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ZrN multilayered interlayer
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Improving the mechanical and tribological properties of amorphous carbonbased films by an a-C/Zr/ZrN multilayered interlayer Ju Ronga,b, Jiahao Wua, Jinping Tuc, Xiaohua Yua, Jing Fengb, , Zhentao Yuana ⁎
a
Faculty of Materials Science & Engineering, Kunming University of Science and Technology, Kunming 650093, China Institute of Metal Research, Chinese Academy of Sciences, Shenyang National Laboratory for Materials Science, Shenyang 110016, China c State Key Laboratory of Silicon Materials, School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China b
ARTICLE INFO
ABSTRACT
Keywords: Medical implants Multilayered films Tribological performance A-C/Zr/ZrN
Amorphous carbon (a-C) films have been widely investigated to reduce the wear of medical implants due to their excellent tribological performance; however, the high internal stress of a-C films generated during the fabrication process remains an important scientific problem. Herein, we report novel a-C-based films with an a-C/Zr/ ZrN multilayered interlayer. Our results reveal that, with increasing thickness of the multilayered interlayer, the hardness of the films decreased while their toughness and adhesion were improved. The Zr layers could act as a ductile phase, providing a toughening effect. A film with a 2:1 thickness ratio of multilayered interlayer to a-C top layer exhibited favorable tribological properties at various applied loads, especially at high applied load. The results indicated that by introducing a multilayered interlayer into a-C based films, the toughness and adhesion could be significantly improved without adversely sacrificing hardness. The tribological properties could be optimized by carefully tailoring the thickness ratio of multilayered interlayer to a-C top layer.
1. Introduction Medical implants are of great value in that they can maintain the mobility and reduce the pain of osteoarthritis patients [1]. However, the wear debris generated during motion can cause inflammatory responses and aseptic loosening [2,3]. In order to improve the safety performance of medical implants, their mechanical and tribological properties must be further studied. Amorphous carbon (a-C) films have been reported to reduce the wear of medical implants due to their high hardness, while also showing good bio- and hemo-compatibility [4–8]. Unfortunately, high internal stress of a-C films generated during the fabrication process is a major concern [9–16]. Diamond-like carbon (a kind of amorphous carbon) has also been investigated as a protective film on artificial joints, but shows short lifespans due to a delamination problem [17,18]. Researchers have found that artificial hip joints can be subject to an extremely high contact stress of over 1 GPa (edge loading) [1], which is also a huge challenge for a-C films with poor mechanical properties. Recently, much attention in film development has been paid to multifunctionality. Multilayered films can offer enhanced properties over their single-layered counterparts by combining the properties of various materials within one film. For a-C-based films, the multilayered structure should be designed to obtain excellent mechanical properties ⁎
without sacrificing the intrinsic tribological performance of a-C [19–26]. Therefore, the multilayered structure should be carefully and appropriately selected to obtain optimal mechanical and tribological properties. Ceramic-type films, such as those of nitrides and carbides, possess remarkably high hardness, high stiffness, and low fracture resistance. Metal films are much more resistant to fracture, but their hardness and wear resistance are very low [23]. In other words, design of a multilayered structure with a-C, ceramic, and metal layers, combining their advantages, could be an optimal solution. Herein, we report the design and assembly of a novel a-C/Zr/ZrN multilayered structure. Nyberg et al. reported that a carbon coating could improve the running-in performance during the wear process of an implant [27]. Hence an a-C top layer was fabricated as the outermost layer and an a-C/Zr/ZrN multilayer was designed as an interlayer. A series of films with various thickness ratios of multilayered interlayer versus a-C top layer was prepared and investigated with the purpose of optimizing the film architecture. To the best of our knowledge, a multilayered structure acting as an interlayer has rarely been studied [28], and further investigation is needed. 2. Materials and methods The a-C-based films were prepared by means of a closed-field
Corresponding author. E-mail address: [email protected] (J. Feng).
https://doi.org/10.1016/j.ceramint.2018.12.193 Received 19 November 2018; Received in revised form 20 December 2018; Accepted 25 December 2018 0272-8842/ © 2018 Published by Elsevier Ltd.
Please cite this article as: Rong, J., Ceramics International, https://doi.org/10.1016/j.ceramint.2018.12.193
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Fig. 1. Schematic depiction of the deposition process.
unbalanced magnetron sputtering system fitted with graphite targets (purity 99.9%) and zirconium (purity 99.9%). Films deposited on Si wafers (Hs = 12.4 GPa, Es = 198 GPa, ts = 525 ± 20 µm) were employed to analyze microstructure and hardness, and films on Ti6Al4V alloy (Shaanxi Cxmet Technology Co. Ltd.; chemical composition of the alloy (wt%): Al 6, V 4, Fe 0.05, C 0.01, N 0.009, H 0.001, O 0.06, and Ti balance) were used for scratch and tribological tests. In order to maintain homogeneous deposition, the substrates were ultrasonically washed and mounted on a turning holder (10 rpm). The base pressure and temperature in the vacuum chamber were maintained at 3 × 10−3 Pa and 100 °C prior to deposition. Ar as the working gas (purity 99.99%) and nitrogen as the reactive gas (purity 99.99%) were used during the deposition process. Additionally, the substrate was biased by a negative bias pulsed at a frequency of 60 kHz with a duty cycle of 35%. To remove surface contamination from the substrates, all samples were cleaned by sputtering in an Ar plasma (at −500 V for 30 min). The deposition process (as shown in Fig. 1) can be described as follows. i) To enhance the adhesion strength between films and substrates, a Zr buffer layer was firstly deposited on the substrate over a period of 10 min ii) the multilayered interlayer of a-C/Zr/ZrN was sequentially deposited on the buffer layer. During the process, Ar gas was used to deposit the a-C and Zr layers, and a mixed gas of Ar and N2 was used to deposit the ZrN layer. iii) An a-C layer was finally deposited on the interlayer as a top layer. The thicknesses of the layers were determined by the deposition times. The Ar flow was kept constant at 50 sccm, while the N2 flow was controlled by setting the optical emission intensity at 50%. A series of films with different thickness ratios of multilayered interlayer to the top a-C layer was fabricated to assess the effect of different multilayered interlayers on the mechanical and tribological behavior. A film with a top a-C layer of thickness 70 nm was labeled as F1. Films with different thickness ratios of the multilayered interlayer to the top a-C layer were designated as F2 (2:1), F3 (1:1), and F4 (1:2). A film without a multilayered interlayer was named as F5. A scanning electron microscope (SEM, Gemini Sigma 300, Germany) was used to characterize the morphologies of the films at low magnification. The fine microstructure of films was observed by transmission electron microscopy (TEM, FEI Tecnai G2F20), whereby the operating voltage was 200 kV and the point resolution was 0.19 nm. The technique for preparation of thin specimens for TEM observation was based on a focused ion beam (FEI xP200, USA), in which a Ga ion
