growth of N/Ti ratio from
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Fig. 1. Schematic configuration of the nano-impact system (a); nano-impact mode and typical data of impact depth versus time (b).
Fig. 2. Variation of Ra roughness and N/Ti ratio with nitrogen flow rate.
0.951 to 0.995. Fig. 2 also showed a nonlinear evolution of the surface roughness Ra with nitrogen flow rate. The Ra values of all films were less than 10 nm, and as the nitrogen flow rate increased, the Ra value decreased dramatically from 7.29 nm (TiN-8) to 4.57 nm (TiN-16) and 5.05 nm (TiN-20), but then it increased to 6.17 nm (TiN-26). In addition, the surface morphology of TiN films was displayed in Fig. 3. Unsurprisingly, with increasing nitrogen flow rate, a good agreement was found between the variation of surface morphology
and Ra roughness. The TiN-8 film (Fig. 3 (a)) showed a rough and nonuniform mounds-like structure. This should be attributed to the severe nitrogen deficiency and the coexistence of several TixNy (x > y) phases [2]. As nitrogen flow rate increased to 16 sccm and 20 sccm, there were appropriate nitrogen atoms to react with titanium atoms, leading to the efficient synthesis of TiN phase, so the mounds-like structure became more smooth and uniform (Fig. 3 (b)e(c)). Nevertheless, with the continuous increase of nitrogen flow rate, the considerable energy consumed by the frequent collisions between atoms reduced the deposition rate and bombardment energy, which was unfavorable to form a densely packed crystalline structure [31,32,48,49]. Accordingly, the surface morphology displayed a tendency of roughing (Fig. 3 (d)). Fig. 4 presented the FESEM cross-sectional micrographs of the as-deposited films. It can be seen that the thickness of TiN-8, TiN16, TiN-20 and TiN-26 films were 5.478 mm, 5.676 mm, 5.304 mm and 5.238 mm, respectively. Thus the average thickness of all films was 5.424 mm with a deviation of just 0.196 mm, indicating that all films have a similar thickness due to the same deposition time of an hour.
3.2. Phase compositions and mechanical properties The GIXRD patterns and phase constituent analysis were presented in Fig. 5. It is obvious that the phase constituents of TiN-8 film were distinctive to those of other films. Lacking of TiN phase, the TiN-8 film was mainly composed of TiN0.30 phase and a small quantity of Ti2N phase. This should be ascribed to the severe nitrogen deficiency in TiN-8 film [2]. While for TiN-16 film, due to the sharply increased N/Ti ratio in the film, the severe non-
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Fig. 3. AFM graphs for surface morphology of (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN-26 films.
Fig. 4. Cross- sectional micrographs of (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN-26 films.
stoichiometric phases of TiN0.30 and Ti2N phases disappeared. The typical peaks of (111), (200), (220), (311) and (222) indicated the formation of TiN phase. However, ascribing to the coexistence of
TiN0.61 and TiN phases in the film, the diffraction patterns overlapped to some extent, thus the diffraction peaks of TiN-16 film were broad and asymmetric. For those diffraction peaks, we firstly
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Fig. 5. GIXRD patterns and phase constituents of as-deposited films.
