Influence of pulsed bias duty cycle variations on structural and mechanical
properties of arc evaporated (Al,Cr) 2 O3 coatings
Markus Pohler, Robert Franz, Jürgen Ramm, Peter Polcik, Christian
Mitterer
PII:
DOI:
Reference:
S0257-8972(15)30293-0
doi: 10.1016/j.surfcoat.2015.09.055
SCT 20608
To appear in:
Surface & Coatings Technology
Received date:
Revised date:
Accepted date:
27 May 2015
28 September 2015
29 September 2015
Please cite this article as: Markus Pohler, Robert Franz, Jürgen Ramm, Peter Polcik,
Christian Mitterer, Influence of pulsed bias duty cycle variations on structural and mechanical properties of arc evaporated (Al,Cr)2 O3 coatings, Surface & Coatings Technology
(2015), doi: 10.1016/j.surfcoat.2015.09.055
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Influence of pulsed bias duty cycle variations on structural and
mechanical properties of arc evaporated (Al,Cr)2O3 coatings
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Markus Pohler a,1, Robert Franz a, Jürgen Ramm b, Peter Polcik c, Christian Mitterer a
Christian Doppler Laboratory for Advanced Hard Coatings at the Department of Physical
Metallurgy and Materials Testing, Montanuniversität Leoben, Franz-Josef-Straße 18, A-8700
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Leoben, Austria
Oerlikon Surface Solutions AG, Iramali 18, LI-9469 Balzers, Principality of Liechtenstein
c
PLANSEE Composite Materials GmbH, Siebenbürgerstraße 23, D-86983 Lechbruck am
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b
D
See, Germany
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Abstract
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(AlxCr1-x)2O3 coatings were synthesised by cathodic arc evaporation applying DC and pulsed
substrate bias voltages to study the impact of different bias duty cycle settings on structure
and morphology as well as mechanical and tribological properties. X-ray diffraction revealed
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a corundum-type (AlxCr1-x)2O3 phase structure for Al contents x ≤ 0.5 and the formation of an
additional face-centred cubic (Al-Cr-O) phase for x ≥ 0.5 for the pulsed bias situation,
independently of the chosen duty cycle. In general, changes of the duty cycle in the manner of
decreasing the electron current time duration along with an extension of the ion bombardment
during growth lead to a reduction of the coatings’ domain size, the formation of compressive
residual stress as well as increasing coating hardness and Young’s modulus. An
amorphous-like coating microstructure accompanied with diminished mechanical properties
1
Corresponding author.
Email address: markus.pohler@ceratizit.com (Markus Pohler)
Present address: CERATIZIT Austria GmbH, A-6600 Reutte, Austria
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was observed when a DC substrate bias voltage was applied. The investigations illustrate the
potential of tuning the properties of arc evaporated oxide coatings by adjusting the energetic
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particle bombardment during growth.
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Keywords: Arc evaporation, (Al,Cr)2O3, pulsed bias, microstructure, mechanical properties,
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tribology
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1 Introduction
During the last years, (AlxCr1-x)2O3 hard coatings grown by physical vapour deposition (PVD)
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processes have gained increased attention due to the fact that they can be synthesised as a
solid solution over a broad compositional range with good mechanical properties, high
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thermal stability and chemical inertness [1-4]. The addition of Cr2O3 facilitates the formation
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of the desired corundum-type Al2O3 structure at moderate deposition temperatures while other
polymorphs, which are described in literature and were observed in the synthesis of
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non-alloyed Al2O3 are suppressed [5,6]. Different PVD deposition techniques such as reactive
sputter deposition and cathodic arc evaporation can be used to synthesise such (AlxCr1-x)2O3
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coatings at substrate temperatures below 600 °C. This represents a major advantage as it
enables the use of temperature sensitive substrate materials like hardened and tempered high
speed steels. Such materials are incompatible with standard PVD and chemical vapour
deposition (CVD) processes for the synthesis of corundum-type Al2O3 based coatings, since
substrate temperatures above 700 °C are required [7,8]. Thus, the ternary (AlxCr1-x)2O3 oxides
are promising candidates for applications as wear protective tool coatings, high-temperature
solar absorbing coatings or hydrogen diffusion barriers [9,10]. However, up to now the
desired thermodynamically stable corundum-type crystal structure can only be grown up to an
Al content of x ~0.7 [1,11] or x ~0.85 if a feasible seed layer is utilised [12]. At higher Al
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contents, an additional metastable face-centred cubic (fcc) (Al-Cr-O) polymorph is formed or
an amorphous-like coating structure arises [13-16].