source (operated at 30 kV and 40 pA) was used to avoid radiation damage. Raman spectra (LABRAM, HR-800), obtained using a 514.5 nm laser as the light source, were used to analyze the atomic bonding of the films. A nano-indentor (Agilent Technologies, G-200) fitted with a Berkovich diamond indenter was employed to test nano-indentation. The hardness and elastic moduli of the films were calculated by the Oliver-Pharr method. The depth of penetration was limited to 100 nm to avoid the influence of the substrate during measurements. Furthermore, to ensure reliability of the data, more than six indentations at various locations were made in each sample. The toughness of the films was evaluated by the Vicker's indentation method, applying normal loads of 50, 100, and 200 g. Scratch tests were performed to evaluate the cohesive and adhesion strengths between the films and the Ti6Al4V alloy substrate. A Rockwell diamond pin (radius 0.2 mm) served as the scratch indenter, which was drawn (at a constant speed of 1 mm min-1) across the surface of the films, as the load was linearly increased from 0 to 80 N. At least six such scratch tests were carried out on each sample. Contact angles (CAs) were measured with a contact angle meter (NS810, 3nh Tech., Shenzhen) by the sessile drop method (3 μL drops of deionized water). Furthermore, to ensure clean surfaces, CA tests were performed within 10 min of sample preparation. CAs at six different locations were measured on each sample to obtain reliable data. A ball-on-disk tribometer was employed to evaluate the tribological behavior. Tests were performed in ambient air (room temperature) and fetal bovine serum (FBS) at a sliding velocity of 0.1 m s−1. A 4 mm diameter Si3N4 ceramic ball (hardness 1500 HV) served as the counterpart in these tests. The tribological tests were conducted at different applied loads (1, 5, and 10 N) for a duration of 10 min. The wear rates of the films, defined as the volume loss per sliding distance, were calculated from the test wear tracks. 3. Results and discussion SEM morphologies of the films are displayed in Fig. 2. All films consisted of a Zr buffer layer and an interlayer with a multilayered structure. The thickness of the Zr buffer layer was measured as approximately 200 nm, implying a deposition rate of 20 nm min−1. The thicknesses of the multilayered interlayer and the a-C top layer were controlled by their deposition times. As can be seen, the total thickness of each film was tailored to be about 2 µm. The deposition rates of the a2
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Intensity / a.u.
Fig. 2. SEM images of the films: a) F1; b) F2; c) F3; d) F4; e) F5; the insets are surface SEM images.
F1 F3 F5 Fitting curves D peak G peak
D peak
G peak
F1 F3 F5
1000
1200
1400
1600
1800
-1
Raman shift cm
Fig. 4. Peak-fitting of the Raman spectra of F1, F3, and F5. Fig. 3. TEM and HRTEM images of F3. 3
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Table 1 Data from Raman spectra of the films. Sample
D peak position (cm−1)
FWHM of D peak (cm−1)
G peak position (cm−1)
FWHM of G peak (cm−1)
ID/IG
F1 F3 F5
1371.7 1370.2 1368.3
324.6 319.9 323.3
1552.7 1548.7 1547.3
137.7 142.3 147.9
1.42 1.06 1.01
Fig. 5. SEM images of Vicker's indentation of the films tested at 50 g (a–e), 100 g (f–j), and 200 g (k–o).
Table 2 Mechanical properties of the films. Sample
H (GPa)
E (GPa)
H/E
H3/E2 (GPa)
F1 F2 F3 F4 F5
18.1 18.6 18.8 19.3 19.7
196.1 208.4 214.5 226.5 235.2
0.092 0.089 0.088 0.085 0.084
0.154 0.148 0.144 0.140 0.138
the D peak and the G peak [29–32]. Table 1 presents information gleaned from the fitted Raman spectra. It can be seen that the position and the full-width at half-maximum (FWHM) of the D peaks for F1, F3, and F5 show little difference. However, the value of ID/IG for F1 is 1.42, significantly higher than those for F3 (1.06) and F5 (1.01). As previously reported in the literature, the introduction of Ti can lead to a decrease in the sp3 fraction [33]. For F1, the sp3 fraction in the a-C top layer (70 nm thickness) is affected by the adjacent Zr monolayer. However, for F3 and F5, the thicknesses of the a-C top layer range from several hundreds to thousands of nanometers, making the effect of the adjacent Zr layer insignificant. It was also found that the Zr monolayers in the multilayered interlayer could affect the adjacent a-C monolayers and lead to a decrease in their sp3 fraction. Fig. 5 presents Vicker's indentations of the films, obtained at normal loads of 50 g (a–e), 100 g (f–j), and 200 g (k–o) [34,35]. At a normal load of 50 g, radial cracks formed on the films were very small. At 100 g, although radial cracks were still barely discernible, cracks extending from the edge of the indentation were visible for F2, F3, F4, and F5. Further increasing the normal load to 200 g led to different results.
Fig. 6. Cross-section TEM image of Vicker's indentation of F1 tested at 50 g.
C top layer and interlayer were 12 and 11 nm min−1, respectively. In addition, the surfaces of each of the films showed uniform clusters of size 50–100 nm. In order to study the layer morphology in detail, further microstructure characterization was carried out by TEM (Fig. 3). For the multilayered structure, the thicknesses of a-C (amorphous structure), the Zr buffer layer, and the ZrN monolayer (nanocrystalline structure) were 70, 15, and 35 nm, respectively. The ZrN monolayer was located between two Zr monolayers, forming a “sandwich” structure, and the corresponding thicknesses are shown in Fig. 3. The Raman spectra of F1, F3, and F5 are shown in Fig. 4, which feature typical wide peaks of a-C between 1000 and 1800 cm−1. As indicated in Fig. 4, the wide peaks can be fitted by two Gaussian bands: 4
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Fig. 7. Optical images of the mid-regions (a–e) and terminal regions (f–j) of scratch tracks in the films.
Fig. 8. Optical images of water contact angle determinations of the films.
