Growth behaviors and microstructural characteristics of AlCrTiSiN [622435]

Applied Surface Science
Manuscript Draft

Manuscript Number: APSUSC -D-19-01951

Title: Growth behaviors and microstructural characteristics of AlCrTiSiN
high-entropy alloy nitri de coatings consisting of single -phase solid
solution

Article Type: Full Length Article

Keywords: High -entropy alloy nitride coating; AlCrTiSiN; Single -phase
solid solution; Interfacial structure; Growth behavior.

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Highlights:
(1) AlCrTiSiN HEA coating s consist ing of single -phase fcc -AlCrTiSi N solid solution
were designed and prepared ;
(2) The Cr etching induced an amorphous loose layer on the substrate;
(3) High substrate bias voltage promoted the transformation of amorphous phase to
(Cr,Fe)N solid solution;
(4) Fine microstructures at atomic scale were studied;
(5) Growth behavior of coatings is controlled by the surface energy and then by the
strain energy.
*Highlights (for review)

Growth behaviors and microstructural characteristic s of AlCrTiSiN high-entropy alloy
nitride coating s consisting of single -phase solid solution

Wanglin Chen a, Bingxin Li a, Chengyong Wang a, An Yan a, Yang Deng a, Hui Xiao b*,
Daoda Zhang c, Xianna Meng d*
a School of Electromechanical Engineering, Guangdong University of Technology,
Guangzhou City, Guangdong Province 510006 , PR China.
b The State Key Laboratory of Advanced Design and Manufacturing for Vehicle Body,
Hunan University, Changsha 410082, PR China .
c Jiangxi Mechanical Scientific Institute, Nanchang City, Jiangxi Province 330000, PR
China.
d School of Mechanical and Automotive Engineering, South China University of
Technology , Guangzhou City, Guangdong Province 510006 , PR China.
*Corresponding author: [anonimizat] ; [anonimizat] .

Abstract :
AlCrTiSiN high -entropy alloy (HEA) nitrid e coatings, consisting of single -phase
fcc-AlCrTiSiN solid solution, were designed and prepared . The formation of
single -phase fcc -AlCrTiSiN solid solution is due to high mixing entropy, low mixing
enthalpy and low Al concentration . The Cr ion etching during the pretreatment
process induces to form a thin amorphous loose layer (LL) on the substrate . Phase
transformation of amorphous phase to the (Cr,Fe)N solid solution occurring in the LL *Manuscript
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ensures high adhesion strength. Moreover, the M 23C6 carbides within the steel
substrate act as nucleation sites for (Cr,Fe)N phase of the LL , resulting in an increase
in adhesion strength. At initial stage of coating deposition, the growth behavior of the
HEA nitride coatings is mainly controlled by the surface energy, producing a (200)
preferred orientation. Later, the growth behavior is mainly controlled by the strain
energy. To lower the strain energy, some neighbor ing columnar grains keep
epitaxial -growth characteristic s, accompanied by the formation of high -density
dislocations and dislocation loops along the {111} fcc crystal planes which further
decreases the strain energy of coatings. Furthermore , the formation of high -density
fcc-AlCrTiSiN nano -twins and the hcp -AlN precipitation at nano -twin boundaries
produces a positive effect for reducing the strain energy. Nano -multilayered structure,
the (111) preferred orientation and high -density nano -twins make the HEA nitride
coating show simultaneously high hardness and fracture toughness.

Key words : High -entropy alloy nitride coating; AlCrTiSiN; Single -phase solid
solution; Interfacial structure ; Grow th behavior.

1 Introduction
High-speed machining has been carried out for enhancing production efficiency ,
dimensional accuracy and surface finishing quality of machined work -pieces due to
high removal rates and low cutting forces [1-2]. However, high cutting speed
inevitably produces high cutting temperature up to 1000 °C and causes severe wear of
cutting tools . Therefore , high -speed cutting puts forward s the higher demands for
cutting performance and durability of cutting tools . Physical vapor deposition (PVD)
hard coatings have been widely used for protecti on of cutting tools due to excellent
wear resistance, thermal stability and oxidation resistance [3-6]. So far , binary, ternary
and quaternary medium -entropy alloy (MEA) nitride coatings have been designed and
developed [7-9].
Recently, a few investigations have been paid to the multi -component
high-entropy alloy (HEA) nitride coatings [10-13]. These HEA nitride coatings
usually show excellent wear resistance, thermal stability and oxidation resistance
[14-15]. Moreover, o ur studies found that under high -speed dry milling of DIN 1.2311
hardened steel [5] and 40CrNiMoA steel [6], the AlCrTiSiN HEA coated cutting tools
show ed the longer service life when compared with the AlCrN, TiSiN , TiCrSiN and
AlCrSiN coated cutting tools . However, compared with the MEA nitride coatings , the
studies of HEA nitride coatings for cutting tools are not systematic. The design ,
preparation and study of HEA nitride coatings are issues of current interest for
exploring advanced -performance coated cutting tools.
It has been demonstrated, b y adjustment of chemical composition s of targets, the

HEA nitride coatings can be usually fabricated by PVD methods in following two
ways: direct deposition with a multicomponent HEA target, which has a good control
in coating stoichiometry; and the co-deposition by multiple metal targets, which has a
wide range in chemical compositions [15]. Because of difficulty in multicomponent
HEA -target preparation and narrow chemical compositional range in HEA nitride
coatings by the first way, the second way for deposition of HEA nitride coatings is
preferentially accepted in industrial applications . Up to now, some AlCrTiSiN HEA
coatings with different chemical composition and microstructures have been designed
and deposited by the co-deposition method simultaneous ly with pure Cr and
Al60Ti30Si10 targets [16], or Al70Cr30 and Al 60Ti33Si7 targets [10,17], or Ti70Cr30,
Al12Si88 and pure Cr targets [18], or Ti 64Al36 and Cr 90Si10 targets [19]. However, these
HEA nitride coatings have a common point in microstructur al characteristic , namely
consist ing of nano -sized grains embedded into Si 3N4 amorphous matrix [10, 16-19].
Such nanocomposite structure cannot endure high temperature up to above 900 °C,
otherwise the recrystallization happen s, leading to notable loss es in hardness and
adhesion strength [10,17]. Therefore, the higher thermal stability of AlCrTiSiN HEA
nitride coatings should be designed and fabricated. According to HEA theor ies [14]
and thermodynamics theor ies, increasing mixing e ntropy and reducing the Gibbs free
energy of PVD coatings are beneficial for obtaining single -phase solid solution with
excellent thermal stability . However, such AlCrTiSiN HEA nitride coatings,
consisting of single -phase fcc solid solution , have been not reported yet .
In addition, HEA nitride coatings have commonly five or more c omponent s,

