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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
Contents lists available at ScienceDirect
International Journal of Refractory Metals
& Hard Materials
journal homepage: www.elsevier.com/locate/IJRMHM
Residual stresses and HRTEM phase interface structures of AlN structured
PVD coatings on plate-grained WC–Co substrate
T
⁎
Li Zhang , Zhi-Qiang Zhong, Yi Chen, Qiao-Ping Xiao, Guo-Kai Luo, Jun-Jie Zhu
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
A R T I C LE I N FO
A B S T R A C T
Keywords:
Plate-grained cemented carbide
Coating
Residual stress
Phase transformation
Interface microstructure
HRTEM
A fine plate-like grained WC–12Co based alloy with (0001)WC texture coefficient of 2.73 for the sinter skin was
used as the substrate. Single layer Al0.66Cr0.34N and two layers Al0.65Cr0.30Si0.04W0.01N/Al0.66Cr0.34N
(AlCrSiWN/ AlCrN) approximate 1.6 μm in total thickness for each were separately deposited on the substrate in
an industrial-scale arc evaporation system. This study aims to investigate the strengthening mechanism for the
adhesion and the resistance against cohesive failure of the films and to prove a hypothesis of stress buffering
function of the binder phase in the substrate to the film. Accordingly, coating related phase structure and residual stress change in the substrate and HRTEM phase interface structures of the film/substrate system were
investigated. The results show that stress buffering by the Co phase (binder phase) is realized by energy absorption in a form of residual stress build-up and fcc to hcp phase transformation, accompanied by an atomic
disorder Co zone in several atomic layer thickness adjacent to the AlCrN/Co interface. Coherent and semicoherent epitaxial growth of AlCrN on WC surfaces is observed and the superiority of plate-grained WC–Co
substrate is confirmed. Methods of exploiting the potential of the stress buffering function are proposed.
1. Introduction
Due to the growing demands for high speed, high efficiency and
high precision machining, there is a strong diving force to exploit the
full potential of wear resistant hard coatings [1]. With the continuous
progress of coating characterization technology, the development of
coating technology has been greatly promoted. Among all the characterization technologies, e.g., chemical composition and microstructure [2, 3], mechanical and tribological properties [4–6], thermophysical properties [7, 8], the contribution of microstructure characterization is particularly prominent in the progress of coating technology.
Cemented carbide substrate dominates the phase interface microstructure and the synergistic effect of coating (film)/ substrate system.
Therefore, it has been widely accepted that without comprehensive
research on the substrate, the ultimate coating performance optimization is difficult to achieve [9, 10]. Nevertheless, research on the phase
interface microstructures between the film and the hard/binder phase
in cemented carbide substrate was rarely reported, and the synergistic
effect of the hard phase and the ductile binder phase in cemented
carbide substrate on the coatings is still undefined. Admittedly, due to
the substantial heterogeneity between the film and the substrate, during
the transmission electron microscopy (TEM) specimen preparation
⁎
process, interface trap is easily formed. Certainly, this trap can lead to
an uncertainty for the interface investigation, especially for the detailed
atomic arrangement at the plane of the interface. Irregularity of the
interface caused by inappropriate surface treatment of the substrate
before the film deposition aggravates this adverse situation.
We have reported [11] that using a fine plate-like grained cemented
carbide as the substrate, > 100 N resistance against cohesive failure
strength for AlN based coatings are achieved. Nevertheless, the related
strengthening mechanism is not still clarified. By comparing the resistance against cohesive failure strength between a binderless carbide
hard material and a WC–Co cemented carbide with the same coating,
we have assumed a stress buffering function of the ductile binder phase
to the film [10]. Nevertheless, this hypothesis still needs to be confirmed.
The existence of stress buffering function of the cobalt based binder
phase (Co phase) facilitates the improvement of the resistance against
cohesive failure of the film. The coherent and/or semi-coherent epitaxial growth of the film/WC or film/Co surfaces facilitates the improvement of the adhesion strength and the resistance against cohesive
failure of the film. In this work, from the aspects of (1) residual stress
change in cobalt phase of the substrate in pre- and as-deposition states,
(2) the Co phase crystal structure determined by X-ray diffraction
(XRD) and high resolution transmission electron microscopy (HRTEM)
Corresponding author.