For F1, the radial cracks were still very small, but for F5 they were as long as several micrometers. Additionally, it is worth noting that with decreasing thickness of the multilayered interlayer, the length of the radial cracks increased, indicating that introduction of the multilayered interlayer can improve the fracture toughness of the films. Fig. 6 presents a cross-section TEM image of the indentation of F1 obtained at a normal load of 50 g. It can be seen that cracks mainly appeared in the aC layer and did not propagate across the interface of the a-C layer and Zr layer. It has been reported in previous papers that the interfaces in a multilayer can hinder the propagation of cracks [19,20,36]. In this work, the Zr layers acted as a ductile phase and imparted a toughening effect. Values of hardness (H), elastic modulus (E), H/E, and H3/E2 for the films are listed in Table 2. F5 without a multilayered interlayer exhibited the highest hardness of about 19.5 GPa. There is an inverse relationship between the thickness of the multilayered interlayer and hardness. Nevertheless, F1 retained a relatively high hardness of 18.2 GPa. With increasing thickness of the multilayered interlayers, the hardness and elastic moduli of the films decreased. The decrease in hardness can be mainly attributed to the presence of soft Zr layers in the multilayered interlayer. The sp3 fraction may also have affected the hardness of the films. F1 showed the highest H/E ratio of 0.093 and an H3/E2 ratio of 0.157. The H3/E2 ratios of the films decreased with increasing thickness of the a-C top layer, in accordance with the results of the Vicker's indentation tests [37].
In order to investigate the adhesion of the films, scratch tests were performed. Typical images of scratch tracks are presented in Fig. 7, corresponding to the mid-regions (a–e) and the terminal regions (f–j). For F1, almost no adhesive failure could be discerned in the mid-region of the scratch track. However, adhesion failure could be observed alongside the scratch track for F2. The adhesion failure became increasingly pronounced with decreasing thickness of the multilayered interlayer. Around the terminal regions of the scratch tracks, few cracks were seen for F1, but they became increasingly prevalent with decreasing thickness of the multilayered interlayer. The films simultaneously bore compression and friction forces during the scratch tests, which could lead to crack generation and propagation. Gioti et al. suggest that intrinsic stress in a thick a-C film can be dissipated by a multilayer structure, whereby layers with low compressive stress can act as buffers and relieve stress [38]. In this work, the Zr layers may serve as buffers and dissipate stress. Consequently, adhesion failure increased (from F2 to F5) with decreasing thickness of the interlayer. As discussed above, F1 with a multilayered structure exhibited the highest toughness as well as the best crack propagation resistance, and thus showed the best adhesion. The results imply that the multilayered interlayer improves the adhesion of the film. Water CAs are presented in Fig. 8. F1 showed a CA of about 76°, while F5 exhibited a slightly lower value of about 68°. The water CAs of the other films were between 68° and 76°, thus showing no significant variation. Previous studies have shown that hydrophobicity can 5
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0.20
10 0.10 5
Coefficient of friction
F1
F2
F3
F4
F5
0.20
0
30 Coefficient of fiction Wear rate
5N
0.15
24 18 12
0.10
6 0.05
c)
15
0.15
F1
F2
F3
F4
F5
0.20
0 50
Coefficient of fiction Wear rate
10 N
0.15
40 30
Broken 20
0.10
10 0.05
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
b)
1N
Wear rate / 10-17 m3N-1m-1
Coefficient of friction Wear rate
0.05
slightly higher. Comparing F4 with F5, it can be inferred that introduction of the multilayered interlayer significantly improved the tribological properties of the film. With increasing applied load, F2 showed the lowest wear rate. For a-C-based films, a graphitized tribolayer that readily formed on the counterpart was used to provide a good lubricant effect [40–42]. In addition, a graphite structure may have been generated on the wear track surface. Fig. 10 presents optical images of the wear tracks of films tested at applied loads of 1 N (a–e), 5 N (f–j), and 10 N (k–o) in ambient air. For sample F5, adhesive failures occurred at 1 N and 5 N, and it was fully worn at 10 N. F4 showed relatively smooth wear tracks at 1 N and 5 N, but pronounced grooves could be seen in the wear track at 10 N. The wear tracks in F2 and F3 indicate that the wear resistance improved with increasing thickness of the multilayered interlayer. It is worth noting that even at an applied load of 10 N, the wear track of F2 was still very narrow and smooth. This can be attributed to high H/E ratio, relatively high toughness, and good adhesion. F1 showed the highest H/E ratio and fracture toughness among the as-deposited films, but grooves were generated in the wear track at 10 N. As is evident from the cross-section image, the thickness of the top a-C layer of F1 was 70 nm (the same as the a-C monolayer in the multilayered interlayer). If the top a-C layer is removed when the film is subjected to a large load during the wear process, the Zr layers and ZrN layers will become exposed, thus leading to friction. Hard wear debris may be generated in this process, resulting in the formation of grooves in the wear track. Fig. 11 presents the friction coefficients and wear rates tested in FBS. The friction coefficients of all films were low at 1 N, and their wear rates were also low under these conditions. According to Ghosha et al., for artificial hip joints, a significant amount of wear debris is generated by edge loading, which can cause aseptic loosening [1]. Elkins et al. suggest that significant wear occurs under edge load condition (Hertzian contact load > 1 GPa [43]). Therefore, tests were also performed at 10 N, corresponding to an initial Hertzian stress of about 1 GPa. F4 and F5 were completely broken at 10 N, whereas F1, F2, and F3 exhibited friction coefficients below 0.1 and low wear rates below 1.3 × 10−16 m3 N−1 m-1. In short, the a-C/Zr/ZrN multilayer showed low friction coefficient and wear rates at both low applied load of 1 N and high applied load of 10 N in FBS. Fig. 12 presents the wear tracks of the films tested at 1 N (a–e) and 10 N (f–j) in FBS. The wear tracks of all of the films were smooth and narrow at 1 N, but F4 and F5 exhibited wide and rough wear tracks at 10 N. F2 showed the narrowest wear track. FBS solution supplied a protein-rich environment. Thus, when tribo-tests were carried out in this medium, the adsorbed protein layer reduced direct contact between the counterpart and the film, improving the boundary lubrication. At an applied load of 1 N, F5 was partially broken when tested in ambient air, but it exhibited a very low wear rate of 5.6 × 10−17 m3 N−1 m−1 in FBS, indicating a very effective lubrication effect of the adsorbed protein layer. Under an edge load, the thickness of the multilayered interlayer becomes crucially important. For F4 and F5, the crack propagation resistance is insufficient, and cracks generate and propagate under the applied load, leading to breakage of the films. The H/E and H3/E2 ratios of F4 and F5 were lower than those for the other films, in accordance with the results of tribological tests. It is interesting that F2 exhibited the lowest wear rates of 2.7 × 10−17 m3 N—1 m−1 at 1 N and 7.5 × 10−17 m3 N−1 m−1 at 10 N. This indicates that in order to design an a-C-based film with high tribological performance, one effective solution is to introduce a multilayered interlayer below an a-C top layer with appropriate thickness ratio.