Owing to the high degree of ionisation in the cathodic arc plasma employed to deposit the
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oxide coatings, the application of a substrate bias voltage is a powerful tool to control the
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growth conditions of the coating and, thus, to tune their morphology, microstructure as well
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as the mechanical and physical properties [17-19]. Moreover, the incorporation of
macro-particles can be significantly influenced by the application of a substrate bias voltage
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[20,21]. In general, the substrate can be biased to a DC negative a unipolar pulsed or a bipolar
pulsed voltage. The latter enables boosted substrate heating and may reduce accumulations of
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positive charges on the substrate surface originating from the bombardment with positively
charged ions. This decreases the risk of formation of generally unwanted electric micro-arcs
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on the substrate surface [22] or a reduction of the coating growth rate based on repulsion of
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ions approaching the surface. Therefore, the technique is well suited for the synthesis of poor
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or even non-conductive coatings (e.g. oxide coatings) [23,24]. In general, the resulting coating
properties can be optimised by selecting proper bias parameters, e.g. voltage amplitude, duty
cycle or pulse frequency. The alternating attraction and repulsion of ions and electrons from
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the plasma leads to modulations of their impact energies on the substrate. For rectangular
shaped unipolar bias pulses an extended period of ion attraction leads to higher average ion
impact energies on the substrate and a more monoenergetic ion energy distribution as
compared to small duty cycles as it was shown in refs. [25-27]. Further, the electron
movement within the substrate sheath forced by applying a pulsed bias may also lead to
increased ionisation of the reactive gas and the metal vapour by electron impact ionisation.
The focus of the present work was to further elucidate the effect of the application of a bipolar
pulsed substrate bias voltage on growth and properties of arc-evaporated (AlxCr1-x)2O3
coatings in order to enable further optimisation of the deposition parameters for the synthesis
of such materials. Therefore, the influence of different substrate bias duty cycle settings on
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coating morphology and microstructure as well as mechanical and tribological properties of
(AlxCr1-x)2O3 coatings with x = 0.25, 0.5 and 0.7 was intensively investigated.
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2 Experimental details
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2.1 Coating synthesis
A production-scale Oerlikon Balzers INNOVA cathodic arc evaporation system was used to
synthesise Al-Cr-O coatings from metallic AlxCr1-x composite cathodes with nominal atomic
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ratios of x = 0.25, 0.5 and 0.7. The cathodes were produced by powder metallurgy (pressing
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of the metallic Al and Cr powders and hot forging) [28] with mean grain sizes of the used
powders in the range of 100 µm. The cathodes were circular in shape with a diameter of
150 mm and are arranged at separate heights within the deposition chamber (diameter
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1200 mm, height 980 mm). In the further course of the text, the atomic ratio x of the AlxCr1-x
cathodes will also be used to distinguish between the different coating compositions.
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The residual base pressure within the process chamber was below 10-3 Pa and the deposition
temperature was kept constant at 550 °C using a radiation heating system for all depositions.
