Microstructure and mechanical properties of CrN coating deposited

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
Knoop hardness test
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.
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
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,
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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.
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.
Total pressure (Pa)
N2 (ml/min)
Thickness (μm)
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
Fig. 1. Effects of gas condition on compositions of CrN coatings deposited on Ti6Al4V
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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
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:
ΔGðr Þ = 4πr σ C V +
4 3
πr ΔGV :
temperature, V the atom volume. The critical nucleation radius r⁎ can be
∂ΔGðr Þ
obtained by solving
= 0 as follows:
2σ C V V
r = − CV =
kTe lnðPo = Pe Þ
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.
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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:
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
3 −1
ln Po = ln10 × −20:68 × 10 T o −1:31 lg T o + 13:68 :
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
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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
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
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.
Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 4690–4696
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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