E-mail address: zhangli@csu.edu.cn (L. Zhang).
https://doi.org/10.1016/j.ijrmhm.2018.08.002
Received 20 March 2018; Received in revised form 31 May 2018; Accepted 6 August 2018
Available online 07 August 2018
0263-4368/ © 2018 Elsevier Ltd. All rights reserved.
International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
strength expressed by critical loads LC1 and LC2 for the two coatings are
achieved, respectively, using P12Co as the substrate. Where LC1 corresponds to the first stripping of the film or the initial exposure of substrate and LC2 corresponds to the complete removal of the film from the
substrate [14, 15].
2.2. Characterization
Phase identification and residual stress analysis were conducted on
Rigaku D/Max 2500 X-ray diffractometer and Bruker D8 Discover X-ray
diffractometer, respectively. MDI Jade 6.5 analysis software developed
by Materials Data Inc. was used to analyze the XRD spectrum for the
phase identification. Focused ion beam (FIB) technology (Helios
Nanolab 600i electron/ion double beam microscope system) was used
for the preparation of the TEM observation specimen. Field emission
TEM with spherical aberration correction (Titan G2 60–300) was used
for the microstructure observation. DigitalMicrograph software provided by Gatan Inc. was used to analyze the HRTEM images.
3. Results and discussion
3.1. Coating related phase structure
XRD patterns of the the film/P12Co systems and the analysis results
are shown in Fig. 2. From Fig. 2 and the formerly reported XPS analysis
results [12], it comes to a conclusion that both AlCrN and AlCrSiWN/
AlCrN films are crystallized into an AlN based cubic structure which
belongs to Fm-3 m (225) space group. A strong (111) preferred orientation (20% intensity in card standard value) of the AlN based phase
can be identified. Notably, in addition to a hexagonal structured WC, a
mixture of hcp and fcc structured Co phase is identified for the two
coated alloys.
Although a mixture structure of hcp and fcc in the Co phase of cemented carbides has been reported [16–19], in our case, the coexistence mechanism of hcp and fcc differs. It is well established that
cobalt undergoes a phase transformation from hcp into a fcc structure at
450 °C [20]. Although it is possible for the transformation of fcc-Co to
hcp-Co through the grinding and polishing treatment [21, 22] before
the film deposition, during the PVD deposition process carried out at
500–600 °C for 90–120 min, the reverse transformation of hcp-Co to fccCo can take place [20, 23].
To support the above proposed coexistence mechanism of hcp and
fcc, XRD analysis of the substrate (P12Co alloy) with different surface
states was carried out, shown in Fig. 3. State A corresponds to a state of
Fig. 1. SEM images of sinter skin (a) and polished section (b) of plate-like
grained WC–12Co–0.9Cr3C2–0.4VC–0.05La2O3 (P12Co for short) substrate
[11].
and (3) the film/Co interface microstructure, the stress buffering
function of the Co phase is revealed. From the observed coherent and
semi-coherent epitaxial growth of AlCrN on WC surfaces, the superiority of plate-grained WC–Co substrate is confirmed.
2. Experimental details
2.1. Coated alloy preparation
A fine plate-like grained WC–12Co–0.9Cr3C2–0.4VC–0.05La2O3
alloy (in mass fraction, %, P12Co for short) with a homogenous microstructure shown in Fig. 1 was used as the substrate. The alloy is
characterized as follows: transverse rupture strength 3679 MPa (B type
specimen), Vickers hardness 1680 kgf/mm2 HV30, Palmqvist toughness
12.1 MPa∙m0.5, texture coefficient TC 2.73 of WC(0001) for the sinter
skin and average TC 1.36 of WC(0001) and WC(0002) for the polished
section. All the details were presented in our previous work [11]. Prior
to the deposition, the specimens were gradually ground and quickly
polished to mirror finish surfaces, so as to attain a straight substrate and
film interface which facilitates the HRTEM investigation.