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Mirror polished Si strips (Si (100), 7 mm × 20 mm × 0.35 mm), cemented carbide cutting
inserts (SPGN 120308, WC-6 wt.% Co) and high speed steel (DIN 1.3343, AISI M2) discs
( 30 mm × 10 mm) with a hardness of 65 HRC were used as substrates and were mounted
on a carousel allowing for a two-fold rotation. The minimum distance between substrate and
cathodes was 25 cm. Prior to the deposition of the oxide layers, the substrates were sputteretched in pure Ar plasma and a ~0.3 µm thick TiAlN base layer as well as a graded ~0.1 µm
thick Al,Cr(O,N) transition layer were deposited to improve the adhesion and avoid oxidation
of the substrates. Thereby, the Al,Cr(O,N) layer was deposited during a continuous change of
the reactive gas atmosphere from pure nitrogen to a oxygen rich O2/N2 gas mixture.
All oxide layers were grown in a pure O2 atmosphere with a gas flow of 400 sccm
corresponding to a pressure of ~0.9 Pa, a DC arc current of 180 A at each of the two AlxCr1-x
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cathodes and a deposition time of 80 min. These parameters correspond to a growth rate of
approximately 2.5 µm/h for the (AlxCr1-x)2O3 functional layer.
In order to prevent detrimental effects due to charging of the substrate surface covered with an
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insulating coating, a bipolar-pulsed substrate bias with an amplitude of 60 V and a pulse
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frequency  of 25 kHz was applied for the transition and the top layer. Here, the duty cycle 
t neg
t neg  t pos
,
(1)
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
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is defined as
with tneg the negative pulse time, tpos the positive pulse time and  = tneg + tpos the total pulse
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period. For each coating composition, values of 70, 90 and 95% for  were applied. In
addition a duty cycle of 100%, which is equal to a DC bias voltage, was applied for the
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coatings with x = 0.5 and 0.7. The used bias power supply was specifically developed for the
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INNOVA system to handle high substrate currents in arc deposition systems.
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The time-dependent voltage and current characteristics of the effective substrate bias, as
compared to the settings at the bias power supply, were recorded with a Tektronix TPS 2024
digital oscilloscope connected with the rotating substrate holder during additional evaporation
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processes from two x = 0.5 cathodes with evaporation parameters similar to the ones
described for coating synthesis.
2.2 Coating characterisation
Elastic recoil detection analysis (ERDA) using a 35 MeV Cl7+ ion beam with an analysed area
of 1.5 mm × 1.5 mm and an information depth of ~600 nm was used to determine the
chemical composition of the oxide layer. The obtained spectra were fitted using the ion beam
analysis DataFurnace (NDF) software package [29].
Coating thickness and morphology were characterised from fracture cross-sections of coated
cemented carbide substrates using a Zeiss EVO 50 scanning electron microscope (SEM). The
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surface roughness of the coatings in the as deposied state was determined using a Leica
DCM3D dual core profilometer in confocal mode with an EPI 150X-L objective.
Structural investigations of the coatings were conducted by X-ray diffraction (XRD) in Θ-2Θ
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mode as well as grazing incidence (GIXRD) geometry with an angle of incidence of 2°. The
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measurements were performed on a Bruker-AXS D8 Advance diffractometer applying Cu-Kα
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(wavelength  = 1.54056 nm) radiation. The device was equipped with a Goebel mirror and
an energy-dispersive X-ray detector. For analysing the phases present in the XRD patterns,
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the ICDD database [30] as well as lattice parameters for different (AlxCr1-x)2O3 solid solutions
reported in ref. [31] were used. All presented XRD results were determined from
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measurements on coated Si substrates. In order to determine domain size and lattice parameter
from the diffraction patterns, Rietveld refinement was done utilising the Bruker-AXS Topas
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4.2 software package. For the phase analysis by Rietveld refinement all diffraction peaks
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between 20 and 70° 2Θ were taken into account.
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The room temperature residual stress state was derived from the wafer curvature
measurements on Si strips. Two parallel laser beams were used to measure the substrate
curvature radius from which the biaxial film stress was calculated using the modified Stoney
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equation [32]. The presented stress values are averaged from at least 2 measurements on
different Si strips per coating.
