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Microstructure and mechanical properties of CrN coating deposited by arc ion plating on Ti6Al4V substrate
Surface & Coatings Technology 205 (2011) 4690–4696
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Microstructure and mechanical properties of CrN coating deposited by arc ion plating on Ti6Al4V substrate Z.K. Chang, X.S. Wan, Z.L. Pei, J. Gong, C. Sun ⁎ State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, PR China
a r t i c l e
i n f o
Article history: Received 14 December 2010 Accepted in revised form 5 April 2011 Available online 13 April 2011 Keywords: Knoop hardness test Interfaces Wear Arc ion plating CrN coatings Ti6Al4V alloy
a b s t r a c t CrN coatings have been grown by arc ion plating (AIP) onto Ti6Al4V alloy substrate at various nitrogen pressures (PN2 ). The goals of this investigation are to study the influence of nitrogen pressure content on the composition, structure and mechanical properties of AIP CrN coatings, as well as their tribological properties. With an increase of PN2 , the main phases in the coatings changed from CrN + Cr2N + Cr to CrN, and the texture of CrN was transformed from CrN (111)-oriented to (220)-oriented. Furthermore, the multi-layers including a metal Cr layer, a Cr2N layer and a CrN layer were observed by cross-sectional TEM (XTEM), besides an “unbalanced” state transition layer at the interface of CrN/substrate which was analyzed by nucleation thermodynamics subsequently. An increase in nitrogen pressure also resulted in a change of micro-hardness due to the variation in composition and structure. Finally, the tribological properties of the Ti6Al4V substrate and the CrN/Ti6Al4V coating system have also been explored, which shows that CrN coatings can act as good wear resistance layer for Ti6Al4V substrate. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The Ti6Al4V alloy has been used extensively in the aerospace, automotive and biomedical industries due to its attractive strengthto-weight ratio, excellent mechanical reliability, corrosion resistance and biocompatibility. However, its poor tribological behavior has limited the extension of Ti6Al4V in application areas related to wear resistance [1,2]. Transition metal nitrides, especially CrN and TiN coating, have usually been used to enhance the weak surface performance of the substrate as good wear resistance materials [3,4]. Yet, the thermal stability and corrosion resistance of CrN film are better than those of TiN film [5,6] as well as the thicker coating forming ability of CrN due to a low compressive stress state in CrN coating in contrast with a high compressive stress state in TiN coating [7], besides high micro-hardness and toughness [8,9]. Therefore, the use of CrN coating on titanium alloys could be widely used to improve friction properties and lifetime of the components in industrial application progressively. Furthermore, numerous advanced surface techniques, such as nitriding [10], ion implantation [11,12], plasma spraying [13] and physical vapor deposition [14–17], have been studied with the aim of enhancing the surface properties of the substrate. Among these, physical vapor deposition (PVD), due to its environmentally friendly characteristic, convenience and precision in deposition, has been one
⁎ Corresponding author. Tel.: + 86 24 83978081; fax: + 86 24 23843436. E-mail address: [email protected] (C. Sun). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.04.037
of the favorable techniques [18,19]. Nevertheless, only few researches [3] reported on the surface modification of a CrN-coated Ti6Al4V alloy by means of an arc ion plating (AIP) process, one of the PVD processes, and neither the interface structure between the CrN layer and the Ti6Al4V substrate nor their wear mechanisms have been investigated yet. In this study, we have deposited CrN coatings, with a Cr transitional layer in order to enhance the adhesion strength of the film/substrate [20], by arc ion plating on a Ti6Al4V substrate, and have studied not only the interfacial microstructure of the CrN/Ti6Al4V coating system and the tribological properties but also the effect of nitrogen pressure on the chemical composition, structure and mechanical performance of AIP CrN coatings. 2. Experimental details 2.1. Sample preparation All coatings were deposited in a MIP-8-800 arc ion plating system, using an evacuated chamber fitted with a round target (diameter 64 mm). The cathode target material was metallic chromium (99.9% purity). Disk samples of a commercial Ti6Al4V alloy (Al: 6.02 wt.%, V: 4.10 wt.%, Fe: 0.16 wt.%, C: 0.04 wt.% and Ti: balance) with dimensions of 15 mm in diameter and 2 mm in thickness were used as the substrate. The samples were ground with 800-mesh SiC paper and sandblasted in a wet atmosphere (200-mesh glass balls), and then ultrasonically cleaned sequentially in a metal detergent, acetone and deionized water,
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respectively. The samples were placed on the substrate holder opposite the target surface in the vacuum chamber. The targetsubstrate distance was approximately 200 mm. Prior to deposition, ion bombardment cleaning of the substrates was carried out under −900 V pulse negative bias voltage for 3 min after the base pressure of the chamber was pumped below 7.0×10− 3 Pa. After cleaning, the pulse bias voltage was reduced to −150 V in order to deposit a Cr interlayer for 5 min. During the ion bombardment cleaning and Cr interlayer deposition procedures, the atmosphere of the deposition chamber was Ar gas (99.99% purity) at 0.2 Pa. Then N2 gas (99.99% purity) was quickly introduced to maintain the chamber pressure during the CrN deposition and Ar was closed off at the same time. The deposition parameters are summarized in Table 1. A 60 A current was applied on the Cr target during the deposition. A composite power supply (pulse bias voltage of −150 V and DC bias voltage of −100 V) was employed to the substrates. The frequency of the pulse bias voltage was 20 kHz, and the duty cycle (ratio of the pulse duration time to a complete cycle period) was kept at 30%. The deposition time was 90 min.
2.2. Characterization of the coatings The surface chemical compositions of the coatings were obtained using Electron-probe microanalysis (EPMA; EPMA-1610, Shimadzu, Japan). X-ray photoelectron spectroscopy (XPS; ESCALAB 250) was carried out to observe the chemical bonding status in the Cr–N films. The XPS spectra, obtained after removing the surface layer of samples by sputtering with Ar+ ion for 60 s, were calibrated by carbon peak C 1s at 284.5 eV. The phase structures of the coatings were characterized by conventional Bragg–Brentano X-ray diffraction (XRD) using a D/max-RA type diffractometer (Cu Kα radiation, λ = 1.54056 Å). The morphology and microstructure of coatings were observed by a scanning electron microscope (SEM; S-3000N, Hitachi, Japan) coupled with emission dispersive spectroscopy (EDS; Oxford ISIS, UK). The cross-sectional morphology and diffraction patterns were obtained by a Tecnai G2 F30 transmission electron microscope (TEM). Knoop hardness (HK) measurements, using a load of 50 g and a dwelling time of 15 s, were performed using an Automatic Microindentation Hardness Testing System (Model AMH43, Japan). The reciprocating sliding wear tests were performed on a CETR UMT-2 micro-tribometer under ambient atmospheric conditions (25 ± 5 °C and 50 ± 5% RH). During the wear tests, an actual dynamic coefficient of friction was able to be obtained by the servo-controlled normal load. Si3N4 balls with a diameter of 4 mm, a surface roughness Ra of 0.02 μm and a hardness of HK50g 1600 were chosen as the wear counterparts. The wear test parameters were as follows: a normal load of 3 N; a sliding displacement amplitude of 4 mm; a sliding frequency of 4 Hz and a testing duration of 5 min. After the wear test, the wear scars of the coatings were evaluated by an Optical Surface Profiler (OSP; MicroXAM-3D, KLA-Tencor Corporation) based on the principle of light interference, Stylus-based Surface Profiler (SSP; Alpha-Step IQ, KLA-Tencor Corporation), SEM and EDS, respectively. In the stylus profilometry measurements, the scan speed, stylus force and scan length were 50 μm/s, 0.12 mN and 4 mm, respectively. The diameter of the stylus tip used in this study was 5 μm.