Single
layer
Al0.66Cr0.34N
(AlCrN)
and
two
layers
Al0.65Cr0.30Si0.04W0.01N/ Al0.66Cr0.34N (contact layer) (AlCrSiWN/
AlCrN) approximate 1.6 μm in total thickness for each were separately
deposited on the substrate. The deposition was conducted in an
Oerlikon Balzers' Rapid Coating System based on cathodic arc evaporation technology. For the film deposition, the vacuum chamber was
evacuated to a pressure of 1 × 10–3 Pa, and the substrates were heated
to a temperature of 500 to 600 °C to outgas the surfaces. Then the
substrates were sputter-etched at an Ar pressure of 1 Pa, − 100 V DC
substrate bias for 30 min. Finally, a bias voltage of – 40 to – 70 V was
applied to the substrate to start coating deposition, which was carried
out in N2 atmosphere (purity > 99.99%) at 1 Pa. The deposition temperature and residence time were 500 to 600 °C and 90 to 120 min,
respectively. All the details were presented in our previous work [12,
13].
Coating adhesions were evaluated using a CSM Revetest Scratch
Tester with a normal force range of 1 N to 200 N. The results shown
that > 100 N resistance against cohesive failure strength and adhesion
Fig. 2. XRD patterns film/P12Co systems and the analysis results. Where marks
“●” and “♥” specifically emphasize the existence of hcp-Co and fcc-Co, respectively. Mark
emphasizes a strong (111) preferred orientation of AlN
based phase. The number expressed as a percentage represents the intensity in
card standard value.
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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
the use of stainless steel container for wet-milling, contamination of Fe
and Cr in the ready to press powder is inevitable [27]. Secondly, the
composition of the investigated alloy is characterized with a reasonable
high Cr3C2 content, i.e., 0.9 wt%. Thirdly, most of the Cr3C2 additive is
dissolved in Co [27]. The interaction of Fe, Cr, V and W atoms and the
effect of additional mechanical force affect the martensitic transformation of Co.
3.2. Coating related residual stress change in substrate
To investigate the aforementioned reason for the coexistence mechanism of hcp and fcc and the assumed stress buffering of the ductile
binder phase to the film [10], residual stresses at room temperature in
WC phase of the substrate in states of pre- and as-deposition were
measured. Here, pre-deposition state corresponds to a grinding and
polishing state before the film deposition and as-deposition state to a
coating state.
Diffraction measurements of residual stress can be performed by
both the X-rays [28] and the neutrons [29] which are characterized
with higher penetration depth compared with X-rays and hence the
effect of the surface state on the measurement results can be avoided.
Taking the superiority of the sensitive characteristic to the surface state,
XRD technology was chosen for the surface stress analysis using sin2ψ
method [28, 30]. Parameters used for the residual stress measurement
and calculation of WC phase are listed in Table 1.
According to the principal stress distributions, at room temperature,
there are both tension stress and compression stress in WC phase and Co
phase of cemented carbides; nevertheless, the mean stresses in WC
phase and Co phase are in states of compression and tension, respectively [32]. Direct measurement of the stress in the Co phase is not
possible because the synergetic dissolution of W, C and other alloying
elements in the Co lattice during the sintering process expands the cell
parameter by an unknown amount [27] and there is a great variable in
the hcp to fcc ratio [33]. The mean stress in the Co phase (σCo) is usually
determined on the basis of the mean stress in WC (σWC) and the force
balance relationship [31, 34, 35] according to:
Fig. 3. XRD patterns P12Co alloy with different surface states and the analysis
results. Where marks “●” and “♥” specifically emphasize the existence of hcpCo and fcc-Co, respectively. A: heavily ground and polished to mirror finish
surface; B: state A + treated at 500 °C for 90 min in 4 N Ar atmosphere; C: sinter
skin in as-sintered state.