which bring great difficulties for microstructural investigation s. In past years, some
traditional characterization methods, including X-ray diffraction, scanning electron
microscopy , X-ray photoelectron spectroscopy and mass /diffraction -contrast
transmission electron microscopy (TEM) , have been widely carried out to study the
microstructures of HEA nitride coatings. By these traditional methods with large
uncertainties, it is difficult for obtaining fine microstructures of PVD coatings ,
including orientation relationship, interfacial structur al characteristics , substructure
characteristics and chemical compositional distribution s. However, fine
microstructur al characteristics could provide some crucia l information for
understand ing the growth behavior s and the essence s of both high hardness and
adhesion strength. Therefore, more characterizations, by an analytical high-resolution
TEM (HRTEM) equipped with a high-angle -annular -dark-field imaging in scanning
transmission electron microscopy (HAADF -STEM) and an X -ray energy dispersive
spectrometer (EDS), are needed to show the fine microstructures , interfac ial structures
and chemical compositions . Furthermore, nano -hardness and adhesion strength should
be characterized to understand the relationships between structural characteristic and
property.
In this work, multi-component AlCrTiSiN HEA nitride coatings , consisting of
single -phase solid solution , were designed and fabricated by adjustments of target
compositions and the substrate bias voltage . An a nalytical high-resolution TEM
(HRTEM) was carried out to study fine microstructural details of the AlCrTiSiN HEA
nitride coatings , and further to understand the growth behaviors and the essence s of

both high hardness and adhesion strength.
2 Material and methods
2.1 PVD sample preparation
Some commercial M2 high -speed steel plates ( 15 mm×15 mm×3mm ) were
mechanically polished to an average roughness (Ra) of 0.05 μm, and then
ultrasonically cleaned for 30 min at room temperature. Subsequently , the cleaned
samples were installed on homemade Ar-assisted cathodic vacuum arc equipment
with four targets holders, and each hold er contain ed one alloy target (Φ160 mm) .
Before the coating deposition, the chamber was pumped down to a low pressure of 1
×10-4 Pa, and meanwhile the substrate was heated to 420 °C. Then, the samples were
etched by Cr ions for 8 min under a bias voltage of -800 V and an etching current of
160 A. The workpiece turntable was rotated at a speed of 3.5 rpm . To increase the
adhesion strength, a CrN buffer layer (BL) was first deposited by p ure Cr target , and
then a nano -multilayered CrN/ AlCrN transition layer (TL) was deposited by pure Cr
and Al 30Cr70 targets . Finally, a AlCrTiSiN HEA nitride layer (HEA -NL) was
co-deposited simultaneously by Al30Cr70 and Ti 82Si18 targets . The deposition
parameters of the AlCrTiSiN coatings are shown in Table 1. For convenience, the
AlCrTiSiN HEA coatings deposited under the substrate bias voltages of -50 V , -80 V
and -110 V were labelled as HEA -50, HEA -80 and HEA -110 samples, respectively.
2.2 Characterizations
Phase structure of the coated samples was tested by a grazing incidence X-ray
diffract ometer (XRD , Bruker D8 advance) at an incident angle of 2 °. The scanning

angle ranged from 30 ° 100 °. The chemical compositions of the Al,Cr, Ti, Si and N
elements were measured by an X -ray photoelectron spectroscopy (XPS , Escalab
250Xi) operating Al K α with 15 kV and 25 W. Fine c ross-sectional microstructures
were characterized by an analytical high-resolution transmission electron microscopy
(HRTEM , FEI Tecnai F20) , which was equipped with an X -ray energy dispersive
spectrometer (EDS, EDAX) and a high-angle -annular -dark-field imaging in scanning
transmission electron microscopy (HAADF -STEM) . The samples for TEM
characterization were prepared by a dual -beam focused -ion beam (FIB, FEI Helios
Nanolab 600 ).
The adhesion strength was tested by a Rockwell -C indention tester using a load
of 150 kg f and a scratch test er (CSM Revetest) using a conical diamond tip of 0.2 mm
radius under a load of 0 120 N . Both indentation impress es and scratch tracks were
observed by an optical microscopy (OM , Leica DM2700M ). Hardness (H) and elastic
modulus (E) were tested by a nanoindent er (MTS XP ) equipped with a
Berkovich -type diamond indenter.
3. Results
3.1 XRD analysis
Fig.1 shows typical GIXRD patterns of the HEA -50, HEA -80 and HEA -110
samples. Four diffraction peaks are observed for each XRD pattern, which are
recognized as the (111) fcc, (200) fcc, (220) fcc and (311) fcc crystal planes of a
single -phase fcc solid solution. This indicates that these HEA coatings basically
consist of single -phase fcc-AlCrTiSiN solid solution . This result is notably different

from the previous reports that the AlCrTiSiN coating s consisted of nano -sized
fcc-AlCrTiN and hcp -AlN embedded into amorphous Si 3N4 matrix [10, 17, 20]. It has
been demonstrated, when the Al/(Al+Me) (Me=Cr+Ti) ratio exceeds 0.67 [21] or 0.7
[22], the hcp-AlN will exist . The formation of single -phase solid solution is associated
with mixing entropy ( ), mixing e nthalpy ( ) and atomic size difference ( ). High
is inclined to obtain single -phase fcc solid solution , and l ow can ensure high
stability of such solid solution , and large atomic size difference ( ) also favor s the
single -phase structure [15]. The can be calculated by the open quantum materials
database (OQMD) [23], and the and can be calculated by following equation s
[14-15] :

(1)

(2)

(3)
where n , , , and are the number of components, the atomic radius, the average
atomic radius and the ato mic fraction of the ith element, respectively. Calculated
values are shown in Table 2. The Al/(Al+Me) ratios of the HEA -50, HEA -80 and
HEA -110 samples are 0.5 3, 0.49 and 0.4 7 respectively , much lower tha n the critical
value (0.67 [21] or 0.7 [22]). Moreover, the HEA -50, HEA -80 and HEA -110 samples
have high , low and large values. Abovementioned causes can be used to
explain why only single -phase fcc solid solution structure occur s in the HEA -50,
HEA -80 and HEA -110 samples .
3.2 HRTEM characterization s of the HEA -80 sample