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3. Results and discussion 3.1. Chemical composition and structure The influence of nitrogen pressures (PN2 ) on the chemical composition of the as-deposited CrN films was analyzed by EPMA, which was the result of five different points of each sample, as shown in Fig. 1. When PN2 was 0.4 Pa, a low N content (CN ≈ 33.3 at.%) was found. With PN2 increasing over 0.8 Pa, the N content saturated at ~45 at.%. At the same time, it can be seen that the Cr content in the CrN films varied in the range 67–55% corresponding to the N content. Fig. 2 shows the XRD patterns of the CrN coatings as deposited with varying PN2 . These were composed primarily of the fcc-CrN phase (JCPDS No. 76-2494) and the mixture with hexagonal-Cr2N (JCPDS No. 35-0803) and bcc-Cr (JCPDS No. 06-0694) phases [21–23]. The coating deposited at a lower N2 pressure (0.4 Pa) shows a (111) preferred oriented CrN phase, from which a texture developed towards the (220) orientation as the transformation occurs with increasing PN2 . At the same time, diffraction patterns can be analyzed using pseudo-Voigt function profile fitting [14,15,24,25], which was inserted in Fig. 2. The intensities of the peaks corresponding to the Cr2N (111) and Cr (110) can be seen to decrease while those corresponding to the CrN (200) increased with increasing PN2 . Fig. 3 shows the fraction of sub-peaks in the Cr 2p3/2 and the N 1s XPS spectra for the CrN coatings deposited at various PN2 . The spectra were fitted by the least-squares method using a Gaussian–Lorentzian envelope. The Cr 2p3/2 spectrum can be interpreted as being composed of three species: metallic Cr0 (573.7–574.4 eV [26]), Cr in a Cr–N environment (CrN, Cr2N, 574.5 eV [27,28]), and Cr in a Cr–O environment (e.g. Cr2O3, 575.8–576.5 eV [26]). The N 1s peak was taken to be composed of 3 groups of different chemical species: CrN (396.9 eV [28]), Cr2N (397.5 eV [27]), and the smaller peaks at 399.4 ± 0.4 eV and 401.9 ± 0.4 eV which occurred in chromium nitrites/nitrates [29]. Oxygen incorporation in the Cr–O [Fig. 3(a)], found in the Cr 2p3/2 spectrum, could originate from the residual oxygen gas in the chamber in which the base pressure was at a level of 10− 3 Pa [30,31]. Meanwhile, there was a much higher fraction of Cr0 (~ 35%) in sample 1 than that of other three samples, due to insufficient reaction of N and Cr in the low nitrogen pressure condition. The fraction of Cr–N increased quickly from 47.6% to 57.9% [Fig. 3(a)] when the PN2 was increased from 0.4 to 0.8 Pa, and it was approximately 60% between 0.8 and 1.2 Pa in the Cr 2p3/2 spectrum. For the N 1s spectrum, including the CrN and Cr2N subpeaks, the relative concentration of CrN increased significantly at
Table 1 Deposition parameters and coating thickness. The errors indicate one standard deviation. Sample
Total pressure (Pa)
N2 (ml/min)
Thickness (μm)
No. No. No. No.
0.4 0.8 1.0 1.2
257 ± 3 275 ± 3 282 ± 4.5 309 ± 6
3.9 ± 0.3 4.8 ± 0.5 5.9 ± 1.0 6.3 ± 0.7
1 2 3 4
Fig. 1. Effects of gas condition on compositions of CrN coatings deposited on Ti6Al4V substrate.
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Fig. 2. XRD spectra of CrN coatings deposited on a Ti6Al4V substrate at different nitrogen pressures and the inserted image of pseudo-Voigt function profile fitting of the mixture with CrN (220), Cr2N (111) and Cr (110).
nitrogen pressures ranging from 0.4 to 1.0 Pa, while the Cr2N exhibited the opposite trend over the same range, and then both of them changed very little [Fig. 3(b)]. These phenomena are consistent with the EPMA results [Fig. 1], which mean that increasing the
nitrogen pressure results in an increase in the N content and also an increase in the number of CrN bonds. By means of the Cr–N phase diagram [32], it can be concluded that the phases present in the Cr–N coatings undergo a change from α-Cr → α-Cr + β-Cr2N → β-Cr2N → β-Cr2N + CrN → CrN with increasing nitrogen content. In our experiments, the N content of sample 1 is about 33 at.% which is in the (β-Cr2N + CrN) phase region of Cr-N phase diagram. As a result, the proportion of the Cr2N phase in sample 1 is much higher than that in the other samples [Fig. 3(b)]. It is worth noting that, although the N contents of samples 2–4 are all in the CrN single-phase region of the phase diagram, it can be still found the Cr2N phase in the XRD patterns [Fig. 2] and the Cr2N bond in the XPS results [Fig. 3]. Similarly, despite the fact that all four samples are in the (β-Cr2N + CrN) or CrN phase regions, there is also the Cr phase in the XRD patterns [Fig. 2] and Cr0 bond in the XPS [Fig. 3(a)] patterns. Insufficient reaction [14] and metal droplets [14,19] during the deposition are thought to be the reasons for the existence of the Cr2N phase and the Cr phase, respectively. Additionally, the formation of (Cr + Cr2N) transition layers is another important factor and this will be discussed in the next section below.
3.2. Interfacial structure
Fig. 3. Variations in relative intensity ratios of different chemical species of Cr and N elements in the CrN films affected by different nitrogen pressures: (a) Cr–N, Cr0 and Cr–O in Cr 2p3/2; (b) CrN and Cr2N in N 1s.