20
F1
F2
F3
F4
F5
0
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
a)
Fig. 9. Friction coefficients and wear rates of the films in ambient air at applied loads of a) 1 N; b) 5 N; c) 10 N.
effectively improve protein adsorption and wear properties [39]. In this work, as the CAs of the films did not vary significantly, we surmise that the influence of an adsorbed protein on the bio-tribological properties would be small. With regard to the tribological behavior of the films, their friction coefficients and the corresponding specific wear rates are shown in Fig. 9. All films exhibited very low friction coefficients of 0.09–0.15 under loads of 1 N and 5 N, although the friction coefficient of F5 was slightly higher than those of the other films. At a high applied load of 10 N, F5 was completely broken and showed a friction coefficient of about 0.5, attributable to the Ti6Al4V substrate, whereas the other films still exhibited low friction coefficients. F5 was partially broken even at a low applied load of 1 N, and the wear rate was impractically large compared to those of the other films. Therefore, the wear rate of F5 is not shown in Fig. 9. F1, F2, and F3 showed similar wear rates at a low applied load of 1 N, whereas the wear rate of F4 was
4. Conclusions This work has revealed that the mechanical and tribological properties of a-C-based films can be improved by a-C/Zr/ZrN multilayers. Additionally, the thickness ratio of the multilayered interlayer to the aC top layer is closely related with the mechanical and tribological 6
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Fig. 10. Optical images of the wear tracks of the films tested at applied loads of 1 N (a–e), 5 N (f–j) and 10 N (k–o) in ambient air.
0.20
Coefficient of friction
1N
0.15
6
0.10
3
F1
F2
F3
F4
F5
0
20
0.20 Coefficient of friction Wear rate
10 N 15
0.15 10
Broken 0.10
5
0.05
Wear rate / 10-17 m3N-1m-1
Coefficient of friction Wear rate
0.05
b)
9
F1
F2
F3
F4
F5
0
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
a)
Fig. 11. Friction coefficients and wear rates of the films in FBS at applied loads of a) 1 N; b) 10 N.
properties of the films. With increasing thickness of the multilayered interlayer, toughness and adhesion are increased. A film with a 2:1 thickness ratio of multilayered interlayer to a-C top layer exhibited excellent tribological properties in both ambient air and FBS under various applied loads, especially at high applied loads. In conclusion,
we have shown that by introducing a multilayered interlayer into a-Cbased films, the toughness and adhesion can be significantly improved without adversely sacrificing hardness. The tribological properties may be optimized by carefully tailoring the thickness ratio of the multilayered interlayer to the a-C top layer. 7
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Fig. 12. Optical images of the wear tracks of the films tested at 1 N (a–e) and 10 N (f–j) in FBS.
Acknowledgements The authors gratefully acknowledge financial support from the National Natural Science Foundation of China (Grant Nos. 51665022, 51801086 and 51601081).
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of the size-dependent cytotoxicity of DLC (Diamondlike Carbon) wear debris in arthroplasty applications, ACS Biomater. Sci. Eng. 3 (2017) 530–539. C.V. Falub, U. Müller, G. Thorwarth, M. Parlinska-Wojtan, C. Voisard, R. Hauert, In vitro studies of the adhesion of diamond-like carbon thin films on CoCrMo biomedical implant alloy, Acta Mater. 59 (2011) 4678–4689. L. Li, L.L. Liu, X. Li, P. Guo, P. Ke, A. Wang, Enhanced tribocorrosion performance of Cr/GLC multilayered films for marine protective application, ACS Appl. Mater. Interfaces 10 (2018) 13187–13198. L. Major, R. Major, M. Kot, J.M. Lackner, B. Major, Ex situ and in situ nanoscale wear mechanisms characterization of Zr/ZrxN tribological coatings, Wear 404 (2018) 82–91. K.M. Jiang, D.Q. Zhao, X. Jiang, Q. Huang, L.J. Miao, H.M. Lu, Y. Li, Electronicstructure, corrosion and mechanical properties of nc-CrC/aC: H films deposited by multi-arc ion plating, J. Alloy. Compd. 750 (2018) 560–569. W. Dai, X. Gao, J. Liu, S.H. Kwon, Q. Wang, Compositionally modulated multilayer diamond-like carbon coatings with AlTiSi multi-doping by reactive high power impulse magnetron sputtering, Appl. Surf. Sci. 425 (2017) 855–861. S. Zhou, J. Lei, Z. Xiong, X. Dai, J. Guo, Z. Gu, Synthesis of Fep/Cu-Cup/Fe duplex composite coatings by laser cladding, Mater. Des. 97 (2016) 431–436. C.Q. Guo, Z.L. Pei, D. Fan, R.D. Liu, J. Gong, C. Sun, Predicting multilayer film's residual stress from its monolayers, Mater. Des. 110 (2016) 858–864. W.Q. Bai, L.L. Li, R.L. Li, C.D. Gu, X.L. Wang, G. Jin, D.G. Liu, J.P. Tu, Deposition and characterization of a ZrN/Zr/aC multilayer: implication on bio-tribological and corrosion behaviors, Surf. Coat. Technol. 324 (2017) 509–517. A. Kelling, K.R. Mangipudi, I. Knorr, T. Liese, H.U. Krebs, C.A. Volkert, Investigating fracture of nanoscale metal–ceramic multilayers in the transmission electron microscope, Scr. Mater. 115 (2016) 42–45. H. Nyberg, T. Tokoroyama, U. Wiklund, S. Jacobson, Design of low-friction PVD coating systems with enhanced running-in performance—carbon overcoats on TaC/ aC coatings, Surf. Coat. Technol. 222 (2013) 48–54. M.A. Domínguez-Crespo, A.M. Torres-Huerta, E. Rodríguez, A. GonzálezHernández, S.B. Brachetti-Sibaja, H.J. Dorantes-Rosales, A.B. López-Oyama, Effect of deposition parameters on structural, mechanical and electrochemical properties in Ti/TiN thin films on AISI 316L substrates produced by r. f. magnetron sputtering, J. Alloy. Compd. 