heavily ground and polished to mirror finish surface; state B corresponds to a state of heavily ground and polished to mirror finish surface
+ treated at 500 °C for 90 min in high purity (4 N) Ar atmosphere; state
C corresponds to a state of sinter skin in as-sintered state. The heavily
grinding and polishing treatment aims to create a high ratio of fcc-Co to
hcp-Co phase transformation. A WC + fcc-Co two phase structure is
observed on the sinter skin. The higher peak strength of fcc-Co detected
on the sinter skin (state C) than that on the cross section (state B) is
caused by the slight enrichment of Co on the sinter skin. Phenomena of
the enrichment of Co on the sinter skins are very common and were
reported by different authors [24, 25]. From Fig. 3, we can draw conclusions that heavily ground and polished treatment has resulted in the
fcc-Co to hcp-Co phase transformation, and heating treatment has resulted in a completely inverse phase transformation, i.e., a single fcc-Co
structure.
The results by Tian et al. [26], shown in Fig. 4, indicate that alloying
Co with Cr, Ru and Rh promotes the hcp phase formation, while Fe, Ni
and Pd favor the fcc phase. The effect of Mo and W on the phase
transition differs from the other elements. For concentrations below
10%, the intrinsic stacking fault energy (ISF) is lower than that for pure
fcc Co, whereas above 10% the situation reverses.
With the help of the results by Tian et al. [26], it is easy to understand the results shown in Fig. 3. Firstly, we have reported that due to
f σCo + (1–f ) σWC = 0
(1)
where f is the volume fraction of the Co phase.
For a straight WC–12Co alloy, 12% of the mass fraction of pure Co
corresponds to the volume fraction of 19.5%. For two-phase structured
P12Co alloy, there are a certain number of W, C, Cr and V atoms dissolved in the Co phase. According to our former report [27], for VC and
Cr3C2 combined additives with a reasonable relative ratio and addition
amount, < 40 wt% of the V additive and > 95 wt% of the Cr additive
are dissolved in the Co phase, respectively. Accordingly, a 21% volume
fraction of the Co phase is estimated for our substrate alloy. Neglecting
the influence of the film/substrate interface atomic layer, the residual
stresses in the Co phase is estimated according to Eq. (1). Table 2 shows
the room temperature residual stresses in the WC phase and Co phase
near the surface area of the substrate in a pre-deposition state, and near
the substrate-film interface area. Taking the pre-deposition state as a
reference, after the deposition of AlCrN and AlCrSiWN/AlCrN films,
Table 1
Parameters used for residual stress measurement and calculation of WC phase.
Scan range
122.5°–125.5°
HKL
(201)
Ψ
0°, 15°, 20°, 25°, 30°, 35°,
40°, 45°
Co
35 kV
697 GPa [31]
Side inclination
Step
0.04°
Count time
Current
ν
Measurement time
12 s
40 mA
0.24 [31]
122 min
Radiation
Voltage
E
Measurement
Method
Fig. 4. Variation of the intrinsic stacking fault energy (ISF) of Co-based alloys
as a function of composition. Values are given with respect of the value in pure
fcc Co. Inset shows the concentration interval between 2 and 12 at.% [26].
Remarks: Ψ: tilt angles; E: Young's modulus; ν: Poisson's ratio.
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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
amplitude of variation, the tendency for the stress increment in WC and
Co phases near the interface caused by the film deposition can be established.
It is reported [32, 35] that residual stresses in WC phase and Co
phase decrease linearly with increasing temperature. Accordingly,
during the PVD deposition process carried out at 500–600 °C for
90–120 min, stress restoration can take place. As a result, stress/strain
energy absorption of the substrate from the film during the film growth
process [36, 37] and the pronounced differences in thermal expansion
coefficients among WC, Co and AlN based phases are predominately
responsible for the high stress build-up in the WC and Co phases, which
is different from the mechanical and thermal mixed state of residual
stresses for the pre-deposition state. Accordingly, it is safe to draw a
conclusion that the substantially high stress build-up represents a
strong ability of energy absorption of the substrate, which facilitates the
release of stress in the film and the improvement of the film toughness
and the ability against film cohesion failure. Linking the two facts together, i.e., the existence of hcp-Co and the substantially high stress in
the Co phase, it is reasonable to draw a conclusion that the high stress
in the Co phase is responsible for the existence of hcp-Co and the phase
transformation facilities the energy absorption from the film.