Fig. 4 shows the cross-sectional TEM (XTEM) images and selectedarea electron diffraction (SAED) pattern in the film/substrate interfacial region of the sample 4. Fig. 4(a) shows a bright field (BF) image and a SAED pattern for the CrN coatings. These coatings exhibit multi-layers including a metal Cr layer, a CrN layer and a Cr2N layer (the zone between the white dot lines and will be proved later in Fig. 4 (d)) between them. A metal Cr interlayer with a thickness of ~ 40 nm, formed during the pretreatment stage after the high pulse bias voltage cleaning process, and can be seen in the BF image. The corresponding dark field (DF) image [Fig. 4(b)], obtained from the diffraction spot of fcc-CrN (111) and hexagonal-Cr2N (110) which have the same crystallographic plane distance, exhibits a strong columnar structure consistent with the evident texture shown by XRD [Fig. 2] and SAED [Fig. 4(a)]. It can be seen from the DF image that the width of the CrN columns is about 30–50 nm perpendicular to the interface while the Cr2N layer is approximately 20 nm parallel to the interface. Fig. 4(c) shows the BF image of CrN coatings at a higher magnification. The Cr interlayer with thickness of 30–40 nm can be clearly identified. A transition layer of ~ 10 nm in thickness is observed between the Cr interlayer and the Ti6Al4V substrate (the zone between the white dot lines in Fig. 4(c)). The phase in the
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Fig. 4. XTEM micrograph of CrN coating deposited on Ti6Al4V substrate at 1.2 Pa N2 pressure: (a) bright field image (BF) with SAED (f: fcc-CrN; h: hexagonal-Cr 2 N); (b) (CrNg = (111) + Cr2 Ng = (110) ) dark field image; (c) high magnification bright field image; and (d) HRTEM image and inset image of the lattice planes for the Cr2 N interlayer network.
transition layer is not fully identified since it is too thin to be characterized by the SAED technique alone. However, we can speculate that it originates from the process of coating deposition by arc ion plating. Before the CrN coating depositing, the substrate needs to endure ion bombardment with high energy in the mode of high pulse bias voltage in order to obtain a clean substrate with some crystallographic defects (the arc cleaning process). More importantly, a nanocrystalline/amorphous layer, like the transition thin layer shown in Fig. 4(c), has been also found by others [18,20,33–35]. Petrov et al. [18] suggested that the formation of a nanocrystalline/amorphous interfacial layer might result from the high density of residual defect concentrations caused by the use of high energy ions during the etching process. In addition to the factors discussed above, the thermodynamic approach based on nucleation theory is introduced to explain the formation of this “unbalanced” state transition layer. The reversible work for crystal cluster formation ΔG(r) can be expressed as a sum of two contributions: 2
ΔGðr Þ = 4πr σ C V +
4 3 πr ΔGV : 3
temperature, V the atom volume. The critical nucleation radius r⁎ can be ∂ΔGðr Þ obtained by solving = 0 as follows: ∂r 2σ 2σ C V V : r = − CV = kTe lnðPo = Pe Þ ΔGV
ð5Þ
ð4Þ
Where σCV is the interfacial free energy per unit area between the condensed phase and the vapor phase; where ΔGV = (− kTe/V)ln(P0/Pe) is the free energy difference per unit volume between the supersaturated vapor pressure Po and the equilibrium vapor pressure Pe, k the Boltzmann constant, Te the equilibrium temperature or the substrate
Fig. 5. Effects of gas condition on micro-hardness of CrN coatings deposited on the Ti6Al4V substrate.
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3.3. Mechanical properties
Fig. 6. Friction coefficient curve of the Ti6Al4V substrate and the CrN/Ti6Al4V coating system during the ball-on-disk test.
For metal Cr, the value of vapor pressure P during 298–2130 K can be obtained from the expression [36]: 3 −1
lgðP = kPaÞ = −20:68 × 10 T
−1:31 lg T + 13:68:
ð6Þ
Using Eq. (5) in conjunction with the vapor pressure P in Eq. (6) gives the derivative of r⁎ with respect to Te: dr ⁎ = dTe =
2σ CV V ð30:19− ln Po −3:02 lg T e Þ
2 ð7Þ kTe2 ln Po − ln10× −20:68×103 T −1 e −1:31 lg T e + 13:68
and 3 −1 ln Po = ln10 × −20:68 × 10 T o −1:31 lg T o + 13:68 :
ð8Þ
At the same time, the equilibrium temperature Te is lower than To, which is in the level of 102–103 K according to the supersaturated vapor pressure Po. As a result, since dr ⁎ /dTe ≻ 0, r⁎ increases with increase Te. In the arc cleaning process, it can be concluded that as the substrate is bombarded by ions of high energy due to the high pulse bias voltage and as the substrate temperature Te rises, the critical nucleation radius r⁎ increases. In addition, the effects of resputtering of the condensing film and sputtering of the substrate material are enhanced with increasing bias voltage [33]. These two factors make it difficult for the formation and growth of a stable nucleus. As a result an “unbalanced” state transition layer is formed during this high energy ion bombardment stage. In order to explore further the details of the transition layer at the interface between the CrN and substrate, a high-resolution TEM (HRTEM) investigation has also been performed [Fig. 4(d)]. Identification of the Cr2N (111) plane between the Cr layer and the CrN layer was obtained from the lattice spacing as shown in the inset micrograph, which has also been documented by the X-ray diffraction studies [Fig. 2]. In a very short time after the deposit pretreatment of the Cr interlayer, N2 gas was gradually introduced into the vacuum chamber while Ar gas concentration was reduced to zero gradually. A thin Cr2N interlayer [Fig. 4(a) and (b)] was formed because of the low nitrogen concentration in the system at first. On the top surface of the bcc-Cr interlayer, the hexagonal-Cr2N transition layer establishes its epitaxial growth during the process of adding N2 and reducing the Ar gradually. More importantly, Cr2N phase and CrN phase have the similar interplanar spacing (dCr2 Nð110Þ ≈2:40) and dCrNð111Þ ≈2:39 )), as shown in Fig. 2 and Fig. 4(a) and (b). So this Cr2N transition layer could reduce the internal stress in maximum, which was caused by lattice mismatch between the Cr interlayer and the fcc-CrN coating.
Results from Knoop micro-hardness measurements on the Ti6Al4V substrate and the CrN/Ti6Al4V coating system are shown in Fig. 5. Due to the relatively large error, ten different points of each sample were analyzed. The hardness of Ti6Al4V substrate (HK ~ 400) was low. However, the micro-hardness is improved greatly with the CrN coatings deposited by AIP. Additionally, with the increase of the nitrogen pressure from 0.4 to 1.2 Pa, the micro-hardness monotonously increases, but which fluctuated substantially due to the rough surface caused by sandblasting and some macroparticles on the coatings. This variation may originate from three factors: firstly, solidsolution strengthening arises from the increasing of the N content [37]; secondly, the higher hardness of single phase coating than the multiphase coating [37,38] for the aforementioned reason (discussed in Section 3.1); thirdly, residual stresses [39]. In discussing the second reason, Tian and Liu [40] suggested that besides Fermi energy the energy of bonding and anti-bonding electron in the d-band of the material will be changed by composition deviation from stoichiometry. This could further influence the bonding strength and result in the phenomenon that the hardness of single phase coating is little higher than that of a multiphase coating. Consequently, the microhardness of CrN/Ti6Al4V coatings system is strongly connected with PN2 determining the structure of the coatings.