carried out the process of peak-resolving and fitting, then the ICDD/ JCPDS PDF retrievals were conducted on the resolved peaks. With the disappearance of TiN0.61 phase, there was only TiN phase in the TiN-20 and TiN-26 films, and the diffraction peaks became symmetric and sharp. It's worth noting that when compared with the standard card of TiN, the TiN-20 film exhibited a preferred growth orientation of TiN (111). Fig. 6 illustrated the hardness, elastic modulus and H3/E2 ratio of as-deposited films. It was observed that attributing to the absence of stoichiometric TiN phase, the TiN-8 film showed the absolutely lower H and E of 16.99 GPa and 185.41 GPa (Fig. 6 (a)) than other films. As the stoichiometric TiN phase gradually generated in the TiN-16 film, both H and E showed a significant increase to 25.53 GPa and 345.51 GPa, respectively. For the TiN-20 and TiN-26 films, because of the disappearance of non-stoichiometric TixNy (x > y) phases, the hardness further grew to 29.98 GPa and 28.32 GPa, respectively. To qualitatively evaluate the toughness of as-deposited films, the H3/E2 ratios of as-deposited films were obtained and presented in Fig. 6 (b). Unlike the variation trend of H and E, the H3/E2 ratio showed a monotone tendency of decrease with increasing N/Ti ratio. Due to the existence of large amount of non-stoichiometric
5
TixNy (x > y) phases, the TiN-8 and TiN-16 films possessed relatively higher H3/E2 ratios of 0.143 GPa and 0.140 GPa, respectively. As the non-stoichiometric phases disappeared, the H3/E2 ratios of TiN-20 and TiN-26 films obviously decreased to 0.128 GPa and 0.114 GPa, respectively. It can be inferred that the mechanical properties of TiN films could be modified by adjusting the phase constituents in films. At ambient temperature and normal pressure, TiN crystallizes in a face-centered cubic (fcc) structure, in which Ti atoms occupy the angular tops and face centers of the lattice, while N atoms fill the edge centers of the cube [50e53]. In the stoichiometric TiN phase, the titanium atoms and nitrogen atoms are bonded by covalent and ionic bonds [50]. However, in the non-stoichiometric TixNy (x > y) phases, due to the severe deficiency of nitrogen atoms, a certain number of N-vacancies will form in the lattice. Guemmaz et al. reported that the metallic bond of Ti-Ti, absent in the stoichiometric TiN phase, would established between the neighboring Ti atoms through N vacancies [51,52]. While when compared with the ionic and covalent bonds, the metallic bond usually possesses lower bonding strength, but better ductility. So the presence of nonstoichiometric TixNy (x > y) phases decreased the hardness, while enhanced the toughness of TiN films. This is in accordance with the researches by Qiao et al. [53]. In addition to the presence of non-stoichiometric TixNy (x > y) phases, the preferred growth orientation also plays an important role in the hardness of TiN films. Since the preferred growth orientation of TiN (111) is the close-packed plane of TiN phase, in which the atoms are closely packed with strong bonding force [54]. Thus the TiN-20 film exhibited denser microstructure (Fig. 3) and slightly higher hardness (Fig. 6 (a)), ultimately resulting in a higher H3/E2 ratio in comparison with TiN-26 film. 3.3. Damage mechanism of films in cyclic nano-impact tests In the cyclic nano-impact tests, the evolution of impact depth vs. impact times was in-situ monitored. Fig. 7 presented some depth evolution curves of the as-deposited films, which were continuously impacted by 200 times with single impact energy ranging from 0.1 mJ to 0.9 mJ. It was observed that during the initial few impacts, all films showed a similar behavior of a rapid increase in depth. This may be attributed to the fact that the initial contact stress was greater than the yield strength of films, causing the instantaneous plastic deformation or crack initiation in films [55]. As the impact continued, obvious differences occurred in the depth evolution. In order to reveal the damage mechanism of TiN films with different N/Ti ratios under the repeated nano-impact
Fig. 6. Mechanical properties of TiN films: (a) nano-hardness and elastic modulus, (b) H3/E2 ratio.
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Fig. 7. Evolution of impact depth vs. impact times in the cyclic nano-impact test with single impact energy ranging from 0.1 mJ to 0.9 mJ: (a) TiN-8, (b) TiN-16, (c) TiN-20 and (d) TiN26 film.
condition, the morphology of corresponding impact craters was displayed in Fig. 8. Obviously, except for the initial impact stage, all the depth curves of TiN-8 and TiN-16 films increased gradually with increasing impact cycles, and the final impact depth increased with increasing impact energy from 0.1 mJ to 0.9 mJ. Meanwhile, the impact craters in Fig. 8 became larger and larger as the impact load increased. Moreover, the higher magnification image in Fig. 9 (a) showed plastic fatigue damage with lots of furrows in the craters and plenty of strip-like removals distributing around the craters [56]. Whereas, the TiN-20 film exhibited two types of depth evolution curves at different impact loads. For the impact energy ranging from 0.1 mJ to 0.5 mJ, the impact depth grew gradually and the corresponding craters displayed plastic fatigue damage. However, as the impact energy increased to 0.7 mJ and even 0.9 mJ, there was a sudden mutation in depth occurring at the 155th and 120th impact, respectively. Correspondingly, the higher magnification morphology in Fig. 9(b) illustrated typical brittle fracture failure with multiple brittle spallations and severe fragmentations distributing on the periphery of craters [56]. Similarly, the TiN-26 film also presented two types of damage mechanism. However, the impact energy threshold leading to brittle fracture failure was only 0.3 mJ, and during the tests with impact energies of 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ, the impact depth increased abruptly at the 135th, 52nd, 30th and 15th impact, respectively. As expected, the corresponding morphology of craters showed severe brittle fracture failure with quite extensive spallations and fragments. Intriguingly, Fig. 6 showed that the TiN-20 and TiN-26 films exhibited higher hardness than the TiN-16 film. However, at the
impact energy of 0.9 mJ, the impact depths of TiN-20 and TiN-26 films before fracture were much higher than that of TiN-16 film, as shown in Fig. 7. This should be explained by the fact that the sudden mutation in depth curve is an indication of brittle fracture, which means that the film has already experienced crack initiation and propagation processes. Accordingly, due to the severe brittleness of TiN-20 and TiN-26 film, cracks may have initiated during the initial few impacts, thus resulting in the higher initial impact depth than TiN-16 film. It can be informed that attributing to the unpredictability of crack initiation, propagation and even fracture, once brittle fracture occurs, both of the initial impact depth and the final impact depth would probably no longer show regular relationship with the hardness of films. In addition, by calculating the product of single impact energy and the impact cycles before fracture (Eq. (1)), the ability of absorbing impact energy before fracture, which refers to the toughness of films was obtained.