Table 2
Summary of residual stresses in WC phase and Co based binder phase (Co
phase) of P12Co substrate in pre- and as-deposition states, at room temperature,
MPa.
State
Pre-deposition
Coated with
AlCrN
Coated with
AlCrSiWN/
AlCrN
WC phase
Co phase
Residual
stress
Measurement
error
Residual
stress
Estimated
error
– 1135
– 1869
± 57
± 57
4270
7032
± 214
± 213
– 1669
± 67
6277
± 253
tensile stresses in the Co phase adjacent to the interface of the AlCrN/
substrate increase by 65% and 47%, respectively. Stress in the substrate
is closely related to the stress in the coating. One way for effectively
reducing the stress in the coating is to use a multi-layer deposition
technique. Accordingly, the two-layer AlCrSiWN/AlCrN film corresponds to a lower tensile stress increment in the Co phase.
The calculated residual stresses of Co phase shown in Table 2 are
much higher than the assumed yield limit 1050 MPa of pure Co [32]. It
is worth mentioning that in cemented carbide, the Co phase is existed in
a form of solid solution, instead of pure Co. Therefore, the yield limit Co
phase must be higher than that of pure Co.
Due to the existence of the adjacent AlCrN film or the deformation
caused by grinding and polishing, using Eq. (1) can result in big error in
the calculation of the stress in Co phase. Nevertheless, due to the big
3.3. HRTEM phase interface structures of film/substrate system
TEM images of AlCrSiWN/AlCrN/P12Co (film/substrate) system are
shown in Fig. 5. Columnar crystal, layer-by-layer growth mode and AlN
based fcc structure can be identified in both AlCrN and AlCrSiWN
layers, with a preferred orientation along the vertical direction of the
film/substrate interface. Significant grain refinement caused by the
Fig. 5. TEM images of film/substrate system (a), AlCrN layer (b), AlCrSiWN layer (c) and AlCrSiWN/AlCrN interface zone (d). The corresponding SAED pattern is
shown in an embedded way.
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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
Fig. 6. TEM images of film/substrate system, showing the observation zones of AlCrN/WC interfaces 1# (a) and 2# (b) and AlCrN/Co interface 3# (c). AlCrN/WC
interface 1# is also shown in (c).
Fig. 7. HRTEM image comprising AlCrN/WC interface 1# (a) and the corresponding FFT patterns of AlCrN (b) and WC (c).
Fig. 8. HRTEM image comprising AlCrN/WC interface 2# (a) and the corresponding FFT patterns of AlCrN (b) and WC (c).
interactions of Al, Cr, Si and W atoms can be identified. Selected area
electron diffraction (SAED) patterns (selected area diameter
Φ ≈ 200 nm) of the two layers show the characteristic of polycrystalline diffraction ring. Nevertheless, due to the larger grain size, incompleteness of the polycrystalline ring is identified for the AlCrN layer
(Fig. 5b). Additionally, diffuse halo is identified for the AlCrSiWN layer
(Fig. 5c), which is caused by the presence of nano-scale amorphous
phase [38–40]. Owing to a coherent epitaxial growth relationship of
AlCrSiWN on AlCrN lattice, it is hard to locate exactly the interface.
TEM images of the selected observation zones of AlCrN/WC interfaces 1# and 2# and AlCrN/Co interface 3# are shown in Fig. 6.