3.4. Tribological properties To see the effect of the CrN coating on the tribological behavior of the Ti6Al4V substrate, tribological tests of the CrN/Ti6Al4V system (sample 2) and of the bare Ti6Al4V substrate were performed using a micro-tribometer. Fig. 6 shows the friction coefficient of the two kinds of samples mentioned above during a ball-on-disk test. Quite different variations of the friction coefficient values were observed. The change of friction coefficients for the uncoated substrate could be divided into two stages. In the first stage from the beginning of wearing to 170 s, the friction (~ 0.35) is somewhat lower and steadier than in the second stage from 170 s to the end of wearing. In the second stage, the friction coefficients show a modest gain and the fluctuation range is more obvious than the first stage. It's attributed to that, as the test continues, the wear of the ball will slowly increase the contact area and the friction and wear become unstable due to the accumulated wear debris inside the groove [41]. For the CrN/Ti6Al4V system, it can be seen that the initial friction coefficient increased dramatically and, after some sliding, then declined to a stable value of 0.4, after reaching a maximum value of 0.65. There are a lot of macroparticles on the surface of the CrN coatings deposited by the AIP method and the surface of CrN/Ti6Al4V system becomes rougher as a result of sandblasting. The very large measured friction coefficient at the beginning of the wearing stems from the high local pressure caused by contact of the macroparticles on the coating surface and the Si3N4 balls which results in stress concentration at the surrounding of the macroparticles. These macroparticles were easily stripped by the load and the frictional force. As the result, with the prolongation of grinding-time, the surface morphology trends to become smooth and most of macroparticles on the CrN coating were stripped and flatten [Fig. 7(b) and (c)]; therefore, the friction coefficient between the coating surface and the Si3N4 balls decreased and stabilized gradually. Fig. 7(a)–(c) shows the SEM micrographs of the wear tracks for a Ti6Al4V substrate and a CrN/Ti6Al4V coating system after 5 min of friction and wear, and Fig. 7(d) shows the local chemical elements obtained by EDS analysis. The microhardness of the Si3N4 balls (HK 1400–1700) is much higher than that of Ti6Al4V substrate (HK ~ 400). As a result, the micro-morphology of the wear cracks appears furrowlike and the substrate suffered serious abrasive wear due to the
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Fig. 7. SEM micrographs and EDS analysis of wear tracks under normal load of 3 N: (a) Ti6Al4V substrate, (b) and (c) CrN/Ti6Al4V coating system, and (d) EDS spectrum corresponding to the point X denoted in (b).
cutting by rough peaks on the surface of the Si3N4 balls during the process of friction and wear, as shown in Fig. 7(a). The local surface morphology of the CrN/Ti6Al4V coating system after wear is represented in Fig. 7(b) and (c). It is noted that the surface of wear tracks presents a scale-like morphology as well as coating delamination and cracks [Fig. 7(c)]. During the abrasion, the accumulation and propagation of fatigue tensile cracks caused by stress on the CrN coatings under the action of the Si3N4 balls lead to the detachment of the coatings [41]. The CrN coatings were stripped and Si3N4 wear debris was milled into scale-like morphology. The EDS analysis of the “scale” in the X-position [Fig. 7(c)] is shown in Fig. 7(d). Besides the coatings' elements, there was a relatively high content of Si from the Si3N4 balls and O from a thin oxide tribolayer formed by the high flash temperature during sliding [16]. No sharp groove was observed in the wear scar, which is mainly attributed to the fact that the micro-hardness of CrN coatings (HK ~1800) is higher than that of the Si3N4 ball. The ball-on-disk wear mechanisms of the CrN coatings on the Ti6Al4V substrate are identified as stress cracks, coating stripping and oxidative wear. In order to acquire the qualitative comparison of abrasion loss between the Ti6Al4V substrate and the CrN/Ti6Al4V coating system, an Optical Surface Profiler (OSP) was used to observe the wear track morphologies. Fig. 8(a) and (b) displays the 3D wear track morphologies of Ti6Al4V substrate and CrN/Ti6Al4V coating system respectively. The light and deep color zones stand for higher and lower positions on the Z-axis respectively. From the wear scar profiles shown in Fig. 8(c) and (d), deep grinding cracks (~7 μm) were produced on the Ti6Al4V substrate, whereas there was no obvious wear track in the CrN/Ti6Al4V coating system. This result demonstrates that CrN coatings are an excellent wear-resistance material to effectively protect Ti6Al4V substrate. 4. Conclusions We have presented experimental results and discussed the mechanisms of the way which nitrogen pressure affects the chemical
composition, structure and mechanical performance of AIP CrN coatings. In addition, an analysis of the interfacial microstructure of the film/substrate system and comparison of the tribological properties of Ti6Al4V substrates with and without coating have been carried out. The main results can be summarized as follows: 1. The N contents in CrN coating increase with increase in PN2 . Accordingly, the main phases in the films are transformed from CrN + Cr2N + Cr to CrN with increasing nitrogen pressure. This is based on the phase diagram for the Cr–N system. And an initial (111)-dominated texture of CrN is changes into the (220) surfaces gradually. 2. Multi-layer structure of (Cr + Cr2N) interlays and CrN layer has been observed by XTEM and HRTEM. In particularly, we have observed and theoretically proved the existence of a thin nanocrystalline/amorphous like transition layer between coating and the substrate. 3. The micro-hardness of AIP CrN films increases from 1550 to 2100 (HK50g) as nitrogen pressure increases from 0.4 to 1.2 Pa, following the same trend as N and CrN contents, which is mainly ascribed to the variation of composition and structure. 4. For CrN coatings, a combination of several wear mechanisms by stress cracks, delamination and oxidative has been considered under a normal load of 3 N. The application of AIP CrN coatings with their anti-abrasive and high micro-hardness performance could significantly expand the application range of Ti6Al4V alloys.
Acknowledgments The authors thank Dr. D. M. Tang, Dr. C. F. Li, Dr. B. Yang and Dr. S. J. Wang for valuable discussions about TEM during this work at IMR. Furthermore, the authors acknowledge Professor William Alan Oates (The University of Newcastle, Australia) and Dr. Z. S. You (IMR) for checking this paper.
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Fig. 8. 3D-surface topographies and wear track profiles of Ti6Al4V substrate (a), (c) and CrN/Ti6Al4V coating system (b), (d).