Coating hardness (H) and Young’s modulus (E) were obtained by means of nanoindentation
using a UMIS device (Fischer-Cripps Laboratories) which was equipped with a Berkovich
diamond tip. The obtained load-displacement curves were corrected according to the Oliver
and Pharr method [33]. Prior to the measurements, the coating surfaces were polished with a
1 µm diamond paste in order to reduce the influence of the surface roughness on the
nanoindentation results. The hardness values and moduli were calculated from indentations
using maximum loads in the range from 4 to 16 mN resulting in maximum indentation depths
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of 220 nm, which is less than 10% of the thickness of the oxide layer. The presented values
are averaged from a minimum of 10 indents per sample.
A CSM ball-on-disc tribometer was used to investigate the tribological behaviour of the
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coatings at room temperature. The tests were performed in ambient atmosphere with a
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temperature between 22 and 25 °C and a relative humidity between 20 and 45%. All coatings
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were tested against alumina ball counterparts ( 6 mm). Normal load, sliding speed, sliding
distance and radius of the wear track were kept constant at 5 N, 10 cm/s, 500 m and 7 mm,
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respectively. The evaluation of the wear tracks was done by optical profilometry using a
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Veeco Wyko NT1000 white light interferometer.
3 Results and discussion
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3.1 Bias voltage and current
The time dependent characteristics of substrate bias voltage and substrate current are plotted
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in Fig. 1 for a nominal duty cycle setting of 90%. According to the voltage signals (measured
and nominal), it is obvious that the actual substrate bias voltage deviates from the bipolar and
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symmetrically pulsed voltage that was applied to the substrate. Beginning at the end of the
negative pulse regime (about −60 V), the voltage is decreased with a high slew rate until a
reduction of the voltage gradient at about −20 to −15 V arises. The voltage decrease proceeds
until the end of the “positive” pulse portion (tpos) where the minimum of the negative voltage
of about −10 V is reached, which is slightly below the floating potential of approximately
−18 V, as measured on the unbiased substrate. The low positive pulse portion (as referred to
the floating potential) can be understood by the high mobility of the electrons in the plasma.
Even at a slightly negative substrate potential, energetic electrons can reach the substrate.
Upon applying a positive voltage on the substrate, more and more electrons are attracted
resulting in a strong increase of electron current which, in turn, delays establishing a higher
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positive potential on the substrate. Therefore, no saturation current for the electrons can be
observed for the chosen short positive pulse time.
The substrate current rises with a constant slope during the nominal positive pulse regardless
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of the duty cycle. For the exemplified duty cycle of 90% a maximum electron current of
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~12 A was reached. At the beginning of the nominal negative pulse (tneg), the situation is
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reversed and the current decreases until the ion saturation current for the chosen arc
parameters at about −7.5 A is reached. Simultaneously, the applied substrate voltage slowly
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decreases, but remains close to the floating potential as long as the substrate current is
positive.
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However, the time dependence of the substrate voltage and current signals is based on
charging effects of the insulating (Al1-xCrx)2O3 coating and the subsequent formation of
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electric fields as well as the variation of electron and ion densities within the sheath [25]. In
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addition, capacitances and inductances present in the electrical circuit of the used deposition
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system might also contribute to the observed voltage and current evolution. In terms of thin
film growth, the most important outcome is the high electron current during the positive pulse
portion which is even increased for increasing tpos (not shown) and the deviation between the
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nominal and the effective substrate bias duty cycle, i.e. the duration of the ion bombardment.
However, it has to be mentioned, that the maximum ion current and particle energy cannot be
influenced to a big extent by changing the duty cycle as it is the case for, e.g., varying the
amplitude of the bias voltage or the ion charge state and energy distribution by changing the
pressure of the used background gas [25,34-37].