Because of low adhesion strength of the HEA -50 sample , both HEA -80 and
HEA -110 samples are selected for cross -sectional TEM (HRTEM) characterizations
in more details . As shown in Fig.2a, c ross-sectional microstructure s of a HEA -80
sample consist of an inner BL, an intermediate TL and an outermost HEA -NL,
showing typical columnar characteristic across coating depth. Fig.2b reveals a thin
loose layer (LL) between BL and steel substrate, with about 25nm in thickness.
Careful observation shows that there is irregular interface between LL and TL,
indicating a n imperfect cohesion between them. Line-scan EDS profiles ( Fig.2c)
display that the LL is rich in Fe, Cr and N atoms, and the BL enriche s Cr and N atoms ,
almost without Fe atoms . This indicates that the formation of LL shou ld be related to
the Cr -etching pretreatment under high substrate bias voltage. Both HRTEM and fast
Fourier transformation (FFT) images (Fig. 2d and e) illustrate that the LL contains
many nano -sized particles with less than 6nm in diameter, which are recognized as
fcc-(Cr,Fe)N solid solution. Furthermore, a nanocomposite structure is also observed
around the fcc-(Cr,Fe)N solid solution , which consists of fcc -clusters with several
atomic layers in thickness embedded into amorphous phase (Fig. 2f). Obviously, the
LL is an incomplete recrystallization zone. Fig.2g reveals a HRTEM image of two
neighboring columnar grains (Grain -A and Grain -B) within the BL , with its FFT
image inset. The (200) Grain-A crystal plane of the Grain -A is exact parallel to the
(200) Grain-B of the Grain -B, and they have the same interplanar spacing . The
neighboring columnar grains keep a following orientation relationship:
(200) Grain-A//(200) Grain-B and [001] Grain-A//[011] Grain-B. The (200) fcc crystal has lowest

surface energy than the (220) fcc and (111) fcc crystal planes [24-25]. Such specific
orientation relationship is beneficial for reducing the surface energy of the BL as the
growth behavior is mainly controlled by the surface energy at the initial stage of
PVD -coating deposition . The TL shows a typical nano -multilayered structure, with a
modulation period of about 12nm (Fig.2b). T he EDS analysis has also shown that
there are notable fluctuations in Cr and Al concentrations occurring in neighboring
sublayers of the TL due to the inter -diffusion of Cr and Al atoms occurring already in
the PVD process ( Fig.2c). The HRTEM image (Fig.2h) demonstrates that the TL
consists of alternative rich-Cr (dark contrast) and rich-Al sublayer s (bright contrast),
and two sublayers keep coherent relationship s, showing epitaxial growth
characteristic . Indexed FFT image inserted in Fig. 2h further confirms that both
sublayers are also composed of single -phase fcc-AlCrN solid solution .
Fig.3a shows a selected area electron diffraction pattern (SEAPD) of the
outermost HEA -NL, showing p olycrystalline diffraction rings, which is identified as
single -phase fcc solid solution. Fig.3b shows typical bright field (BF) TEM image of
several columnar grains within the HEA -NL, viewing along [011] fcc direction . It
shows that the columnar grains are also composed of nano -multilayered sublayers
with a modulation period of about 8 nm. The Grain -2, Grain -3, Grain -4 and Grain -5
basically show approximate diffraction contrast images, indicating that these grains
should have approximate orientation relationship. At less 20 BF-TEM images ,
viewing along [011] fcc direction , were carried out to s tatistical ly calculate the width of
columnar grains, and the average width of columns is about 65nm. Line-scan EDS

profiles (Fig.3c) show irregular fluctuations in Al, Cr, Ti, Si and N concentrations
between neighboring sublayers , indicat ing that the inter-diffusion has occur red in Al,
Cr, Ti and Si atoms during the AlCrTiSiN HEA -NL deposition. Careful observation
shows that there a narrow bright -contrast zone occurring in the Grain -4 which
surrounded by the Grain -2 and Grain -3. HRTEM image (Fig. 3d) of the framed zone in
Fig.3b shows that the lattice fringes of such bright zone are unclear . This indicates
that the crystal orientation of such bright zone is far from [011] fcc direction .
High-magnification image (Fig.3e) of the framed zone in Fig.3d shows that the (11一
1)Grain -4, (111一
)Grain -4 and (200) Grain -4 crystal planes of the Grain -4 are parallel to the (11一
1)Grain -2, (11 1一
)Grain -2 and (200) Grain -2 crystal planes of the Grain -2. This indicates that
both Grain -2 and Grain -4 keep coherent interfacial relationship. Moreover, v iewing
along [ 11一
1]fcc, [11 1一
]fcc and [200] fcc directions, inverse fast Fourier transformation
(IFFT) images (Figs. 3f-h) show some dislocation s and dislocation loops (as
highlighted by an e llipse in Fig. 3f) along the {111} fcc and {200} crystal planes.
Similarly, both Grain -2 and Grain -3 also remain such coherent interface. HRTEM
image ( Fig.3i) further confirm s that the columnar Grain -2 also consists of rich -Ti and
rich-AlCr sublayers , which consists of fcc AlCrTiSiN solid solution. The (1 1一
1)rich-Ti,
(111一
)rich-Ti and (200) rich-Ti crystal plane s of rich -Ti sublayer is exactly parallel to the
cognominal -index crystal planes of rich -AlCr sublayer , showing an epitaxial growth
characteristic . The orientation relationship between them is : (11一
1)rich-Ti//(11一
1)rich-AlCr,
(111一
)rich-Ti//(111一
)rich-AlCr, (200) rich-Ti //(200) rich-AlCr and [01 1一
] rich-Ti //[01 1一
] rich-AlCr.
Fig.4a also displays a BF-TEM image of several columns within the HEA -NL

also viewing along [011] fcc direction . As can be seen, the Grain -6, Grain -7 and
Grain -8 have approximate orientation relationship . Moreover, the fcc -AlCrTiSiN
grains have a broad interface and a narrow interface, which are approximate parallel
to the (111) fcc and (200) fcc crystal planes of fcc -AlCrTiSiN grains , respectively . In
addition , high strain fields ar e clearly observed within the se grains , indicating the
existence of high -density dislocation s. Fig.4b shows a HRTEM image of the framed
zone -A in Fig. 4a, and its FFT image demonstrates that the diffraction spots from
Grain -7 and Grain -8 completely overlap . Indexed results confirm that both Grain -7
and Grain -8 are composed of single -phase fcc-AlCrTiSiN solid solution . Fig.4 c shows
the HRTEM image of the framed zone -B in Fig. 4a. It shows that there are gradient
changes in crystal orientation from the Zone-D to the Zone-G. The formation of such
gradient crystal orientation should be closely associates with high -density dislocations
(as shown in Discussion). Similarly, t he (11一
1)Grain -7, (111一
)Grain -7 and ( 200) Grain -7 crystal
planes of the Grain -7 are also exact parallel to the cognominal index planes of the
Grain -8. Fig.4 d shows the HRTEM image of a bright zone ( the framed zone -C in
Fig.4a ). There is a narrow zone adjacent to the Grain -8, with less than 2nm in
thickness , which keeps coherent orientation relationships with the Grain -8, but the
rest of the bright zone within the Grain -7 is far from the [011] fcc direction .
In HAADF -STEM imaging model, the Z -contrast image is proportional to Z1.7-2.0
(Z represents the atomic numbers) [10,26]. Hence the HAADF -STEM imag ing
technique can be used to interpret the difference s in microstructure and chemical
composition. Fig.5a shows a typical HAAD F-STEM image of several columns within