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[21] W.D. Munz, M. Schenkel, S. Kunkel, K. Bewilogua, M. Keunecke, R. Wittorf, SVC 50th Annual Technical Conference Proceedings, 2007, p. 155. [22] T. Hurkmans, D.B. Lewis, H. Paritong, J.S. Brooks, W.D. Munz, Surf. Coat. Technol. 114 (1999) 52. [23] T. Hurkmans, D.B. Lewis, J.S. Brooks, W.D. Munz, Surf. Coat. Technol. 86–87 (1996) 192. [24] H. Toraya, J. Appl. Crystallogr. 23 (1990) 485. [25] W.I.F. David, J. Appl. Crystallogr. 19 (1986) 63. [26] B. Stypula, J. Stoch, Corros. Sci. 36 (1994) 2159. [27] I. Bertóti, Surf. Coat. Technol. 151–152 (2002) 194. [28] C. Emery, A.R. Chourasia, P. Yashar, J. Electron. Spectrosc. Relat. Phenom. 104 (1999) 91. [29] http://stdata.nist.gov/xps/. [30] Q.M. Wang, K.H. Kim, J. Vac, Sci. Technol. A 26 (2008) 1267. [31] A. Lippitz, T. Hübert, Surf. Coat. Technol. 200 (2005) 250. [32] V.G. Ivanchenko, T.V. Mel'nichenko, Metallofizika 13 (1991) 23. [33] C. Schonjahn, H. Paritong, W.D. Munz, R.D. Twesten, I. Petrov, J. Vac, Sci. Technol. A 19 (2001) 1392. [34] A.P. Ehiasarian, J.G. Wen, I. Petrov, J. Appl. Phys. 101 (2007) 54301. [35] S.S. Zhao, Y. Yang, J.B. Li, J. Gong, C. Sun, Surf. Coat. Technol. 202 (2008) 5185. [36] O. Kubaschewski, C.B. Alcock (Eds.), Metallurgical Thermochemistry, 5th ed, Pergamon Press, Oxford, 1979. [37] C. Rebholz, H. Ziegele, A. Leyland, A. Matthews, Surf. Coat. Technol. 115 (1999) 222. [38] P. Hones, R. Sanjines, F. Levy, Surf. Coat. Technol. 94–95 (1997) 398. [39] X.S. Wan, S.S. Zhao, Y. Yang, J. Gong, C. Sun, Surf. Coat. Technol. 204 (2010) 1800. [40] M.B. Tian, D.L. Liu (Eds.), Handbook of Thin Film Science and Technology, China Machine Press, Peking, 1991. [41] E. Martínez, J. Romero, A. Lousa, J. Esteve, Surf. Coat. Technol. 163–164 (2003) 571.
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Microstructure and mechanical properties of CrN coating deposited by arc ion plating on Ti6Al4V substrate Z.K. Chang, X.S. Wan, Z.L. Pei, J. Gong, C. Sun ⁎ State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, PR China
a r t i c l e
i n f o
Article history: Received 14 December 2010 Accepted in revised form 5 April 2011 Available online 13 April 2011 Keywords: Knoop hardness test Interfaces Wear Arc ion plating CrN coatings Ti6Al4V alloy
a b s t r a c t CrN coatings have been grown by arc ion plating (AIP) onto Ti6Al4V alloy substrate at various nitrogen pressures (PN2 ). The goals of this investigation are to study the influence of nitrogen pressure content on the composition, structure and mechanical properties of AIP CrN coatings, as well as their tribological properties. With an increase of PN2 , the main phases in the coatings changed from CrN + Cr2N + Cr to CrN, and the texture of CrN was transformed from CrN (111)-oriented to (220)-oriented. Furthermore, the multi-layers including a metal Cr layer, a Cr2N layer and a CrN layer were observed by cross-sectional TEM (XTEM), besides an “unbalanced” state transition layer at the interface of CrN/substrate which was analyzed by nucleation thermodynamics subsequently. An increase in nitrogen pressure also resulted in a change of micro-hardness due to the variation in composition and structure. Finally, the tribological properties of the Ti6Al4V substrate and the CrN/Ti6Al4V coating system have also been explored, which shows that CrN coatings can act as good wear resistance layer for Ti6Al4V substrate. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The Ti6Al4V alloy has been used extensively in the aerospace, automotive and biomedical industries due to its attractive strengthto-weight ratio, excellent mechanical reliability, corrosion resistance and biocompatibility. However, its poor tribological behavior has limited the extension of Ti6Al4V in application areas related to wear resistance [1,2]. Transition metal nitrides, especially CrN and TiN coating, have usually been used to enhance the weak surface performance of the substrate as good wear resistance materials [3,4]. Yet, the thermal stability and corrosion resistance of CrN film are better than those of TiN film [5,6] as well as the thicker coating forming ability of CrN due to a low compressive stress state in CrN coating in contrast with a high compressive stress state in TiN coating [7], besides high micro-hardness and toughness [8,9]. Therefore, the use of CrN coating on titanium alloys could be widely used to improve friction properties and lifetime of the components in industrial application progressively. Furthermore, numerous advanced surface techniques, such as nitriding [10], ion implantation [11,12], plasma spraying [13] and physical vapor deposition [14–17], have been studied with the aim of enhancing the surface properties of the substrate. Among these, physical vapor deposition (PVD), due to its environmentally friendly characteristic, convenience and precision in deposition, has been one
⁎ Corresponding author. Tel.: + 86 24 83978081; fax: + 86 24 23843436. E-mail address: [email protected] (C. Sun). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.04.037
of the favorable techniques [18,19]. Nevertheless, only few researches [3] reported on the surface modification of a CrN-coated Ti6Al4V alloy by means of an arc ion plating (AIP) process, one of the PVD processes, and neither the interface structure between the CrN layer and the Ti6Al4V substrate nor their wear mechanisms have been investigated yet. In this study, we have deposited CrN coatings, with a Cr transitional layer in order to enhance the adhesion strength of the film/substrate [20], by arc ion plating on a Ti6Al4V substrate, and have studied not only the interfacial microstructure of the CrN/Ti6Al4V coating system and the tribological properties but also the effect of nitrogen pressure on the chemical composition, structure and mechanical performance of AIP CrN coatings. 2. Experimental details 2.1. Sample preparation All coatings were deposited in a MIP-8-800 arc ion plating system, using an evacuated chamber fitted with a round target (diameter 64 mm). The cathode target material was metallic chromium (99.9% purity). Disk samples of a commercial Ti6Al4V alloy (Al: 6.02 wt.%, V: 4.10 wt.%, Fe: 0.16 wt.%, C: 0.04 wt.% and Ti: balance) with dimensions of 15 mm in diameter and 2 mm in thickness were used as the substrate. The samples were ground with 800-mesh SiC paper and sandblasted in a wet atmosphere (200-mesh glass balls), and then ultrasonically cleaned sequentially in a metal detergent, acetone and deionized water,
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respectively. The samples were placed on the substrate holder opposite the target surface in the vacuum chamber. The targetsubstrate distance was approximately 200 mm. Prior to deposition, ion bombardment cleaning of the substrates was carried out under −900 V pulse negative bias voltage for 3 min after the base pressure of the chamber was pumped below 7.0×10− 3 Pa. After cleaning, the pulse bias voltage was reduced to −150 V in order to deposit a Cr interlayer for 5 min. During the ion bombardment cleaning and Cr interlayer deposition procedures, the atmosphere of the deposition chamber was Ar gas (99.99% purity) at 0.2 Pa. Then N2 gas (99.99% purity) was quickly introduced to maintain the chamber pressure during the CrN deposition and Ar was closed off at the same time. The deposition parameters are summarized in Table 1. A 60 A current was applied on the Cr target during the deposition. A composite power supply (pulse bias voltage of −150 V and DC bias voltage of −100 V) was employed to the substrates. The frequency of the pulse bias voltage was 20 kHz, and the duty cycle (ratio of the pulse duration time to a complete cycle period) was kept at 30%. The deposition time was 90 min.