HRTEM images comprising the selected AlCrN/WC interfaces 1#
and 2# and the corresponding fast Fourier transform (FFT) patterns of
AlCrN and WC are shown in Fig. 7 and Fig. 8, respectively. The FFT
patterns of AlCrN and WC are corresponding to zone A and zone B in
Fig. 7a and Fig. 8a, respectively. After FFT and inverse FFT (IFFT) of the
marked transition zones C shown in Fig. 7a and Fig. 8a, clear AlTiN/WC
interface microstructures are attained, shown in Fig. 9. Both α and Δα
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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
Fig. 9. IFFT images comprising AlCrN/WC interfaces corresponding to the marked zones C shown in Fig. 7a and Fig. 8a, respectively, showing epitaxial growth
characteristics of coherent interface 1# (a) and semi-coherent interface 2# (b).
overlaying of the TEM interface image and the IFFE image.
As shown in Fig. 9, the following orientation relationship is established, i.e., (111)AlCrN // (0001)WC for interfaces 1# and 2#. The same
epitaxial growth characteristic of AlCrN grain on WC surface is observed, i.e., in the direction of [211]AlCrN which is parallel to the (111)
AlCrN plane, a preferred orientation of AlN based phase established by
the XRD analysis shown in Fig. 2. As shown in Fig. 9, there is a certain
angle of 16° between the AlCrN epitaxial lattice and the WC lattice. In
this case, a lower interface misfit can be attained, as compared with the
situation of no angle deflection. This is a result of self-consistent for
mutually matching [41, 42]. For the estimation of the lattice misfit |δ|,
different Eqs. [41] are used, depending on the actual interface arrangements:
Table 3
Comparison of lattice misfits (|δ|) between AlCrN and WC phase at the interfaces.
No.
Orientation
relationship
Interplanar spacing (nm) and
tilt angle
|δ| (%)
Interface type
1#
α = 0°, Δα = 16°
4.0
Coherent
α = 21°, Δα = 16°
6.3
Semi-coherent
1A#
(111)AlCrN //
(0001)WC
[011]AlCrN //
[1210]WC
(111)AlCrN //
(0001)WC
[011]AlCrN //
[1210]WC
Same as 1#
8.1
Semi-coherent
2A#
Same as 2#
d(111)AlCrN = 0.2385
d(1010)WC = 0.2578
d(111)AlCrN = 0.2376
d(1010)WC = 0.2576
8.4
Semi-coherent
2#
|δ| =
Note: |δ| of interfaces 1# and 2# are calculated according to Eq. (2); |δ| of the
assumed interfaces 1A# and 2A# are calculated according to Eq. (3).
cos α
−1
cos(|α − Δα|)
(2)
|δ| =
|ds − dc |
dc
(3)
|δ| =
|ds − 2dc |
2dc
(4)
where dc and ds refer to the interplanar spacing of the film and the
substrate phase, respectively. When the lattice of the film is tilted with a
certain angle Δα relative to the substrate plane, Eq. (2) is used. When
angle deflection Δα is negligible and dc and ds is close, Eq. (3) is used.
When Δα is negligible and there is an approximate two times relationship between ds and dc, Eq. (4) is used. The calculation results of
the misfits are shown in Table 3. |δ| of 1A# and 2A# are calculated
under the assumption that AlCrN grows without angle deflection.
As shown in Table 3, the |δ| of interface 1# is 4.0% (coherent interface), which is only half the |δ| of the assumed interface 1A# (8.1%,
semi-coherent interface). Compared with the assumed interface 2A#
(|δ| = 8.4%), the lattice misfit of interface 2# (|δ| = 6.3%) is reduced
by up to 33%. Because of the relatively large lattice distortion, edge
dislocations are observed in the interface 2# comprising zone in the side
of AlCrN, shown in Fig. 9b.
HRTEM images comprising AlCrN/Co interface 3# and the corresponding FFT patterns of AlCrN and Co are shown in Fig. 10. The FFT
patterns of AlCrN and Co are corresponding to zone A and zone B in
Fig. 10a, respectively. The interface transition zone C is specifically
demonstrated with quite large magnification, shown in an embedded
way in Fig. 10a. Notably, at least two sets of FFT patterns for AlCrN can
be identified from Fig. 10b, which is caused by the nucleation stimulation effect of the Co phase. It is confirmed that the Co phase near the
AlCrN/Co interface has a hcp single-phase structure (Fig. 10c). Unlike
Fig. 10. HRTEM image comprising AlCrN/Co interface 3# (a) and FFT patterns
of AlCrN (b) and Co (c). The interface transition zone C is specifically demonstrated with quite large magnification, shown in an embedded way.