the HEA -NL of a HEA -80 sample , mainly showing two types of Z -contrast images:
bright and gray . Line-scan EDS profiles (Fig.5b) illustrate that the bright zone
enriches Cr but relatively depletes N and Al atoms. HRTEM image of the “I” framed
zone in Fig.5a confirms that the gray zone corresponds to single -phase AlCrTiSiN
solid solution, without exceptions. HRTEM image (Fig.5d) of the “II” framed zone in
Fig.5a reveals a large n umber of nano -twins occurring in the bright zone, with 13
atomic layers in thickness. Its indexed SAEDP (Fig.5e) confirms that the hcp -AlN
phase is formed besides the nano -twins of rich-Cr fcc-AlCrTiSiN solid solution. The
(0001) hcp plane of hcp -AlN is always parallel to the (111) fcc of fcc -AlCrTiSiN solid
solution and its interplanar spacing of the (0002) hcp is basically equal to that of the
(111) fcc plane. Orientation relationship among these phases are: ( 1一
11一
)M//(1一
1
1一
)T//(000 2)hcp and [011] M //[01一
1一
]T//[2一
110] hcp. Further high -magnification image
(Fig.5f) shows that the hcp-AlN clusters with 2 atomic layers in thickness always
precipitate at the twin boundaries (TBs) .
3.3 HRTEM characterization of the HEA -110 sample
Fig.6a and b show typical cross -sectional microstructure s of a HEA -110 sample ,
also consisting of an inner BL, a n intermediate TL, and an outermost HEA -NL. It
shows that there is a dense LL between substrate and BL. Its fine microstructure s at
atomic scale are shown in Fig.6c and d. Steel substrate consists of α-Fe and carbides,
hence two types of interface s should exist : α-Fe/(Cr,Fe) N interface and
carbide/ (Cr,Fe)N interface. Fig.6c shows that t he α-Fe/(Cr,Fe)N interface is a typical
big-angle grain boundary (GB) . Fig.6d shows an irregular interface between carbide

and (Cr,Fe)N solid solution , and its FFT image is shown in Fig.6e. Its simulated
SAEDP (Fig.6f) confirms that the carbide corresponds to the M23C6 (M=Cr, W, Mo
and Fe), and the (020) fcc plane of (Cr,Fe)N solid solution is approximately parallel to
the ( 200)plane of M23C6-typed carbide. The orientation relationship between (Cr,Fe)N
and M23C6-typed carbide is: [001] fcc//[011] M23C6 and (020) fcc//(200) M23C6 with a
deviation of 2.9 ° (Fig.6 e and f). A lot of HRTEM observations further show that the
LL is indeed a complete crystallization zone consist ing of (Cr,Fe)N solid solution.
Fig.7a shows typical SAEDP of the HEA -NL of the HEA -110 sample , and its
indexed results confirm that the HEA -NL is also composed of single -phase
fcc-AlCrTiSiN solid solution. Centered dark field (CDF) image (Fig.7b) , viewing
along (111) fcc direction , shows that the average thickness of columns within the
HEA -110 sample is about 120nm. Fig.7c shows a BF-TEM image of several columns
within the HEA -NL, viewing along [011] fcc direction. It shows that the Grain -A,
Grain -B and Grain -C have approximate orientation relationship, and the high -density
dislocations also occur in the se grains . Both HRTEM image (Fig.7d) of the
rectangular framed zone in Fig.7c and its FFT image (Fig.7d) conform that the
interface between Grain -A and Grain -B is also a coherent interface . Similarly, the
coherent interface between Grain -B and Grain -C is also observed . HRTEM image
(Fig.7e) demonstrates that the Grain -A also consists of alternative rich -Ti and
rich-AlCr sublay ers consisting of single -phase AlCrTiSiN solid solution. Fig.7f shows
a typical HAADF -STEM image of the HEA -NL, also showing two types of
Z-contrast images: bright and black. Fig.7g and h also confirm that the bright zone

contains a large number of nano -twins and the hcp-AlN clusters formed at TBs of the
micro -twins .
3.4 Property measurements
Surface hardness , fracture toughness and adhesion strength are three of key
mechanical properties for PVD coatings , which directly affect the practical utilization
of PVD -coated cutting tools. For protective PVD coatings, high fracture toughness
means high H3/E*2 value [27-28].
Surface hardness (H) , elastic modulus (E) , effective elastic modulus E* (E*=
E/(1-v2), v is the Posson s ratio) and H3/E*2 values are shown in Table 3. It shows that
the H and E values of the HEA -80 sample are about 33 GPa and 239 GPa, the higher
than those (about 32 GPa and 217 GPa) of the HEA -50 sample, but the lower than
those (about 37 GPa and 246 GPa) of the HEA -110 sample. The H3/E*2 values of the
HEA -50, HEA -80 and HEA -110 samples are 0.584 GPa, 0.521 GPa and 0.665 GPa,
respectively.
High substrate bias voltage s not only produce a high residual compressive stress,
but also obtain high dense microstructure s, resulting in a strengthening effect in
hardness but a loss in fracture toughness. So, the HEA -80 sample shows a higher H
but a lower H3/E*2 value when compared with the HEA -50 sample. Moreover, the
formation of nano -twins also produces a positive effect on hardness and fracture
toughness. In addition, XRD results show that the (111) preferred orientation (as
shown in discussion section) occurs in the HEA -80 sample, but almost do not appears
in the HEA -50 sample . The (111) fcc crystal plane is a close -packed plane of fcc -phase,

and the formation of the (111) fcc plane is commonly beneficial for increasing the
fracture toughness of coatings. The HEA -110 sample shows a higher surface hardness
when compared with HEA -80 sample. This indicates that the hardening effect i s
dominant. However, due to more obvious (111) fcc preferred orientation , the HEA -110
sample reveals a higher H3/E*2 value when compared with the HEA -80 sample.
The optical micrographs of Rockwell indentation s of the HE A-50, HEA -80 and
HEA -110 samples are shown in Fig. 8a-c. For the HEA -50 sample , some coarse cracks
around indentation impress and spalling zones are observed at the interfaces between
indentation impress and HEA coating (Fig. 8a). For the HEA -80 sample , a few
spalling zones are observed (Fig.8b). For the HEA -110 sample , the interfaces between
indentation impress and coating are smooth, basically without coating spalling
(Fig. 8c). Further q uantitative measurements of adhesion strength were carried out by
scratch tester, and the tested values are shown in Table 3 , and the micrographs of
scratch impresses are shown in Fig.7d -f. Crack formation stage, peeling off stage and
complete peeling off stage are three stages of t he overall scratch process , and their
critical loads are defined as LC1, LC2 and LC3 [29-31], respectively. The LC2 is
claimed as adhesion strength of coatings [30-32]. The LC2 values of the HEA -50,
HEA -80 and HEA -110 samples are about 44N, 56N and 78N , respectively. The
results of the scratch are well consistent with the results of Rockwell indentation .
According to abovementioned results of mechanical properties, it can be found
that the HEA -110 sample shows the high surface hardness, fracture toughness and
adhesion strength , indicating that its coated cutting tools should show the excellent