2.2. Characterization of the coatings The surface chemical compositions of the coatings were obtained using Electron-probe microanalysis (EPMA; EPMA-1610, Shimadzu, Japan). X-ray photoelectron spectroscopy (XPS; ESCALAB 250) was carried out to observe the chemical bonding status in the Cr–N films. The XPS spectra, obtained after removing the surface layer of samples by sputtering with Ar+ ion for 60 s, were calibrated by carbon peak C 1s at 284.5 eV. The phase structures of the coatings were characterized by conventional Bragg–Brentano X-ray diffraction (XRD) using a D/max-RA type diffractometer (Cu Kα radiation, λ = 1.54056 Å). The morphology and microstructure of coatings were observed by a scanning electron microscope (SEM; S-3000N, Hitachi, Japan) coupled with emission dispersive spectroscopy (EDS; Oxford ISIS, UK). The cross-sectional morphology and diffraction patterns were obtained by a Tecnai G2 F30 transmission electron microscope (TEM). Knoop hardness (HK) measurements, using a load of 50 g and a dwelling time of 15 s, were performed using an Automatic Microindentation Hardness Testing System (Model AMH43, Japan). The reciprocating sliding wear tests were performed on a CETR UMT-2 micro-tribometer under ambient atmospheric conditions (25 ± 5 °C and 50 ± 5% RH). During the wear tests, an actual dynamic coefficient of friction was able to be obtained by the servo-controlled normal load. Si3N4 balls with a diameter of 4 mm, a surface roughness Ra of 0.02 μm and a hardness of HK50g 1600 were chosen as the wear counterparts. The wear test parameters were as follows: a normal load of 3 N; a sliding displacement amplitude of 4 mm; a sliding frequency of 4 Hz and a testing duration of 5 min. After the wear test, the wear scars of the coatings were evaluated by an Optical Surface Profiler (OSP; MicroXAM-3D, KLA-Tencor Corporation) based on the principle of light interference, Stylus-based Surface Profiler (SSP; Alpha-Step IQ, KLA-Tencor Corporation), SEM and EDS, respectively. In the stylus profilometry measurements, the scan speed, stylus force and scan length were 50 μm/s, 0.12 mN and 4 mm, respectively. The diameter of the stylus tip used in this study was 5 μm.
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3. Results and discussion 3.1. Chemical composition and structure The influence of nitrogen pressures (PN2 ) on the chemical composition of the as-deposited CrN films was analyzed by EPMA, which was the result of five different points of each sample, as shown in Fig. 1. When PN2 was 0.4 Pa, a low N content (CN ≈ 33.3 at.%) was found. With PN2 increasing over 0.8 Pa, the N content saturated at ~45 at.%. At the same time, it can be seen that the Cr content in the CrN films varied in the range 67–55% corresponding to the N content. Fig. 2 shows the XRD patterns of the CrN coatings as deposited with varying PN2 . These were composed primarily of the fcc-CrN phase (JCPDS No. 76-2494) and the mixture with hexagonal-Cr2N (JCPDS No. 35-0803) and bcc-Cr (JCPDS No. 06-0694) phases [21–23]. The coating deposited at a lower N2 pressure (0.4 Pa) shows a (111) preferred oriented CrN phase, from which a texture developed towards the (220) orientation as the transformation occurs with increasing PN2 . At the same time, diffraction patterns can be analyzed using pseudo-Voigt function profile fitting [14,15,24,25], which was inserted in Fig. 2. The intensities of the peaks corresponding to the Cr2N (111) and Cr (110) can be seen to decrease while those corresponding to the CrN (200) increased with increasing PN2 . Fig. 3 shows the fraction of sub-peaks in the Cr 2p3/2 and the N 1s XPS spectra for the CrN coatings deposited at various PN2 . The spectra were fitted by the least-squares method using a Gaussian–Lorentzian envelope. The Cr 2p3/2 spectrum can be interpreted as being composed of three species: metallic Cr0 (573.7–574.4 eV [26]), Cr in a Cr–N environment (CrN, Cr2N, 574.5 eV [27,28]), and Cr in a Cr–O environment (e.g. Cr2O3, 575.8–576.5 eV [26]). The N 1s peak was taken to be composed of 3 groups of different chemical species: CrN (396.9 eV [28]), Cr2N (397.5 eV [27]), and the smaller peaks at 399.4 ± 0.4 eV and 401.9 ± 0.4 eV which occurred in chromium nitrites/nitrates [29]. Oxygen incorporation in the Cr–O [Fig. 3(a)], found in the Cr 2p3/2 spectrum, could originate from the residual oxygen gas in the chamber in which the base pressure was at a level of 10− 3 Pa [30,31]. Meanwhile, there was a much higher fraction of Cr0 (~ 35%) in sample 1 than that of other three samples, due to insufficient reaction of N and Cr in the low nitrogen pressure condition. The fraction of Cr–N increased quickly from 47.6% to 57.9% [Fig. 3(a)] when the PN2 was increased from 0.4 to 0.8 Pa, and it was approximately 60% between 0.8 and 1.2 Pa in the Cr 2p3/2 spectrum. For the N 1s spectrum, including the CrN and Cr2N subpeaks, the relative concentration of CrN increased significantly at
Table 1 Deposition parameters and coating thickness. The errors indicate one standard deviation. Sample
Total pressure (Pa)
N2 (ml/min)
Thickness (μm)
No. No. No. No.
0.4 0.8 1.0 1.2
257 ± 3 275 ± 3 282 ± 4.5 309 ± 6
3.9 ± 0.3 4.8 ± 0.5 5.9 ± 1.0 6.3 ± 0.7
1 2 3 4
Fig. 1. Effects of gas condition on compositions of CrN coatings deposited on Ti6Al4V substrate.
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Fig. 2. XRD spectra of CrN coatings deposited on a Ti6Al4V substrate at different nitrogen pressures and the inserted image of pseudo-Voigt function profile fitting of the mixture with CrN (220), Cr2N (111) and Cr (110).
nitrogen pressures ranging from 0.4 to 1.0 Pa, while the Cr2N exhibited the opposite trend over the same range, and then both of them changed very little [Fig. 3(b)]. These phenomena are consistent with the EPMA results [Fig. 1], which mean that increasing the
nitrogen pressure results in an increase in the N content and also an increase in the number of CrN bonds. By means of the Cr–N phase diagram [32], it can be concluded that the phases present in the Cr–N coatings undergo a change from α-Cr → α-Cr + β-Cr2N → β-Cr2N → β-Cr2N + CrN → CrN with increasing nitrogen content. In our experiments, the N content of sample 1 is about 33 at.% which is in the (β-Cr2N + CrN) phase region of Cr-N phase diagram. As a result, the proportion of the Cr2N phase in sample 1 is much higher than that in the other samples [Fig. 3(b)]. It is worth noting that, although the N contents of samples 2–4 are all in the CrN single-phase region of the phase diagram, it can be still found the Cr2N phase in the XRD patterns [Fig. 2] and the Cr2N bond in the XPS results [Fig. 3]. Similarly, despite the fact that all four samples are in the (β-Cr2N + CrN) or CrN phase regions, there is also the Cr phase in the XRD patterns [Fig. 2] and Cr0 bond in the XPS [Fig. 3(a)] patterns. Insufficient reaction [14] and metal droplets [14,19] during the deposition are thought to be the reasons for the existence of the Cr2N phase and the Cr phase, respectively. Additionally, the formation of (Cr + Cr2N) transition layers is another important factor and this will be discussed in the next section below.