are marked in Fig. 9, where α refers to the inclination angle of the
substrate crystallographic plane relative to the interface, and Δα refers
to the intersection angle between the crystallographic planes of film
and the substrate. The inclination angle α is determined through the
parallel shift of the observed AlCrN/WC interface and the image
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International Journal of Refractory Metals & Hard Materials 77 (2018) 97–104
L. Zhang et al.
the Co phase can be highlighted through the formation of atomic
disorder zone adjacent to the AlCrN/Co interface and the targeted
regulation and control of the stacking fault energy of the Co phase.
As a result, the resistance against cohesive failure of the films can be
synchronously and effectively strengthened.
(4) AlCrN grain growth exhibits a strong ability of self-consistent for
mutually matching with the WC phase. Plate-grained WC–Co substrate facilitates the formation of coherent and semi-coherent WC/
AlCrN interfaces, which can effectively strengthen the film adhesion and the resistance against cohesive failure.
the AlCrN/WC interface, a direct epitaxial growth relationship is not
existed for the AlCrN/Co interface. Nevertheless, it is established that
AlCrN film grows continuously on the Co phase, accompanied by an
atomic disorder zone approximately 2 nm in wide in the Co phase side
of AlCrN/Co interface shown in Fig. 10a. The substantially high stress
state near the interface and the phase transformation are thought to be
responsible for the formation of the atomic disorder zone.
By measuring the interplanar spacing values of zone A and zone B
based on the FFT pattern shown in Fig. 10b and c, we attain the following data, i.e., d(111)AlCrN of 0.2379 nm and d(0001)Co of 0.4071 nm.
From Eq. (4), we get the |δ| of 14.3% between (111)AlCrN and (0001)Co,
a potential semi-coherent relationship. Clearly, the inexistence of the
semi-coherent AlCrN/Co interface and the existence of the atomic disorder transition zone are strongly related to the substantially high stress
build-up in the Co phase and the stress/strain-induced gradual process
of fcc to hcp-Co phase transformation. Admittedly, the formation of the
atomic disorder transition zone facilitates the relaxing of internal strain
energy caused by the lattice misfit.
The formation of atomic disorder zone adjacent to the AlCrN/Co
interface, the single hcp-Co structure near the interface indicated by
HRTEM analysis, the mixture structure of fcc-Co and hcp-Co established
by XRD analysis, as well as the stress/strain build-up in the Co phase
can provide a strong support for the film deposition induced fcc to hcpCo phase transformation. Not only in the process of film deposition, but
also in the service process of the cutting tool, the step by step fcc to hcpCo phase transformation facilitates the energy absorption and hence the
stress buffering, which facilitates the improvement of the tool life and
its stability.
Based on the result from first principles study of the stacking fault
energies for fcc Co-based binary alloys, Tian et al. [26] reported that
alloying Co with Cr, Ru, and Rh promotes the hcp phase formation.
Marx et al. [43] reported that strain induced fcc to hcp-Co phase
transformation can enhance the ductility and toughness of the substrate. From this aspect, it can be deduced that Cr3C2 additive in cemented carbide substrate can delay the crack onset of the film/substrate
system. Tolédano et al. [44] reported that a stacking disorder is always
present in the hcp-Co phase even during the early stage of the transformation. Owing to the high concentration of stacking faults, which is
consistent with the low stacking-fault energy of cobalt enhanced by
Cr3C2 additive, the lattice fringe image of hcp-Co shown in Fig. 10a is
quite similar to the reports by Waitz and Liu et al. [45, 46].
Acknowledgments
The National Natural Science Foundation of China (grant no.
51574292), the Open-End Fund for the Valuable and Precision
Instruments of Central South University (grant no. 2017gxjj014) and
the foundations of the State Key Laboratory of Powder Metallurgy
(grant no. 2017zzkt21) financially supported this project.
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103
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