746 (2018) 688–698. X. Sui, R. Xu, J. Liu, S. Zhang, Y. Wu, J. Yang, J. Hao, Tailoring the Tribocorrosion and antifouling performance of (Cr, Cu)-GLC coatings for marine application, ACS Appl. Mater. Interfaces 10 (2018) 36531–36539. C. Casiraghi, A.C. Ferrari, J. Robertson, Raman spectroscopy of hydrogenated amorphous carbons, Phys. Rev. B 72 (2005) 085401. A.C. Ferrari, J. Robertson, Interpretation of Raman spectra of disordered and amorphous carbon, Phys. Rev. B 61 (2000) 14095–14107. J.R. Webb, A.A. Martin, R.P. Johnson, M.B. Joseph, M.E. Newton, Fabrication of a single sub-micron pore spanning a single crystal (100) diamond membrane and impact on particle translocation, Carbon 122 (2017) 319–328. M. Braic, A. Vladescu, M. Balaceanu, C. Luculescu, S.C. Padmanabhan, L. Constantin, M.A. Morris, V. Braic, C.E.A. Grigorescu, P. Ionescu, M.D. Dracea, C. Logofatu, A comparative study of the structural, mechanical and tribological characteristics of TiSiC-Cr coatings prepared in CH4 and C2H2 reactive atmosphere by cathodic vacuum arc, Appl. Surf. Sci. 400 (2017) 318–328. Y. Qin, J. He, F. Yin, B. Liu, F. Zhang, Effect of Ti particle size on mechanical and tribological properties of TiCN coatings prepared by reactive plasma spraying, Ceram. Int. 43 (2017) 16548–16554. Z. Zhao, P. Hui, T. Wang, X. Wang, Y. Xu, L. Zhong, M. Zhao, New strategy to grow TiC coatings on titanium alloy: contact solid carburization by cast iron, J. Alloy. Compd. 745 (2018) 637–643. Y.X. Ou, H. Chen, Z.Y. Li, J. Lin, W. Pan, M.K. Lei, Microstructure and tribological behavior of TiAlSiN coatings deposited by deep oscillation magnetron sputtering, J. Am. Ceram. Soc. 101 (2018) 5166–5176. F. Cai, M. Chen, M. Li, S. Zhang, Influence of negative bias voltage on microstructure and property of Al-Ti-N films deposited by multi-arc ion plating, Ceram. Int. 43 (2017) 3774–3783. M. Gioti, S. Logothetidis, C. Charitidis, Stress relaxation and stability in thick
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Improving the mechanical and tribological properties of amorphous carbonbased films by an a-C/Zr/ZrN multilayered interlayer Ju Ronga,b, Jiahao Wua, Jinping Tuc, Xiaohua Yua, Jing Fengb, , Zhentao Yuana ⁎
a
Faculty of Materials Science & Engineering, Kunming University of Science and Technology, Kunming 650093, China Institute of Metal Research, Chinese Academy of Sciences, Shenyang National Laboratory for Materials Science, Shenyang 110016, China c State Key Laboratory of Silicon Materials, School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China b
ARTICLE INFO
ABSTRACT
Keywords: Medical implants Multilayered films Tribological performance A-C/Zr/ZrN
Amorphous carbon (a-C) films have been widely investigated to reduce the wear of medical implants due to their excellent tribological performance; however, the high internal stress of a-C films generated during the fabrication process remains an important scientific problem. Herein, we report novel a-C-based films with an a-C/Zr/ ZrN multilayered interlayer. Our results reveal that, with increasing thickness of the multilayered interlayer, the hardness of the films decreased while their toughness and adhesion were improved. The Zr layers could act as a ductile phase, providing a toughening effect. A film with a 2:1 thickness ratio of multilayered interlayer to a-C top layer exhibited favorable tribological properties at various applied loads, especially at high applied load. The results indicated that by introducing a multilayered interlayer into a-C based films, the toughness and adhesion could be significantly improved without adversely sacrificing hardness. The tribological properties could be optimized by carefully tailoring the thickness ratio of multilayered interlayer to a-C top layer.
1. Introduction Medical implants are of great value in that they can maintain the mobility and reduce the pain of osteoarthritis patients [1]. However, the wear debris generated during motion can cause inflammatory responses and aseptic loosening [2,3]. In order to improve the safety performance of medical implants, their mechanical and tribological properties must be further studied. Amorphous carbon (a-C) films have been reported to reduce the wear of medical implants due to their high hardness, while also showing good bio- and hemo-compatibility [4–8]. Unfortunately, high internal stress of a-C films generated during the fabrication process is a major concern [9–16]. Diamond-like carbon (a kind of amorphous carbon) has also been investigated as a protective film on artificial joints, but shows short lifespans due to a delamination problem [17,18]. Researchers have found that artificial hip joints can be subject to an extremely high contact stress of over 1 GPa (edge loading) [1], which is also a huge challenge for a-C films with poor mechanical properties. Recently, much attention in film development has been paid to multifunctionality. Multilayered films can offer enhanced properties over their single-layered counterparts by combining the properties of various materials within one film. For a-C-based films, the multilayered structure should be designed to obtain excellent mechanical properties ⁎
without sacrificing the intrinsic tribological performance of a-C [19–26]. Therefore, the multilayered structure should be carefully and appropriately selected to obtain optimal mechanical and tribological properties. Ceramic-type films, such as those of nitrides and carbides, possess remarkably high hardness, high stiffness, and low fracture resistance. Metal films are much more resistant to fracture, but their hardness and wear resistance are very low [23]. In other words, design of a multilayered structure with a-C, ceramic, and metal layers, combining their advantages, could be an optimal solution. Herein, we report the design and assembly of a novel a-C/Zr/ZrN multilayered structure. Nyberg et al. reported that a carbon coating could improve the running-in performance during the wear process of an implant [27]. Hence an a-C top layer was fabricated as the outermost layer and an a-C/Zr/ZrN multilayer was designed as an interlayer. A series of films with various thickness ratios of multilayered interlayer versus a-C top layer was prepared and investigated with the purpose of optimizing the film architecture. To the best of our knowledge, a multilayered structure acting as an interlayer has rarely been studied [28], and further investigation is needed. 2. Materials and methods The a-C-based films were prepared by means of a closed-field
Corresponding author. E-mail address: [email protected] (J. Feng).
https://doi.org/10.1016/j.ceramint.2018.12.193 Received 19 November 2018; Received in revised form 20 December 2018; Accepted 25 December 2018 0272-8842/ © 2018 Published by Elsevier Ltd.