cutting performance under high -speed cutting conditions, which will be further
systematically studied in further cutting experiments.
4. Discussion
4.1 Formation mechanism of the LL
The a dhesion strength of the coatings to substrate is closely related to the
interf acial structure between coating and substrate. To enhance the adhesion strength ,
metal ion (Cr or Ti ions) etching, being as in situ cleaning method, has been widely
carried out to remove surface contaminations [33-34]. Moreover, metal ions under
high bias voltage conditions can result in metal ion implantation, which promote s
local epitaxial growth and further ensures high adhesion strength [34].
In this work, the Cr ion etching under a high bias voltage of -800 V was carried
out for etching substrate surface prior to coating deposition. The Cr ion etching not
only removes surface contaminations, but only produce s a strong sputtering effect on
steel substrate (Fig.9a) , inevitably result ing in spattering of Fe atom arriving from
steel substrate . Partial spatter ed Fe atoms are taken away by vacuum pump, but partial
Fe atoms are ionized by electric charge s at near substrate zones, which subsequently
are re-deposited on the samples to form a loose layer (LL) on the substrate under
strong electric -field effect (Fig.9b). However, strong bombardment of Cr ion s under
the high substrate bias voltage i nevitably introduces plenty of defects such voids
within the LL. When the concentration of permanent defects exceeds the equilibrium
concentration of the mobile defects, failure of crystallization growth occurs [35].
Therefore, the LL shows amorphous -structure characteristic (Fig.9b) .

High substrate bias voltage of PVD coatings indicates high h eat accumulation of
coatings , which promotes the mobility of Cr and Fe atoms of the LL. Meanwhile,
plenty of active N atoms diffuse into the LL during the BL deposition . Subsequently ,
both Cr/Fe and N atoms migrate to the preferred sites for crystallization growth , and
further to form (Cr,Fe)N solid solution . The crystallization behavior is more sens itive
to the higher substrate bias voltage. Therefore, u nder relative low substrate bias
voltage ( -80V), the crystallization is not enough, so the LL is composed of a
nanocomposite structure consisting of nano -sized (Cr,Fe)N solid solution and
amorphous phase (Fig.9c). Under high substrate bias voltage ( -110V), the
crystallization is enough, so the LL entirely consists of the (Cr,Fe)N solid solution
(Fig.9d) , almost without amorphous phase. In order to reduce the strain energy , there
is a speci fic orientation relationship between (Cr,Fe)N and carbides. Opposite, the LL
of the HEA -50 sample should contain the higher amorphous -phase concentration
when compared with the HEA -80 sample .
The (Cr,Fe)N solid solution shows the same crystal structure and very
approximate lattice parameters with CrN phase, which act s as nucleation sites for the
crystal grow of the CrN BL, and further ensure high adhesion strength. T herefore, due
to the brittleness nature of amorphous phase and high concentration of the (Cr,Fe)N
solid solution of the LL, the HEA -50 sample shows the lower adhesion strength when
compared with the HEA -80 sample . HRTEM observation shows that there is a
specific orientation relationship between M 23C6 carbides and (Cr,Fe)N solid solution :
[001] fcc//[011] M23C6 and (020) fcc//(200) M23C6 with a deviation of 2.9 °. This indicates

that the carbides can also act at nucleation sites for (Cr,Fe)N crystal growth of the LL,
and further increase the adhesion streng th. Moreover, no amorphous phase is observed
in the LL of the HEA -110 sample . So, the HEA -110 sample shows the higher
adhesion strength when compared with the HEA -80 sample .
4.2 Growth behaviors of AlCrTiSiN H EA-NL
It has been demonstrated, the growth behavior s of the coatings are controlled by
the surface energy and the strain energy . For the fcc-phase s, the surface energies of
(200) fcc, (220) fcc and (111) fcc crystal planes follows: S 200＜S220＜S111, whereas strain
energies are: S 111＜S220＜S200 [24-25]. When the surface energy is dominant , the (200)
preferred orientation is formed, and when the strain energy is dominant, the (111)
preferred orientation is formed. The preferred orientation can be determined by the
texture coefficient (TC) from XRD pattern s, which is defined as [36]:
]/[)/1(/
00
hkl hklhkl hkl
I I nI ITC
(4)
Where , the
hklI and
0
hklI represent the integrated intensities of reflections from the
experimental sample and the standard power sample (Taking CrN powder as refere nce
[37]), respectively; the n is the tot al number of reflection planes. The TCs of {200} fcc
crystal plane of the HEA -50, HEA -80 and HEA -110 samples are about 1.9, 1. 8 and
1.7, and the TCs of {111} fcc of them are about 1. 04, 1.24 and 1.4 4 respectively. The
TCs of {200} fcc of the HEA -50, HEA -80 and HEA -110 samples are the higher than
the critical value (1.0) [37], and the TCs of {111} fcc of HEA -80 and HEA -110 samples
are also the higher the critical value (1.0) . These indicate that almost only (200)
preferred orientation occurs in the HEA -50, but both the (200) and (111) preferred

orientation s occur in the HEA -80 and HEA -110 samples . For the HEA -50 sample, its
growth behaviors are basically controlled by the surface energy. For the HEA -80 and
HEA -110 samples, at initial stage of PVD -coating deposition, the growth behaviors
are mainly controlled by the surface energy, resulting in the formation of the (200)
preferred orientation . However , the formation of the (200) preferred orientation
inevitably results in a notable increase in strain energy of coating systems. Moreover,
the increasing of coating thickness and high substrate bias voltage also increase the
strain energy of coating systems. Later, the growth behaviors of the coatings are
mainly controlled by the strain energy, which promotes the formation of the (111)
preferred orientation . As a result, the HEA -80 and HEA -110 samples simultaneously
contain the (200) and (111) preferred orientation s.
Specific growth processes of the fcc-AlCrTiSiN columnar grains can be depict ed
by following diagrammatic sketch es in Fig.1 0 (Taking the Grain -6, Grain -7 and
Grain -8 in Fig.4a for example in Fig.4 ). As shown in Fig.1 0a, due to the energy and
structures fluctuation s, both Grain -6 and Grain-8 preferential ly nucleate and grow
along the low surface -energy habit plane (about (111) fcc crystal plane), but the crystal
nucleus es of the Grain -7 is relative s low development . To reduce the strain energy of
coating systems, the crystal nucleus es of Grain -7 need to keep coherent relationship
with the Grain -8. We experimentally observe such coherent relationship at the growth
front of the Grain -7 (Fig.4d). Then, the crystal nucleus es gradually develop to become
the small Grain -7 (labelled as J in Fig.1 0b). However, to still remain such low
strain -energy coherent interface , a large number of dislocations and dislocation loops