3.2 Chemical composition and coating morphology
All investigated oxide coatings reveal a stoichiometric (AlxCr1-x)2O3 composition within the
measurement accuracy with an oxygen content of about 59  1 at.% as determined by ERDA.
In Table 1 the chemical composition of the coatings with x = 0.25, 0.5 and 0.7 synthesised
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with a substrate bias duty cycle of 90% is presented. The level of H and C impurities was
below 1 at.% for all investigated coatings. No traces of nitrogen were found in the
surface-near region of the coatings. In all cases, the Al/(Al+Cr) atomic ratio of the coatings is
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slightly higher than that of the corresponding cathode. This moderate Cr deficiency is in good
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agreement with recently published results on arc evaporated (Al1-xCrx)2O3 coatings [3,13,38].
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However, no significant influence of the substrate bias duty cycle on the coating composition
was found and, as an example, the comparison between a duty cycle of 70 and 90% for the
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coating with x = 0.5 is also given in Table 1.
Fig. 2 shows SEM top-view micrographs of the coatings with x = 0.5 on cemented carbide
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substrates for all substrate bias duty cycles used. A pronounced surface roughness mainly due
to the incorporation of macro-particles is evident regardless of the duty cycle settings. The
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measured Sq values are in the range of 220  40 nm without apparent correlation to the change
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in duty cycle. However, the surface roughness increases with increasing Al content as the Sq
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values for x = 0.25 and 0.7 are about 135  30 nm and 340  70 nm, respectively.
The corresponding SEM fracture cross-section images of the coatings with x = 0.5 are shown
in Fig. 3 and reveal a coarse and highly columnar fracture morphology for duty cycles of 70
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and 100% whereas, a fine structured appearance is obvious for the coating deposited with the
95% duty cycle setting. A transition state is observed for a duty cycle of 90%. In addition, the
applied duty cycle variation is without apparent influence on the coating thickness (see
Fig. 3). Especially for high duty cycles, a remaining continuous background flow of low
energetic ions toward the substrate can be assumed even for the positive pulse period due to
the expected inertia of the heavy ions. This in combination with the continuous flux of
evaporated neutrals or neutralised ions, which are not influenced by the applied bias, might
allow for the observed roughly duty cycle independent growth rates. Furthermore, a constant
coating thickness also implies equilibrium between coating growth and ion energy supported
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resputtering of adherent particles [39]. Both mechanisms are supposed to be favoured by
increased ion attraction for high duty cycles.
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3.3 Coating microstructure
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3.3.1 Phase composition
The microstructural evolution of the coatings with x = 0.5 as a function of the substrate bias
duty cycle is presented in Fig. 4 where the corresponding GIXRD patterns are shown. All
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coatings exhibit a dual-phase structure of the oxide top layer consisting of corundum-type
(Al0.5Cr0.5)2O3 (space group R-3c) together with a fcc-(Al-Cr-O) phase. Increasing the bias
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duty cycle from 70 to 95% leads to increased peak intensities of the predominant corundumtype (Al0.5Cr0.5)2O3 phase, while the intensities of the fcc-(Al-Cr-O) phase are slightly reduced
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(c.f. Fig. 4). Thus, growth of the corundum phase fraction at the expense of the fcc phase can
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be assumed. However, switching from the bipolar pulsed mode to a DC substrate bias of
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−60 V (i.e. 100% duty cycle) results in a significant reduction of the corundum-type peaks
indicating a preferred growth of a fine grained fcc phase fraction in an amorphous-like matrix.
The results of the structural investigations for the coatings with x = 0.25 and 0.7 are
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summarized in Fig. 5 where the GIXRD scans for substrate bias duty cycles of 70 and 95%
for both coating compositions are presented. Therein, a dual-phase structure with mainly
fcc-(Al-Cr-O) in an amorphous matrix and traces of a nano-crystalline corundum-type phase
as indicated by the broad feature between 32° and 37° 2θ is obvious for the coatings with
x = 0.7. With increased duration of the ion bombardment (i.e. duty cycle 95%), a moderately
increased intensity of this feature is observed, suggesting a progressive growth and
crystallisation of the nano-crystalline or amorphous-like (AlxCr1-x)2O3 domains.