Please cite this article as: Rong, J., Ceramics International, https://doi.org/10.1016/j.ceramint.2018.12.193
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Fig. 1. Schematic depiction of the deposition process.
unbalanced magnetron sputtering system fitted with graphite targets (purity 99.9%) and zirconium (purity 99.9%). Films deposited on Si wafers (Hs = 12.4 GPa, Es = 198 GPa, ts = 525 ± 20 µm) were employed to analyze microstructure and hardness, and films on Ti6Al4V alloy (Shaanxi Cxmet Technology Co. Ltd.; chemical composition of the alloy (wt%): Al 6, V 4, Fe 0.05, C 0.01, N 0.009, H 0.001, O 0.06, and Ti balance) were used for scratch and tribological tests. In order to maintain homogeneous deposition, the substrates were ultrasonically washed and mounted on a turning holder (10 rpm). The base pressure and temperature in the vacuum chamber were maintained at 3 × 10−3 Pa and 100 °C prior to deposition. Ar as the working gas (purity 99.99%) and nitrogen as the reactive gas (purity 99.99%) were used during the deposition process. Additionally, the substrate was biased by a negative bias pulsed at a frequency of 60 kHz with a duty cycle of 35%. To remove surface contamination from the substrates, all samples were cleaned by sputtering in an Ar plasma (at −500 V for 30 min). The deposition process (as shown in Fig. 1) can be described as follows. i) To enhance the adhesion strength between films and substrates, a Zr buffer layer was firstly deposited on the substrate over a period of 10 min ii) the multilayered interlayer of a-C/Zr/ZrN was sequentially deposited on the buffer layer. During the process, Ar gas was used to deposit the a-C and Zr layers, and a mixed gas of Ar and N2 was used to deposit the ZrN layer. iii) An a-C layer was finally deposited on the interlayer as a top layer. The thicknesses of the layers were determined by the deposition times. The Ar flow was kept constant at 50 sccm, while the N2 flow was controlled by setting the optical emission intensity at 50%. A series of films with different thickness ratios of multilayered interlayer to the top a-C layer was fabricated to assess the effect of different multilayered interlayers on the mechanical and tribological behavior. A film with a top a-C layer of thickness 70 nm was labeled as F1. Films with different thickness ratios of the multilayered interlayer to the top a-C layer were designated as F2 (2:1), F3 (1:1), and F4 (1:2). A film without a multilayered interlayer was named as F5. A scanning electron microscope (SEM, Gemini Sigma 300, Germany) was used to characterize the morphologies of the films at low magnification. The fine microstructure of films was observed by transmission electron microscopy (TEM, FEI Tecnai G2F20), whereby the operating voltage was 200 kV and the point resolution was 0.19 nm. The technique for preparation of thin specimens for TEM observation was based on a focused ion beam (FEI xP200, USA), in which a Ga ion
source (operated at 30 kV and 40 pA) was used to avoid radiation damage. Raman spectra (LABRAM, HR-800), obtained using a 514.5 nm laser as the light source, were used to analyze the atomic bonding of the films. A nano-indentor (Agilent Technologies, G-200) fitted with a Berkovich diamond indenter was employed to test nano-indentation. The hardness and elastic moduli of the films were calculated by the Oliver-Pharr method. The depth of penetration was limited to 100 nm to avoid the influence of the substrate during measurements. Furthermore, to ensure reliability of the data, more than six indentations at various locations were made in each sample. The toughness of the films was evaluated by the Vicker's indentation method, applying normal loads of 50, 100, and 200 g. Scratch tests were performed to evaluate the cohesive and adhesion strengths between the films and the Ti6Al4V alloy substrate. A Rockwell diamond pin (radius 0.2 mm) served as the scratch indenter, which was drawn (at a constant speed of 1 mm min-1) across the surface of the films, as the load was linearly increased from 0 to 80 N. At least six such scratch tests were carried out on each sample. Contact angles (CAs) were measured with a contact angle meter (NS810, 3nh Tech., Shenzhen) by the sessile drop method (3 μL drops of deionized water). Furthermore, to ensure clean surfaces, CA tests were performed within 10 min of sample preparation. CAs at six different locations were measured on each sample to obtain reliable data. A ball-on-disk tribometer was employed to evaluate the tribological behavior. Tests were performed in ambient air (room temperature) and fetal bovine serum (FBS) at a sliding velocity of 0.1 m s−1. A 4 mm diameter Si3N4 ceramic ball (hardness 1500 HV) served as the counterpart in these tests. The tribological tests were conducted at different applied loads (1, 5, and 10 N) for a duration of 10 min. The wear rates of the films, defined as the volume loss per sliding distance, were calculated from the test wear tracks. 3. Results and discussion SEM morphologies of the films are displayed in Fig. 2. All films consisted of a Zr buffer layer and an interlayer with a multilayered structure. The thickness of the Zr buffer layer was measured as approximately 200 nm, implying a deposition rate of 20 nm min−1. The thicknesses of the multilayered interlayer and the a-C top layer were controlled by their deposition times. As can be seen, the total thickness of each film was tailored to be about 2 µm. The deposition rates of the a2
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Intensity / a.u.
Fig. 2. SEM images of the films: a) F1; b) F2; c) F3; d) F4; e) F5; the insets are surface SEM images.
F1 F3 F5 Fitting curves D peak G peak
D peak
G peak
F1 F3 F5
1000
1200
1400
1600
1800
-1
Raman shift cm
Fig. 4. Peak-fitting of the Raman spectra of F1, F3, and F5. Fig. 3. TEM and HRTEM images of F3. 3
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Table 1 Data from Raman spectra of the films. Sample
D peak position (cm−1)
FWHM of D peak (cm−1)
G peak position (cm−1)
FWHM of G peak (cm−1)
ID/IG
F1 F3 F5
1371.7 1370.2 1368.3
324.6 319.9 323.3
1552.7 1548.7 1547.3
137.7 142.3 147.9
1.42 1.06 1.01
Fig. 5. SEM images of Vicker's indentation of the films tested at 50 g (a–e), 100 g (f–j), and 200 g (k–o).