are formed along the {111} fcc planes , which must induce the lattice distortion and
crystal o rientation deflection . Therefore, we can commonly observe a gradual change
in crystal orientation occurring in the Grain -7 (Fig.4c ). Crystal o rientation deflection
means the existence of lots of sub -interfaces, resulting in a notable increase in the
surface energy of coating sys tems. With the dynamic grain growth (Fig.1 0c), another
zone (K -zone) with gradual change s in crystal orientation is also formed at the growth
front of the Grain -7, but the crystal orientation of previous zone (J -zone) is gradually
close to the [011] fcc crystal orientation of Grain -8, which is also experimentally
observed in Fig.4b) . Later, with the fast growth of the Grain -7, the surface energ ies of
coating systems are rapidly increased, hence the growth of the Grain -7 is gradually
restrict ed. When the surface -energy effect is dominant, such (111) fcc[011] fcc low
strain -energy growth behavior is stopped , and the rest of the Grain -7 is far from
previous growth model (Fig.1 0d). Such growth -structure characteristic has been also
observed carefully in Fig.4 d.
In addition, we always observe that high -density nano -twins are formed within
the rich -Cr fcc -AlCrTiSiN solid solution. Moreover, lots of hcp -AlN clusters
precipitate at TBs (Fig.5d and Fig.7g) . According to the open quantum materials
database (OQMD) [23], the can be calculated by calculated. Here we put
forwards a hypothesis that the Ti, Si and N concentration s are 10 .0 at.%, 3 .0 at.% and
45.0 at.%, and the sum of both Al and Cr is 42.0 at.%, and calculated values are
shown in Fig.1 1. It shows th at with the Cr -concentration increase, the value is
gradually increased, indicating that the improvement of Cr -concentration in the

AlxCr42-xTi10Si3N45 system plays a negative effect on the thermal stability. Calculated
results show that the rich -Cr zone s (bright zone) in Fig.5a and Fig.7f show the lower
stability in comparison to the gray zones. To enhance the stability relating to the strain
energy, lots of nano -twins are formed within the rich -Cr fcc -AlCrTiSiN solid solution,
which can notably lowe r the strain energy of the coating systems. Moreover, hcp -AlN
clusters precipitate at TBs is also inclined to lower the strain energy.
5. Conclusions
From the results obtained, the following can be concluded .
(1) Due to high , low and low Al concentration, the AlCrTiSiN coatings
basically consists of single -phase fcc-AlCrTiSi N solid solution , which shows typical
nano -multilayered structure characteristic.
(2) The Cr ion etching during the pretreatment process induces to form a thin
amorphous loose layer (LL) on the substrate . High substrate bias voltage of
PVD -coating deposition promotes the transformation of amorphous phase to the
(Cr,Fe)N solid solution , resulting in a notable increase in the adhesion strength .
Moreover, the M 23C6 carbides within the steel substrate act as nucleation sites for
(Cr,Fe)N phase of the LL, resulting in a further increase in adhesion strength. Both
(Cr,Fe)N and M 23C6 have a specific orientation relationship: [001] fcc//[011] M23C6 and
(020) fcc//(200) M23C6 with a deviation of 2.9°.
(3) At initial stage of PVD -coating deposition, the g rowth behavior of the
AlCrTiSiN HEA nitride coatings is mainly controlled by the surface energy , resulting
in the formation of (200) preferred orientation. Later, the growth behavior is mainly

dependent on the strain energy. To lower strain energy of coating systems, some
neighbor ing columnar grains always keep coherent interfacial relationship,
accompanied by the formation of high-density dislocation and dislocation loops al ong
the {111} fcc crystal planes. Such (111) fcc[011] fcc low strain -energy growth behavior
promotes the formation of (111) fcc preferred orientation.
(4) High -density nano -twins are formed within the rich -Cr fcc -AlCrTiSiN solid
solution. Moreover, lots of hcp -AlN clusters precipitate at the TBs. The orientation
relationships among them are: ( 1一
11一
)M//(1一
11一
)T//(0002) hcp and [011] M //[01一
1一
]T//[2一
110] hcp.
Such orientation relationships are inclined to lower the strain energy of coating
systems.

Acknowledgments : This work was supported by the National Natural Science
Foundation of China (No. 51641101); Key Research Project of Jiangxi Province
(20161BBE50064 , 20171BBH80011).

References
[1] B. Wang, Z.Q. Liu, Q.B. Yang, Investigations of yield stress, fracture toughness,
and energy distribution in high speed orthogonal cutting, Int J Mach Tools Manuf
73(2013) 1 -8.
[2] R.S. Pawade, S.S. Joshi, P.K. Brahamnkar, Effect of cutting edge geometry and
machining parameters on surface integrity of high -speed turned Inconel 718, Int J
Mach Tool Manuf 48 (2008) 15 -28.

[3] L.E. Gustavsson, H. Barankova, L. Bardos, Some properties of TiN films
produced in hollow cathode and microwave ECR hybrid, Surf Coat Technol 201(2006)
1464 -1468.
[4] D.H. Yu, C.Y . Wang, X.L. Cheng, F.L. Zhang, Optimization of hybri d PVD
process of TiAlN coatings by Taguchi method, Appl Surf Sci 255(2008) 1865 -1869 .
[5] Y .H. Yuan, Z. Qin, D.H. Yu, C.Y . Wang, J.B. Sui, H.S. Lin, Q.M. Wang,
Relationship of microstructure, mechanical properties and hardened steel cutting
performance of TiSiN -based nanocomposite coated tool, J Manuf Process
28(2017)399 -409.
[6] W.L. Chen, D.D. Zhang, D.C. Yao, S.H. Zhang, W.W. Wu, Investigations on
microstructure and mechanical properties of containing -Si coatings, Surf Eng 33(2017)
536-541.
[7] D. Jakub éczyov á, P. Hvizdo š, M. Seleck á, Investigation of thin layers deposited by
two PVD techniques on high speed steel produced by power metallurgy, Appl. Surf.
Sci. 258(2012) 5105 -5110.
[8] Y .Y . Chang, S.Y . Weng, C.H. Chen, F.X. Fu, High temperature oxidation and
cutting performance of AlCrN, TiVN and multilayered AlCrN/TiVN hard coatings,
Surf. Coat. Technol. 332(2017) 494 -503.
[9] D. Kottfer, M. Ferdinandy, L. Kaczmarek, I. Ma ňková, J. Beňo, Investigation of Ti
and Cr based PVD coatings deposited onto HSS Co 5 twist drill, 282(2013) 770 -776.
[10] W.L. Chen, A. Yan, X.N. Meng, D.Q. Wu, D.C. Yao, D.D. Zhang,
Microstructural change and phase transformation in each individual layer of a