3.2. Interfacial structure
Fig. 3. Variations in relative intensity ratios of different chemical species of Cr and N elements in the CrN films affected by different nitrogen pressures: (a) Cr–N, Cr0 and Cr–O in Cr 2p3/2; (b) CrN and Cr2N in N 1s.
Fig. 4 shows the cross-sectional TEM (XTEM) images and selectedarea electron diffraction (SAED) pattern in the film/substrate interfacial region of the sample 4. Fig. 4(a) shows a bright field (BF) image and a SAED pattern for the CrN coatings. These coatings exhibit multi-layers including a metal Cr layer, a CrN layer and a Cr2N layer (the zone between the white dot lines and will be proved later in Fig. 4 (d)) between them. A metal Cr interlayer with a thickness of ~ 40 nm, formed during the pretreatment stage after the high pulse bias voltage cleaning process, and can be seen in the BF image. The corresponding dark field (DF) image [Fig. 4(b)], obtained from the diffraction spot of fcc-CrN (111) and hexagonal-Cr2N (110) which have the same crystallographic plane distance, exhibits a strong columnar structure consistent with the evident texture shown by XRD [Fig. 2] and SAED [Fig. 4(a)]. It can be seen from the DF image that the width of the CrN columns is about 30–50 nm perpendicular to the interface while the Cr2N layer is approximately 20 nm parallel to the interface. Fig. 4(c) shows the BF image of CrN coatings at a higher magnification. The Cr interlayer with thickness of 30–40 nm can be clearly identified. A transition layer of ~ 10 nm in thickness is observed between the Cr interlayer and the Ti6Al4V substrate (the zone between the white dot lines in Fig. 4(c)). The phase in the
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Fig. 4. XTEM micrograph of CrN coating deposited on Ti6Al4V substrate at 1.2 Pa N2 pressure: (a) bright field image (BF) with SAED (f: fcc-CrN; h: hexagonal-Cr 2 N); (b) (CrNg = (111) + Cr2 Ng = (110) ) dark field image; (c) high magnification bright field image; and (d) HRTEM image and inset image of the lattice planes for the Cr2 N interlayer network.
transition layer is not fully identified since it is too thin to be characterized by the SAED technique alone. However, we can speculate that it originates from the process of coating deposition by arc ion plating. Before the CrN coating depositing, the substrate needs to endure ion bombardment with high energy in the mode of high pulse bias voltage in order to obtain a clean substrate with some crystallographic defects (the arc cleaning process). More importantly, a nanocrystalline/amorphous layer, like the transition thin layer shown in Fig. 4(c), has been also found by others [18,20,33–35]. Petrov et al. [18] suggested that the formation of a nanocrystalline/amorphous interfacial layer might result from the high density of residual defect concentrations caused by the use of high energy ions during the etching process. In addition to the factors discussed above, the thermodynamic approach based on nucleation theory is introduced to explain the formation of this “unbalanced” state transition layer. The reversible work for crystal cluster formation ΔG(r) can be expressed as a sum of two contributions: 2
ΔGðr Þ = 4πr σ C V +
4 3 πr ΔGV : 3
temperature, V the atom volume. The critical nucleation radius r⁎ can be ∂ΔGðr Þ obtained by solving = 0 as follows: ∂r 2σ 2σ C V V : r = − CV = kTe lnðPo = Pe Þ ΔGV
ð5Þ
ð4Þ
Where σCV is the interfacial free energy per unit area between the condensed phase and the vapor phase; where ΔGV = (− kTe/V)ln(P0/Pe) is the free energy difference per unit volume between the supersaturated vapor pressure Po and the equilibrium vapor pressure Pe, k the Boltzmann constant, Te the equilibrium temperature or the substrate
Fig. 5. Effects of gas condition on micro-hardness of CrN coatings deposited on the Ti6Al4V substrate.
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3.3. Mechanical properties
Fig. 6. Friction coefficient curve of the Ti6Al4V substrate and the CrN/Ti6Al4V coating system during the ball-on-disk test.
For metal Cr, the value of vapor pressure P during 298–2130 K can be obtained from the expression [36]: 3 −1
lgðP = kPaÞ = −20:68 × 10 T
−1:31 lg T + 13:68:
ð6Þ
Using Eq. (5) in conjunction with the vapor pressure P in Eq. (6) gives the derivative of r⁎ with respect to Te: dr ⁎ = dTe =
2σ CV V ð30:19− ln Po −3:02 lg T e Þ
2 ð7Þ kTe2 ln Po − ln10× −20:68×103 T −1 e −1:31 lg T e + 13:68
and 3 −1 ln Po = ln10 × −20:68 × 10 T o −1:31 lg T o + 13:68 :
ð8Þ
At the same time, the equilibrium temperature Te is lower than To, which is in the level of 102–103 K according to the supersaturated vapor pressure Po. As a result, since dr ⁎ /dTe ≻ 0, r⁎ increases with increase Te. In the arc cleaning process, it can be concluded that as the substrate is bombarded by ions of high energy due to the high pulse bias voltage and as the substrate temperature Te rises, the critical nucleation radius r⁎ increases. In addition, the effects of resputtering of the condensing film and sputtering of the substrate material are enhanced with increasing bias voltage [33]. These two factors make it difficult for the formation and growth of a stable nucleus. As a result an “unbalanced” state transition layer is formed during this high energy ion bombardment stage. In order to explore further the details of the transition layer at the interface between the CrN and substrate, a high-resolution TEM (HRTEM) investigation has also been performed [Fig. 4(d)]. Identification of the Cr2N (111) plane between the Cr layer and the CrN layer was obtained from the lattice spacing as shown in the inset micrograph, which has also been documented by the X-ray diffraction studies [Fig. 2]. In a very short time after the deposit pretreatment of the Cr interlayer, N2 gas was gradually introduced into the vacuum chamber while Ar gas concentration was reduced to zero gradually. A thin Cr2N interlayer [Fig. 4(a) and (b)] was formed because of the low nitrogen concentration in the system at first. On the top surface of the bcc-Cr interlayer, the hexagonal-Cr2N transition layer establishes its epitaxial growth during the process of adding N2 and reducing the Ar gradually. More importantly, Cr2N phase and CrN phase have the similar interplanar spacing (dCr2 Nð110Þ ≈2:40) and dCrNð111Þ ≈2:39 )), as shown in Fig. 2 and Fig. 4(a) and (b). So this Cr2N transition layer could reduce the internal stress in maximum, which was caused by lattice mismatch between the Cr interlayer and the fcc-CrN coating.