Table 2 Mechanical properties of the films. Sample
H (GPa)
E (GPa)
H/E
H3/E2 (GPa)
F1 F2 F3 F4 F5
18.1 18.6 18.8 19.3 19.7
196.1 208.4 214.5 226.5 235.2
0.092 0.089 0.088 0.085 0.084
0.154 0.148 0.144 0.140 0.138
the D peak and the G peak [29–32]. Table 1 presents information gleaned from the fitted Raman spectra. It can be seen that the position and the full-width at half-maximum (FWHM) of the D peaks for F1, F3, and F5 show little difference. However, the value of ID/IG for F1 is 1.42, significantly higher than those for F3 (1.06) and F5 (1.01). As previously reported in the literature, the introduction of Ti can lead to a decrease in the sp3 fraction [33]. For F1, the sp3 fraction in the a-C top layer (70 nm thickness) is affected by the adjacent Zr monolayer. However, for F3 and F5, the thicknesses of the a-C top layer range from several hundreds to thousands of nanometers, making the effect of the adjacent Zr layer insignificant. It was also found that the Zr monolayers in the multilayered interlayer could affect the adjacent a-C monolayers and lead to a decrease in their sp3 fraction. Fig. 5 presents Vicker's indentations of the films, obtained at normal loads of 50 g (a–e), 100 g (f–j), and 200 g (k–o) [34,35]. At a normal load of 50 g, radial cracks formed on the films were very small. At 100 g, although radial cracks were still barely discernible, cracks extending from the edge of the indentation were visible for F2, F3, F4, and F5. Further increasing the normal load to 200 g led to different results.
Fig. 6. Cross-section TEM image of Vicker's indentation of F1 tested at 50 g.
C top layer and interlayer were 12 and 11 nm min−1, respectively. In addition, the surfaces of each of the films showed uniform clusters of size 50–100 nm. In order to study the layer morphology in detail, further microstructure characterization was carried out by TEM (Fig. 3). For the multilayered structure, the thicknesses of a-C (amorphous structure), the Zr buffer layer, and the ZrN monolayer (nanocrystalline structure) were 70, 15, and 35 nm, respectively. The ZrN monolayer was located between two Zr monolayers, forming a “sandwich” structure, and the corresponding thicknesses are shown in Fig. 3. The Raman spectra of F1, F3, and F5 are shown in Fig. 4, which feature typical wide peaks of a-C between 1000 and 1800 cm−1. As indicated in Fig. 4, the wide peaks can be fitted by two Gaussian bands: 4
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Fig. 7. Optical images of the mid-regions (a–e) and terminal regions (f–j) of scratch tracks in the films.
Fig. 8. Optical images of water contact angle determinations of the films.
For F1, the radial cracks were still very small, but for F5 they were as long as several micrometers. Additionally, it is worth noting that with decreasing thickness of the multilayered interlayer, the length of the radial cracks increased, indicating that introduction of the multilayered interlayer can improve the fracture toughness of the films. Fig. 6 presents a cross-section TEM image of the indentation of F1 obtained at a normal load of 50 g. It can be seen that cracks mainly appeared in the aC layer and did not propagate across the interface of the a-C layer and Zr layer. It has been reported in previous papers that the interfaces in a multilayer can hinder the propagation of cracks [19,20,36]. In this work, the Zr layers acted as a ductile phase and imparted a toughening effect. Values of hardness (H), elastic modulus (E), H/E, and H3/E2 for the films are listed in Table 2. F5 without a multilayered interlayer exhibited the highest hardness of about 19.5 GPa. There is an inverse relationship between the thickness of the multilayered interlayer and hardness. Nevertheless, F1 retained a relatively high hardness of 18.2 GPa. With increasing thickness of the multilayered interlayers, the hardness and elastic moduli of the films decreased. The decrease in hardness can be mainly attributed to the presence of soft Zr layers in the multilayered interlayer. The sp3 fraction may also have affected the hardness of the films. F1 showed the highest H/E ratio of 0.093 and an H3/E2 ratio of 0.157. The H3/E2 ratios of the films decreased with increasing thickness of the a-C top layer, in accordance with the results of the Vicker's indentation tests [37].
In order to investigate the adhesion of the films, scratch tests were performed. Typical images of scratch tracks are presented in Fig. 7, corresponding to the mid-regions (a–e) and the terminal regions (f–j). For F1, almost no adhesive failure could be discerned in the mid-region of the scratch track. However, adhesion failure could be observed alongside the scratch track for F2. The adhesion failure became increasingly pronounced with decreasing thickness of the multilayered interlayer. Around the terminal regions of the scratch tracks, few cracks were seen for F1, but they became increasingly prevalent with decreasing thickness of the multilayered interlayer. The films simultaneously bore compression and friction forces during the scratch tests, which could lead to crack generation and propagation. Gioti et al. suggest that intrinsic stress in a thick a-C film can be dissipated by a multilayer structure, whereby layers with low compressive stress can act as buffers and relieve stress [38]. In this work, the Zr layers may serve as buffers and dissipate stress. Consequently, adhesion failure increased (from F2 to F5) with decreasing thickness of the interlayer. As discussed above, F1 with a multilayered structure exhibited the highest toughness as well as the best crack propagation resistance, and thus showed the best adhesion. The results imply that the multilayered interlayer improves the adhesion of the film. Water CAs are presented in Fig. 8. F1 showed a CA of about 76°, while F5 exhibited a slightly lower value of about 68°. The water CAs of the other films were between 68° and 76°, thus showing no significant variation. Previous studies have shown that hydrophobicity can 5
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0.20
10 0.10 5
Coefficient of friction
F1
F2
F3
F4
F5
0.20
0
30 Coefficient of fiction Wear rate
5N
0.15
24 18 12
0.10
6 0.05
c)
15
0.15
F1
F2
F3
F4
F5
0.20
0 50
Coefficient of fiction Wear rate
10 N
0.15
40 30
Broken 20
0.10
10 0.05
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
b)
1N
Wear rate / 10-17 m3N-1m-1
Coefficient of friction Wear rate
0.05