nano -multilayered AlCrTiSiN high -entropy alloy nitride coating upon annealing,
Applied Surface Science 462 (2018) 1017 -1028.
[11] P. Wongpanya, S. Surinphong, J. Rujisomnapa, Increasing tool life by AlCrTiSiN
film, Adanced Materials Research 853(2014) 217 -222.
[12] S.C. Liang, Z.C. Chang, D.C Tsai, Y .C Lin, H.S. Sung, M.J. Deng, F.S. Shieu,
Effects of substrate temperature on the structure and mechanical properties of
(TiVCrZrHf)N coatings, Appl Surf Sci 257(2011)7709 -7713.
[13] W.J. Shen, M.H. Tsai, Y .S. Chang, J.W. Yeh, Effects of substrate bias on the
structure and mechanical properties (Al 1.5CrNb 0.5Si0.5Ti)N x coatings, Thin Solid Films
520(2012) 6183 -6188.
[14] Y . Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw , Z.P. Lu,
Microstructures and properties high -entropy alloys, Prog. Mater. Sci. 61(2014)1 -93.
[15] W. Li, P. Liu, R.K. Liaw, Microstructures and properties of high -entropy alloy
films and coatings: a review, Mater. Res. Lett. 6(2018) 199 -229.
[16] N. Pani ch, S. Surinphong, D.A. Karpov, Y .K. Tan, Mechancial properties of
AlCrTiSiN coatings developed by cathodic arc for protection applications, Adanced
Materials Research 185(2012)81 -83.
[17] W.L. Chen, X.N. Meng, D.Q. Wu, D.C. Yao, D.D. Zhang, The effect of vacuum
annealing on microstructure, adhesion strength and electrochemical behaviors of
multilayered AlCrTiSiN coatings, Applied Surface Science 467 -468 (2019) 391 -401.
[18] S.K. Kim, P.V . Vinh, J.W. Lee, Deposition of superhard nanolayered TiCrAlSiN
thin f ilms by cathodic arc plasma deposition, Surface and Coatings Technology 202

(2008) 5395 -5399 .
[19] B.J. Xiao, H.X. Li, H.J. Mei, W. Dai, F. Zou, Z.T. Wu, Q.M. Wang, A study of
oxidation behavior of AlTiN -and AlCrN -based multilayer coatings, Surface and
Coatings Technology 333(2018) 229 -237.
[20] W.L. Chen, Y . Lin, J. Zheng, S.H. Zhang, S.Y . Liu, S.C. Kwon, Preparation and
characterization of CrAlN/TiAlSiN nano -multilayers by cathodic vacuum arc, Surf.
Coat. Technol. 265(2015)205 -211.
[21] J.L. Endrino, C. Århammar, A. Guti érrez, R. Gago, D. Horwat, L. Soriano, G.
Fox-Rabinovich, D. Mart íny Marero, J. Guo, J.E. Rubensson, J. Andersson, Spectral
evidence of spinodal decomposition, phase transformation and molecular nitrogen
formation in supersaturated TiAlN fi lms upon annealing, Acta Materialia 59 (2011)
6287 -6296 .
[22] Y . Makino, K. Nogi, Synthesis of pesudobinary Cr -Al-N films with B1 structure
by rf -assisted magnetron sputtering method, Surf Coat Technol 98(1998) 1008 -1012.
[23] http://oqmd.org/materials/composition .
[24] J. Pelleg, L.Z. Zevin, S. Lungo, Reactive -sputter -deposition TiN films on glass
substrates, Thin Solid Films 197(1991)117 -128.
[25] U.C. Oh, J.H. Je, Effects o f strain energy on the preferred orientation of TiN thin
films, J. Appl. Phys. 74(1993) 1692 -1696.
[26] L.A. Giannuzzi, M. Utlaut, Non -monotonic material contrast in scanning ion and
scanning electron images, Ultramicroscopy 111(2011)1564 -1573.
[27] S. Vepĭek, P. Nesl ádek, A. Niederhofer, F. Glatz, M. Jilek, M. Sima, Recent

process in the superhard nanocrystalline composites: towards their industrialization
and understanding of the origin of the superhardness, Surf Coat Technol.
108-109(1998) 138 -147.
[28] W.L. Chen, B. Fang, D.D. Zhang, X.N. Meng, S.H. Zhang, Thermal stability and
mechanical properties of HVOF/PVD duplex ceramic coatings produced by HVOF
and cathodic vacuum arc, Ceramics International 43 (2017) 7415 –7423 .
[29] M. Ali, E. Hamzah, I.A. Oazi , M.R.M. Toff, Effect of cathodic arc PVD
parameters on roughness of TiN coating on steel substrate, Curr. Appl. Phys. 10(2010)
471-474.
[30] W.L. Chen, J. Zheng, X.N. Meng, S.K. Kwon, S.H. Zhang, Investigation on
microstructures and mechanical properties of AlCrN coatings deposited on the surface
of plasma nitrocarburized cool -work tool steels, Vacuum 121 (2015) 194 -201.
[31] W.L. Chen, B. Fang, D.D. Zhang, X.N. Meng, S.H. Zhang, Thermal stability and
mechanical properties of HVOF/PVD duplex ceramic coatings produced by HVOF
and cathodic vacuum arc, Ceramics International 43 (2017) 7415 –7423 .
[32] D.B. Lee, T.D. Nguyen, S.K. Kim, Air -oxidation of nano -multilayered CrAlSiN
thin films between 800 and 1000 °C, Surf. Coat. Technol. 203 (2009) 1199 -1204.
[33] M. Gassner, N. Schalk, B, Sartory, M. Pohler, C. Czettl, C. Mitter, Influence of
Ar ion etching on the surface topography of cemented carbide cutting inserts, Int. J.
Refract. Met. 69 (2017) 234 -239.
[34] C. Schon jahn, D.B. Lewis, W.D. Munz, I. Petrov, Substrate ion etching in
combined steered cathodic arc -UBM deposition system: effects on interface

architecture, adhesion, and tool performance, Surf. Eng. 16(2000) 176 -180.
[35] Y .S Jung , D.W Lee , DY Jeon , Effects of nitrogen concentration on
microstructures of WNx films synthesized by cathodic arc method , Applied Surface
Science , 221 (1)(2004) 136 -142.
[36] J. Gao, W. Jie, Y . Yuan, T. Wang, G. Zha, J. Tong , Dependence of film texture on
substrate and growth conditions for CdTe films deposited by close -spaced sublimation ,
Journal of Vacuum Science & Technology A Vacuum , 29 (5)(2011)0 51507 -051507 -6.
[37] Joint Committee on Powder Diffraction Standards, Powder Diffraction File,
Inorganic Phase.