For the coatings with x = 0.25 (see Fig. 5), a single-phase corundum-type structure is
observed independent of the applied bias duty cycle. Thus, the increased Cr content for the
coating with x = 0.25 facilitates the growth of the corundum-type crystal structure, while a
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high Al content favours the formation of the fcc-(Al-Cr-O) polymorph, which is in good
agreement with literature [12,15].
Similar to the coatings with x = 0.5 (c.f. Fig. 4), also the diffractograms for the coatings with
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x = 0.25 and 0.7 (c.f. Fig. 5) show additional small peaks which are caused by the Ti0.5Al0.5N
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interlayer (lattice parameters taken from ref. [40]) and metallic macro-particles incorporated
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into the oxide coating. More precisely, the existence of AlCr solid solution, tetragonal Al2Cr,
intermetallic Al8Cr5 and Al4Cr compounds is suggested for x = 0.5 and x = 0.7 (with
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diffraction intensities between 35° and 38° 2θ as well as between 40° and 45° 2θ), whereas a
body-centred cubic (bcc) Cr(Al) solid solution can be detected for the high Cr containing
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coatings with x = 0.25 [38,41]. The latter is evidenced by the asymmetric left shoulder of the
peak at ~45° 2θ. In order to allow for an easy readability of Figs. 4 and 5, the exact positions
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of the described intermetallic phases are not displayed.
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In summary, especially the dual-phase coatings with a predominant corundum-type phase
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composition (e.g. x = 0.5) exhibit a pronounced dependency of their microstructure on the
duty cycle. On the other hand, the structure of the single-phase corundum-type (x = 0.25) and
the predominant fcc-(Al-Cr-O) coatings (x = 0.7) seems to be rather unaffected within the
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investigated duty cycles as long as a bipolar pulsed bias is used. Apparently, for the transient
composition of x = 0.5 only a small additional activation energy is needed during film growth
in order to stimulate the corundum-type crystal structure. In fact, the interaction of ion and
electron radiation with the growing coating might contribute sufficient energy to initiate
surface diffusion of adatoms and out-annealing of vacancies, thus promoting the formation of
the energetically favourable corundum phase and suppressing the growth of the metastable
vacancy stabilised fcc structure [14]. The possibility of influencing the mobility of adatoms
and, thus, the coatings microstructure by changing the ion flux to the substrate was also
discussed in more detail in refs. [42,43].
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3.3.2 Biaxial residual stress
A comparison of the total residual stress in the as-deposited state for the different substrate
bias duty cycle settings is shown in Fig. 6. Since the deposition parameters and the thickness
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for the TiAlN and (Al,Cr)O,N interlayer are identical for all coatings, the measured relative
variations in the total stress level originate from the Al-Cr-O top layer which was deposited
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with different substrate bias duty cycle settings. In general, the arising macro-stress level is
rather low with values between ~0.3 GPa tensile and ~0.5 GPa compressive stress. Increasing
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the duty cycle from 70 to 95% shifts the total residual stress towards the compressive regime
for all investigated coating compositions. However, this development is not continued for a
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negative DC substrate bias (100% duty cycle), where a reduction of the compressive stresses
and even the formation of moderate tensile stress (e.g. for the coating with x = 0.5) can be
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observed. This stress evolution can be understood by a reduced particle bombardment as a
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consequence of charging effects on the growing film surface.
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For low duty cycles, on the other hand, the reduction of the compressive stress is most likely
based on defect annihilation due to the enhanced electron bombardment (see e.g. [44]).