Ea ¼ Es ,N
(1)
Where, Ea , Es and N refers to the absorbed energy before fracture, the single impact energy and the impact duration before fracture, respectively. Herein, the work done by gravity and the heat loss during impact were ignored. Since the TiN-8 and TiN-16 films never showed any fracture phenomenon during the 200 times impact, even at the impact energy of 0.9 mJ, the TiN-8 and TiN-16 films could absorb impact energy of at least 0.18 mJ before fracture. The TiN-20 film showed depth mutations at the 155th and 120th impact
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Fig. 8. Morphologies of the as-deposited films after impacted by 200 cycles with single impact energy ranging from 0.1 mJ to 0.9 mJ.
for the tests with impact energy of 0.7 mJ and 0.9 mJ. Correspondingly, the impact energy absorbed by TiN-20 before fracture was about 0.1085 mJ and 0.108 mJ, respectively. Meanwhile, as the single impact energy increased, the impact energy absorbed by TiN-26 film before fracture was about 0.04 mJ, 0.026 mJ, 0.021 mJ and 0.014 mJ, respectively. It indicates that for the films with severe brittleness, the impact energy absorbed by film before fracture will decrease with increasing impact load. Furthermore, a good agreement was found between the ability of absorbing impact energy before fracture and the H3/E2 ratio of films. Thus it was intuitively demonstrated that the H3/E2 ratio was an evaluation indicator of toughness [9]. Hereafter, the toughness of films would be qualitatively compared by the values of H3/E2 ratio. Accordingly, it was informed that in the case of repeated nanoimpact, the damage mechanism of TiN films transformed gradually from plastic fatigue damage to brittle fracture failure with the increase of N/Ti ratio in films. The evolution of plastic fatigue damage can be described as follows: (1) during the initial few impacts, due to the small contact area between the indenter and film, the contact
stress was much greater than the yield strength of the film, so the film underwent significant plastic deformation and the impact depth increased dramatically; (2) With the increase of impact times, the contact area increased gradually, thus the contact stress decreased, which slowed down the growth rate of impact depth; (3) When the contact stress decreased to lower than the yield stress of film, the impact depth was approximately constant. Meanwhile, there were large amounts of dislocations and slides in the plastic deformation zone, and the high density of dislocations would inevitably produce blockings, leading to the work hardening [57,58]. (4) After a fatigue period, the local material would gradually remove in the forms of small sheets or strips, as shown in Fig. 9 (a). For the brittle fracture failure, since some of the films had very limited toughness, cracks were prone to initiate during initial few impacts. As the impact continued, the cracks would propagate and coalescence, and finally resulting in the brittle fracture reflected by the sudden increase in depth. Furthermore, our previous researches demonstrated that during impact, the extremely high tensile stresses, which could easily lead to the cracking of ceramic
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Fig. 9. Two typical damage mechanisms of TiN films subjecting to cyclic impact load.