Results from Knoop micro-hardness measurements on the Ti6Al4V substrate and the CrN/Ti6Al4V coating system are shown in Fig. 5. Due to the relatively large error, ten different points of each sample were analyzed. The hardness of Ti6Al4V substrate (HK ~ 400) was low. However, the micro-hardness is improved greatly with the CrN coatings deposited by AIP. Additionally, with the increase of the nitrogen pressure from 0.4 to 1.2 Pa, the micro-hardness monotonously increases, but which fluctuated substantially due to the rough surface caused by sandblasting and some macroparticles on the coatings. This variation may originate from three factors: firstly, solidsolution strengthening arises from the increasing of the N content [37]; secondly, the higher hardness of single phase coating than the multiphase coating [37,38] for the aforementioned reason (discussed in Section 3.1); thirdly, residual stresses [39]. In discussing the second reason, Tian and Liu [40] suggested that besides Fermi energy the energy of bonding and anti-bonding electron in the d-band of the material will be changed by composition deviation from stoichiometry. This could further influence the bonding strength and result in the phenomenon that the hardness of single phase coating is little higher than that of a multiphase coating. Consequently, the microhardness of CrN/Ti6Al4V coatings system is strongly connected with PN2 determining the structure of the coatings.
3.4. Tribological properties To see the effect of the CrN coating on the tribological behavior of the Ti6Al4V substrate, tribological tests of the CrN/Ti6Al4V system (sample 2) and of the bare Ti6Al4V substrate were performed using a micro-tribometer. Fig. 6 shows the friction coefficient of the two kinds of samples mentioned above during a ball-on-disk test. Quite different variations of the friction coefficient values were observed. The change of friction coefficients for the uncoated substrate could be divided into two stages. In the first stage from the beginning of wearing to 170 s, the friction (~ 0.35) is somewhat lower and steadier than in the second stage from 170 s to the end of wearing. In the second stage, the friction coefficients show a modest gain and the fluctuation range is more obvious than the first stage. It's attributed to that, as the test continues, the wear of the ball will slowly increase the contact area and the friction and wear become unstable due to the accumulated wear debris inside the groove [41]. For the CrN/Ti6Al4V system, it can be seen that the initial friction coefficient increased dramatically and, after some sliding, then declined to a stable value of 0.4, after reaching a maximum value of 0.65. There are a lot of macroparticles on the surface of the CrN coatings deposited by the AIP method and the surface of CrN/Ti6Al4V system becomes rougher as a result of sandblasting. The very large measured friction coefficient at the beginning of the wearing stems from the high local pressure caused by contact of the macroparticles on the coating surface and the Si3N4 balls which results in stress concentration at the surrounding of the macroparticles. These macroparticles were easily stripped by the load and the frictional force. As the result, with the prolongation of grinding-time, the surface morphology trends to become smooth and most of macroparticles on the CrN coating were stripped and flatten [Fig. 7(b) and (c)]; therefore, the friction coefficient between the coating surface and the Si3N4 balls decreased and stabilized gradually. Fig. 7(a)–(c) shows the SEM micrographs of the wear tracks for a Ti6Al4V substrate and a CrN/Ti6Al4V coating system after 5 min of friction and wear, and Fig. 7(d) shows the local chemical elements obtained by EDS analysis. The microhardness of the Si3N4 balls (HK 1400–1700) is much higher than that of Ti6Al4V substrate (HK ~ 400). As a result, the micro-morphology of the wear cracks appears furrowlike and the substrate suffered serious abrasive wear due to the
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Fig. 7. SEM micrographs and EDS analysis of wear tracks under normal load of 3 N: (a) Ti6Al4V substrate, (b) and (c) CrN/Ti6Al4V coating system, and (d) EDS spectrum corresponding to the point X denoted in (b).
cutting by rough peaks on the surface of the Si3N4 balls during the process of friction and wear, as shown in Fig. 7(a). The local surface morphology of the CrN/Ti6Al4V coating system after wear is represented in Fig. 7(b) and (c). It is noted that the surface of wear tracks presents a scale-like morphology as well as coating delamination and cracks [Fig. 7(c)]. During the abrasion, the accumulation and propagation of fatigue tensile cracks caused by stress on the CrN coatings under the action of the Si3N4 balls lead to the detachment of the coatings [41]. The CrN coatings were stripped and Si3N4 wear debris was milled into scale-like morphology. The EDS analysis of the “scale” in the X-position [Fig. 7(c)] is shown in Fig. 7(d). Besides the coatings' elements, there was a relatively high content of Si from the Si3N4 balls and O from a thin oxide tribolayer formed by the high flash temperature during sliding [16]. No sharp groove was observed in the wear scar, which is mainly attributed to the fact that the micro-hardness of CrN coatings (HK ~1800) is higher than that of the Si3N4 ball. The ball-on-disk wear mechanisms of the CrN coatings on the Ti6Al4V substrate are identified as stress cracks, coating stripping and oxidative wear. In order to acquire the qualitative comparison of abrasion loss between the Ti6Al4V substrate and the CrN/Ti6Al4V coating system, an Optical Surface Profiler (OSP) was used to observe the wear track morphologies. Fig. 8(a) and (b) displays the 3D wear track morphologies of Ti6Al4V substrate and CrN/Ti6Al4V coating system respectively. The light and deep color zones stand for higher and lower positions on the Z-axis respectively. From the wear scar profiles shown in Fig. 8(c) and (d), deep grinding cracks (~7 μm) were produced on the Ti6Al4V substrate, whereas there was no obvious wear track in the CrN/Ti6Al4V coating system. This result demonstrates that CrN coatings are an excellent wear-resistance material to effectively protect Ti6Al4V substrate. 4. Conclusions We have presented experimental results and discussed the mechanisms of the way which nitrogen pressure affects the chemical
composition, structure and mechanical performance of AIP CrN coatings. In addition, an analysis of the interfacial microstructure of the film/substrate system and comparison of the tribological properties of Ti6Al4V substrates with and without coating have been carried out. The main results can be summarized as follows: 1. The N contents in CrN coating increase with increase in PN2 . Accordingly, the main phases in the films are transformed from CrN + Cr2N + Cr to CrN with increasing nitrogen pressure. This is based on the phase diagram for the Cr–N system. And an initial (111)-dominated texture of CrN is changes into the (220) surfaces gradually. 2. Multi-layer structure of (Cr + Cr2N) interlays and CrN layer has been observed by XTEM and HRTEM. In particularly, we have observed and theoretically proved the existence of a thin nanocrystalline/amorphous like transition layer between coating and the substrate. 3. The micro-hardness of AIP CrN films increases from 1550 to 2100 (HK50g) as nitrogen pressure increases from 0.4 to 1.2 Pa, following the same trend as N and CrN contents, which is mainly ascribed to the variation of composition and structure. 4. For CrN coatings, a combination of several wear mechanisms by stress cracks, delamination and oxidative has been considered under a normal load of 3 N. The application of AIP CrN coatings with their anti-abrasive and high micro-hardness performance could significantly expand the application range of Ti6Al4V alloys.
Acknowledgments The authors thank Dr. D. M. Tang, Dr. C. F. Li, Dr. B. Yang and Dr. S. J. Wang for valuable discussions about TEM during this work at IMR. Furthermore, the authors acknowledge Professor William Alan Oates (The University of Newcastle, Australia) and Dr. Z. S. You (IMR) for checking this paper.
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Fig. 8. 3D-surface topographies and wear track profiles of Ti6Al4V substrate (a), (c) and CrN/Ti6Al4V coating system (b), (d).
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