Figure captions:
Fig.1 Typical GIXRD spectra of the HEA -50, HEA -80 and HEA -110 samples.
Fig.2 (a) A typical cross -sectional BF-TEM image of a HEA -80 sample; (b) Typical
STEM -HAADF image of the framed zone in (a) ; (c) A typical line -scan concentration
profiles of Fe, Al, Cr and N element along the black line in (b) from steel substrate to
TL; (d) A HRTEM image of the rectangle zon e in (b); (e) The FFT image of (d); (f)
Magnification image of the framed zone in (d); (g) HRTEM image of two neighboring
columnar grains within TL in (b).
Fig. 3 (a) Typical SAEDP of the HEA -NL of a HEA -80 sample; (b) BF-TEM image
of several columnar grains within the HEA -NL; (c) Line-scan EDS profiles of Al, Cr,
Ti, Si and N atoms along the red line in (b) ; (d) HRTEM image of the framed zone in
(b); (e) High -magnification image of the framed zone in (d); (f-h) IFFT im ages of (e),
viewing along g=[1 1一
1]fcc, [111一
]fcc and [200] fcc directions, respectively, with their FFT
image insets; (i) HRTEM image of a columnar grain within the Grain -2, with FFT
image inset.
Fig.4 (a) Typical BF-TEM image of several columns within the HEA -NL layer; (b)
HRTEM image of the framed zone -A in (a) , with its FFT image inset ; (c) HRTEM
image of the framed zone -B in (a), with its FFT image inset ; (d) HRTEM image of the
framed zone -C in (a), with its FFT image inset .
Fig. 5 (a) A typical HAADF -STEM image of a zone within the HEA -NL of a HEA -80
sample; (b) Line-scan EDS profiles of Al, Cr, Ti, Si and N atoms along the yellow line
in (a); (c) HRTEM image of the framed “I” zone in (a), with its FFT image inset; (d)

HRTEM image of the framed “II” zone in (a), with its FFT image inset; (e) Simulated
SAEDP of (d); (f) High -magnification image of a zone within the bright zone in (a ).
Fig. 6 (a) Typical cross -sectional BF-TEM image of a HEA -110 sample; (b)
High -magnification image of the framed zone in (a); (c -d) HRTEM images of the LL
between TL and substrate; (e) FFT image of (d); (f) Simulated SAEDP of (e).
Fig. 7 (a) A typical SAEDP of the HEA -NL of a HEA -110 sample; (b) CDF image of
a zone within the HEA -NL, viewing along g=[111] reflection; (c) BF -TEM image of
another zone within the HEA -NL; (d) HRTEM image of the rectangle framed zone in
(c), with its FFT image ins et. Note that the HRTEM image is rotated by
90 ° counterclockwise relative to the framed zone in (c); (e) A typical HAADF -STEM
image of a zone within the HEA -NL; (f) HRTEM image of the framed zone in (e); (g)
Indexed SAEDP of (f).
Fig. 8 (a-c) Typical opt ical micrographs of Rockwell indentations on tested samples:
(a) HEA -50; (b) HEA -80; (c) HEA -110; and (d -f) Micrographs of scratch tests of (d)
HEA -50, (e) HEA -80 and (f) HEA -110 samples.
Fig.9 Diagrammatic sketch es of (a) Cr ion etching on substrate; (b) LL formation; (c)
Cross-sectional microstructures under relative low substrate bias voltage; (d)
Cross-sectional microstructures under high substrate bias voltage.
Fig.10 Diagrammatic sketch es of columnar grains (Simulat ing the formation process
of Grain -6, Grain -7 and Grain -8 in Fig.4): (a) Initial stage of c rystal nucleus es of the
Grain -7; (b) Formation stage of coherent interface between Grain -7 and Grain -8; (c)
Fast growth stage of the Grain -7; (d) Final formation stage of the Grain -6, Grain -7

and Grain -8. Note that C Is represent coherent interface.
Fig.1 1 Formation enthalpies of different Cr concentration (at.%) in Al 42-xCrxTi10Si3N45
system (0≤x≤42).

Tables

Table 1 Deposition parameters of experimental AlCrTiSiN HEA coating s
Procedure Substrate
bias
voltage
(V) N2
partial
pressure
(Pa) Target current (A)
Substrate
temperature
(°C) Deposition
time (min) Cr AlCr TiSi
1 -120 3 160 0 0 420 3
2 -120 3 160 160 0 420 30
3 -50 2 0 160 160 420 60
-80 2 0 160 160 420 60
-100 2 0 160 160 420 60

Table 2 Chemical composition, , , and values of AlCrTiSiN HEA
coatings.
Al
(at.%) Cr
(at.%) Ti
(at.%) Si
(at.%) N
(at.%) Al/(Al+Me)
ratio
(pm)
(kJ/mol)
(eV/atom)
HEA -50 30.3 16.3 10.1 3.4 39.9 0.53 112.3 27.9 11.4 -1.313
HEA -80 27.4 16.1 11.6 3.2 41.7 0.49 111.6 28.1 11.5 -1.331
HEA -110 26.8 18.0 11.8 3.9 39.5 0.47 112.2 27.7 11.7 -1.306
Ref.[10] 44.2 10.7 2.9 4.1 38.1 0.76 113.7 28.0 10.1 -1.244

Table 3 H, E, E*, H/E, H3/E*2 and LC2 values of tested samples.
H (GPa) E (GPa) E* (GPa) H3/E*2 (GPa) LC2 (N)
HEA -50 32.1±0.8 216.7±3.2 238.2±3.2 0.584±0.045 44.0±1.2
HEA -80 33±0.4 238.8±1.4 262.4±4.5 0.521±0.018 56.0±2.3
HEA -110 36.5±1.1 245.6±2.7 269.9±3.2 0.665±0.042 78.0±1.8
Ref. [10] 39.4 307.2 337.6 0.537 —

Figure-1
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Figure-2
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Figure-3
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Figure-4
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Figure-5
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Figure-6
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Figure-7
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Figure-8
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Figure-9
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Figure-10
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Figure-11
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