materials, always distributing on the periphery of contact region [46], thus the brittle spallation mainly occurred around the impact craters, as shown in Fig. 9 (b). 3.4. Influence of mechanical properties on the anti-impact performance It can be seen from the impact crater morphologies in Fig. 8 that even under the same impact condition, both the damage mechanism and the size of impact craters were different among films. In order to provide a basis for the mechanical optimization of impact resistant TiN films in the future, the influence of mechanical properties on the anti-impact performance of TiN films was investigated. Firstly, dividing the number of brittle fracture events by the total number of tests under the same loading condition, the fracture percentage of films were calculated and graphed in Fig. 10. Possessing the relatively higher toughness with H3/E2 ratios of 0.143 GPa and 0.140 GPa, respectively, the fracture percentage of both TiN-8 and TiN-16 films throughout the impact energy ranging from 0.1 mJ to 0.9 mJ, was always 0%, indicating no brittle fracture event happened. Whereas, owning the lower toughness with H3/E2 ratio of 0.128 GPa, the TiN-20 film showed a critical impact energy leading to brittle fracture of 0.7 mJ. As expected, the TiN-26 film featuring the worst toughness with H3/E2 ratio of 0.114 GPa performed a so severe brittle failure tendency that the critical impact energy of brittle fracture was only 0.3 mJ, and even all of the five repeated tests at the impact energy of 0.9 mJ showed brittle fracture. It was further demonstrated that in the case of repeated nanoimpact, the film with higher toughness, the greater the critical impact energy causing brittle fracture, the lower the possibility of brittle fracture, and the more impact cycles it could withstand.
Fig. 10. Fracture percentage of the as-deposited TiN films at the single impact energy of 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ.
Furthermore, Fig. 11 displays the average final impact depth of the as-deposited TiN films under different impact loads. It was obvious that as the impact load increased, all films showed an increase in the final impact depth. Herein, as shown in Fig. 11, the full range of impact load was divided into the low load (0.1 mJ), medium load (0.3 mJe0.5 mJ) and high load (0.7 mJe0.9 mJ). In the impact condition of low load, since no brittle fracture event occurred in any films, the final impact depths of all films were small. Since hardness is known to be the ability of materials to prevent them from being pressed in by a hard object, the final
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of mechanical properties on the anti-impact performance of TiN films were evaluated by cyclic nano-impact tests.
Fig. 11. Average final impact depth of the as- deposited TiN films from the five repeated tests under the single impact energy of 0.1 mJ, 0.3 mJ, 0.5 mJ, 0.7 mJ and 0.9 mJ.
impact depth of films showed a negative correlation with their hardness that the TiN-20 film featuring the highest hardness showed the smallest final impact depth, followed by TiN 26 and TiN 16 films in turn, while the TiN-8 film exhibited the largest final impact depth. When impacted by medium load, owing to the occurrence of brittle fracture event, the final impact depth of TiN26 film jumped to the maximum value among all films. Meanwhile, the final impact depth rank of other films was the same as that at low load. As the impact energy increased to 0.7 mJ and even higher, both the TiN-20 and TiN-26 films suffered from brittle fracture with different probabilities, thus the final impact depth of them increased obviously to be higher than that of TiN-16 film. It is worth noting that at high load, although the TiN-8 film never underwent brittle fracture, its final impact depth increased evidently to be far greater than that of other three films. This is probably attributed to that the TiN-8 film possessed much lower hardness, which severely limited the ability to prevent it from being pressed in by the indenter, thus resulting in the excessive impact depth. It can be reasonably believed that the anti-impact performance of TiN films is complicated related to the mechanical properties of hardness and toughness. As an indicator of toughness, H3/E2 ratio is an important factor affecting the damage mechanism of TiN film, and the higher H3/E2 ratio, the lower probability of sudden brittle failure. While when no brittle failure occurs, the final impact depth of TiN films was negatively correlated with their hardness, namely, the greater the hardness, the smaller the impact depth. Therefore, sufficient hardness is the prerequisite to improve the impact resistance of TiN films. While for the film with enough high hardness, too low H3/E2 ratio is one of the main factors weakening the impact resistance of TiN films. Above all, the TiN-16 film featuring an appropriate hardness of 25.53 GPa and a higher H3/E2 ratio of 0.140 GPa was not only hard enough to resist the indenter's penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive impact resistance. 4. Conclusions In the current study, the TiN films with different N/Ti ratios have been prepared by controlling nitrogen flow rate. The evolution of phase composition with N/Ti ratio was investigated, and the variation of mechanical properties with phase constituent was further revealed. Furthermore, the damage mechanism and the influence