slightly higher. Comparing F4 with F5, it can be inferred that introduction of the multilayered interlayer significantly improved the tribological properties of the film. With increasing applied load, F2 showed the lowest wear rate. For a-C-based films, a graphitized tribolayer that readily formed on the counterpart was used to provide a good lubricant effect [40–42]. In addition, a graphite structure may have been generated on the wear track surface. Fig. 10 presents optical images of the wear tracks of films tested at applied loads of 1 N (a–e), 5 N (f–j), and 10 N (k–o) in ambient air. For sample F5, adhesive failures occurred at 1 N and 5 N, and it was fully worn at 10 N. F4 showed relatively smooth wear tracks at 1 N and 5 N, but pronounced grooves could be seen in the wear track at 10 N. The wear tracks in F2 and F3 indicate that the wear resistance improved with increasing thickness of the multilayered interlayer. It is worth noting that even at an applied load of 10 N, the wear track of F2 was still very narrow and smooth. This can be attributed to high H/E ratio, relatively high toughness, and good adhesion. F1 showed the highest H/E ratio and fracture toughness among the as-deposited films, but grooves were generated in the wear track at 10 N. As is evident from the cross-section image, the thickness of the top a-C layer of F1 was 70 nm (the same as the a-C monolayer in the multilayered interlayer). If the top a-C layer is removed when the film is subjected to a large load during the wear process, the Zr layers and ZrN layers will become exposed, thus leading to friction. Hard wear debris may be generated in this process, resulting in the formation of grooves in the wear track. Fig. 11 presents the friction coefficients and wear rates tested in FBS. The friction coefficients of all films were low at 1 N, and their wear rates were also low under these conditions. According to Ghosha et al., for artificial hip joints, a significant amount of wear debris is generated by edge loading, which can cause aseptic loosening [1]. Elkins et al. suggest that significant wear occurs under edge load condition (Hertzian contact load > 1 GPa [43]). Therefore, tests were also performed at 10 N, corresponding to an initial Hertzian stress of about 1 GPa. F4 and F5 were completely broken at 10 N, whereas F1, F2, and F3 exhibited friction coefficients below 0.1 and low wear rates below 1.3 × 10−16 m3 N−1 m-1. In short, the a-C/Zr/ZrN multilayer showed low friction coefficient and wear rates at both low applied load of 1 N and high applied load of 10 N in FBS. Fig. 12 presents the wear tracks of the films tested at 1 N (a–e) and 10 N (f–j) in FBS. The wear tracks of all of the films were smooth and narrow at 1 N, but F4 and F5 exhibited wide and rough wear tracks at 10 N. F2 showed the narrowest wear track. FBS solution supplied a protein-rich environment. Thus, when tribo-tests were carried out in this medium, the adsorbed protein layer reduced direct contact between the counterpart and the film, improving the boundary lubrication. At an applied load of 1 N, F5 was partially broken when tested in ambient air, but it exhibited a very low wear rate of 5.6 × 10−17 m3 N−1 m−1 in FBS, indicating a very effective lubrication effect of the adsorbed protein layer. Under an edge load, the thickness of the multilayered interlayer becomes crucially important. For F4 and F5, the crack propagation resistance is insufficient, and cracks generate and propagate under the applied load, leading to breakage of the films. The H/E and H3/E2 ratios of F4 and F5 were lower than those for the other films, in accordance with the results of tribological tests. It is interesting that F2 exhibited the lowest wear rates of 2.7 × 10−17 m3 N—1 m−1 at 1 N and 7.5 × 10−17 m3 N−1 m−1 at 10 N. This indicates that in order to design an a-C-based film with high tribological performance, one effective solution is to introduce a multilayered interlayer below an a-C top layer with appropriate thickness ratio.
20
F1
F2
F3
F4
F5
0
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
a)
Fig. 9. Friction coefficients and wear rates of the films in ambient air at applied loads of a) 1 N; b) 5 N; c) 10 N.
effectively improve protein adsorption and wear properties [39]. In this work, as the CAs of the films did not vary significantly, we surmise that the influence of an adsorbed protein on the bio-tribological properties would be small. With regard to the tribological behavior of the films, their friction coefficients and the corresponding specific wear rates are shown in Fig. 9. All films exhibited very low friction coefficients of 0.09–0.15 under loads of 1 N and 5 N, although the friction coefficient of F5 was slightly higher than those of the other films. At a high applied load of 10 N, F5 was completely broken and showed a friction coefficient of about 0.5, attributable to the Ti6Al4V substrate, whereas the other films still exhibited low friction coefficients. F5 was partially broken even at a low applied load of 1 N, and the wear rate was impractically large compared to those of the other films. Therefore, the wear rate of F5 is not shown in Fig. 9. F1, F2, and F3 showed similar wear rates at a low applied load of 1 N, whereas the wear rate of F4 was
4. Conclusions This work has revealed that the mechanical and tribological properties of a-C-based films can be improved by a-C/Zr/ZrN multilayers. Additionally, the thickness ratio of the multilayered interlayer to the aC top layer is closely related with the mechanical and tribological 6
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Fig. 10. Optical images of the wear tracks of the films tested at applied loads of 1 N (a–e), 5 N (f–j) and 10 N (k–o) in ambient air.
0.20
Coefficient of friction
1N
0.15
6
0.10
3
F1
F2
F3
F4
F5
0
20
0.20 Coefficient of friction Wear rate
10 N 15
0.15 10
Broken 0.10
5
0.05
Wear rate / 10-17 m3N-1m-1
Coefficient of friction Wear rate
0.05
b)
9
F1
F2
F3
F4
F5
0
Wear rate / 10-17 m3N-1m-1
Coefficient of friction
a)
Fig. 11. Friction coefficients and wear rates of the films in FBS at applied loads of a) 1 N; b) 10 N.
properties of the films. With increasing thickness of the multilayered interlayer, toughness and adhesion are increased. A film with a 2:1 thickness ratio of multilayered interlayer to a-C top layer exhibited excellent tribological properties in both ambient air and FBS under various applied loads, especially at high applied loads. In conclusion,
we have shown that by introducing a multilayered interlayer into a-Cbased films, the toughness and adhesion can be significantly improved without adversely sacrificing hardness. The tribological properties may be optimized by carefully tailoring the thickness ratio of the multilayered interlayer to the a-C top layer. 7
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Fig. 12. Optical images of the wear tracks of the films tested at 1 N (a–e) and 10 N (f–j) in FBS.
Acknowledgements The authors gratefully acknowledge financial support from the National Natural Science Foundation of China (Grant Nos. 51665022, 51801086 and 51601081).
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