1) By adjusting the phase constituents in TiN films, the variation of N/Ti ratios could modify the hardness and toughness of TiN films. The hardness increased with increasing TiN phase in films, while as an indicator of toughness, the H3/E2 ratio decreased. 2) In the cyclic nano-impact tests, the damage mechanism of TiN films transformed gradually from plastic fatigue damage to brittle fracture failure with decreasing H3/E2 ratio. 3) The anti-impact performance of TiN films was closely related to both hardness and toughness. High enough hardness was the prerequisite of outstanding impact resistance, and H3/E2 ratio was an important factor affecting the damage mechanism of TiN films. For the films with sufficient hardness, the higher toughness was conducive to lowering the probability of sudden brittle failure, thus resulting in better impact resistance. 4) The TiN-16 film featuring an appropriate hardness of 25.53 GPa and a higher H3/E2 ratio of 0.140 GPa was not only hard enough to resist the indenter's penetration, but also tough enough to prevent the film from fracturing, thus it exhibited the best comprehensive impact resistance. Acknowledgements This work was financially supported by the Key Research and Development Program of Shaanxi Province (Grant No. 2017ZDXMGY-048) and the National Science and Technology Major Project (Grant No. 2017-Ⅶ-0012-0107). We appreciate Mr Huang at instrument Analysis Center of Xi'an Jiaotong University for his assistance with XRD analysis. Also, we thanks Professor Chen Jian and Dr. Shi Xiangru for their help in cyclic nano-impact tests. References [1] V. Merie, M. Pustan, G. Negrea, C.B. Rleanu, Research on titanium nitride thin films deposited by reactive magnetron sputtering for MEMS applications, Appl. Surf. Sci. (2015) 525e532. [2] H. Zhang, Z. Li, C. Ma, W. He, X. Cao, Y. Li, The anti-sand erosion performance of TiN films fabricated by filtered cathodic vacuum arc technique at different nitrogen flow rates, Ceram. Int. 45 (2019) 10819e10825. [3] A.D. Pogrebnjak, V.M. Beresnev, O.V. Bondar, B.O. Postolnyi, K. Zaleski, E. Coy, et al., Superhard CrN/Mon coatings with multilayer architecture, Mater. Des. 153 (2018) 47e59. [4] Q. Yang, R. McKellar, Nanolayered CrAlTiN and multilayered CrAlTiNeAlTiN coatings for solid particle erosion protection, Tribol. Int. 83 (2015) 12e20. [5] D.E. Wolfe, B.M. Gabriel, M.W. Reedy, Nanolayer (Ti,Cr)N coatings for hard particle erosion resistance, Surf. Coat. Technol. 205 (2011) 4569e4576. [6] E. Bemporad, M. Sebastiani, C. Pecchio, S. De Rossi, High thickness Ti/TiN multilayer thin coatings for wear resistant applications, Surf. Coat. Technol. 201 (2006) 2155e2165. [7] L. Swad Ba, A. Maciejny, B. Formanek, P. Liberski, P. Podolski, B. Mendala, et al., Influence of coatings obtained by PVD on the properties of aircraft compressor blades, Surf. Coat. Technol. 78 (1996) 137e143. [8] J. Chen, H. Li, B.D. Beake, Load sensitivity in repetitive nano-impact testing of TiN and AlTiN coatings, Surf. Coat. Technol. 308 (2016) 289e297. [9] Y.X. Xu, H. Riedl, D. Holec, L. Chen, Y. Du, P.H. Mayrhofer, Thermal stability and oxidation resistance of sputtered Ti Al Cr N hard coatings, Surf. Coat. Technol. 324 (2017) 48e56. [10] M. Bartosik, C. Rumeau, R. Hahn, Z.L. Zhang, P.H. Mayrhofer, Fracture toughness and structural evolution in the TiAlN system upon annealing, Sci. Rep.UK 7 (2017). [11] S. Hassani, M. Bielawski, W. Beres, L. Martinu, M. Balazinski, J.E. KlembergSapieha, Predictive tools for the design of erosion resistant coatings, Surf. Coat. Technol. 203 (2008) 204e210. [12] K.D. Bouzakis, G. Skordaris, E. Bouzakis, P. Charalampous, T. Kotsanis, D. Tasoulas, et al., Effect of the interface fatigue strength of NCD coated hardmetal inserts on their cutting performance in milling, Diam. Relat. Mater. 59 (2015) 80e89. [13] K.D. Bouzakis, N. Michailidis, E. Bouzakis, G. Katirtzoglou, S. Makrimallakis, S. Gerardis, et al., Cutting performance of coated tools with various adhesion strength quantified by inclined impact tests, CIRP Annals 60 (2011) 105e108. [14] E. Bousser, L. Martinu, J.E. Klemberg-Sapieha, Solid particle erosion mechanisms of protective coatings for aerospace applications, Surf. Coat. Technol.
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