However, only slight changes in the total coating stress are observed for all coating
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compositions, which indicate only a moderate variation in the maximum energy of the
bombarding ions. This can be attributed to the chosen O2 pressure of 0.9 Pa. According to
Franz et al. [34,37], at high working gas pressure the fraction of multiple-charged ions is
strongly reduced and most of the remaining ions are thermalised due to frequent collisions
with neutrals. As a result, a maximum energy Ei of ~60 eV can be expected which is
independent of the duty cycle. Only an increase in the negative bias voltage will cause distinct
changes in the coating texture as it was, e.g., observed in our previous study where the bias
voltage was changed from −60 V up to −160 V [12].
The above discussion is focused on the influence of the plasma properties with changing duty
cycle on the stress state in the coatings. For a more complete stress analysis further effects,
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e.g. possible differences in the thermal expansion coefficient between the increasing fcc phase
content and the corundum-type (Al-Cr-O) polymorph also need to be taken into account.
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3.3.3 Domain size
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The evolution of the average coherently diffracting domain sizes for the different coatings
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with x = 0.25 and 0.5, as determined from θ-2θ XRD scans (not shown) using the
corundum-type peaks between 2θ angle from 20 to 70°, is presented in Fig. 7. For both
compositions, an extension of the negative bias pulse duration within the bipolar regime leads
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to a moderate reduction of the domain size from ~130 to ~85 nm for the coatings with x = 0.5
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and from ~110 to ~70 nm for x = 0.25. This reduction in the domain size is also related to a
reduction in column width for higher duty cycles and corroborates the assumptions made
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from the above presented SEM cross-sections (c.f. Fig. 3). Stimulation of surface migration of
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deposited particles by surface heating and energy transfer from incident electrons on the
growing surface is supposed to be the main reason for the increased domain size for low
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substrate bias duty cycles, where the highest electron currents were measured [45,46]. In
addition, the increased portion of impinging ions, as it is the case for high duty cycles, would
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cause pronounced generation of growth defects as well as lattice distortion within the coating.
This in turn promotes the formation of additional nucleation sites, a higher re-nucleation rate
and, thus, smaller domains [18]. Such an evolution is well consistent with the described
formation of moderate compressive stress [17,47].
The application of a negative DC substrate bias voltage for the coatings with x = 0.5 is
accompanied by a strong decrease of the domain size down to ~8 nm (c.f. Fig. 7), which is in
good agreement with the decreasing crystallinity of this coating as measured by XRD. An
evaluation of the domain sizes of the corundum-type phase for the coating with x = 0.7 was
not feasible due to low diffraction intensities.
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3.3.4 Lattice parameters
The lattice parameters of the corundum-type phase fraction of the coatings with x = 0.25 and
0.5 are a = 0.492 nm and c = 1.345 nm as well as a = 0.487 nm and c = 1.330 nm,
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respectively, as determined by Rietveld refinement. However, the lattice parameters are
independent from the settings for the substrate bias duty cycle. In general, the obtained values
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are in good agreement with published data for sintered bulk material as well as arc evaporated
coatings with an Al/(Al + Cr) ratio of ~0.25 and ~0.5 and thus corroborate the chemical
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composition as obtained by ERDA [1,31].
The experimentally obtained lattice parameters of the fcc-(Al-Cr-O) phase fraction are
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a = 0.401 and 0.398 nm for the coatings with x = 0.5 and 0.7, respectively, and correspond
well with the calculated lattice parameter in a recently reported theoretical study on a
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metastable B1-like vacancy stabilised (Al0.5Cr0.5)2O3 structure [48]. Due to the decreasing
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lattice parameter with increasing Al content, substitution of Cr by Al ions within this structure
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can be assumed, indicating the existence of a fcc-(Al-Cr-O) solid solution with a wide
compositional existence range which is in good agreement with results published by Khatibi
et al. [15]. This is also corroborated by earlier findings, where a cubic lattice parameter of
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a = 0.404 nm for an AlCrO coating with an Al/(Al+Cr) fraction of ~0.31 was reported
[14,49]. Thus, the nearly linear evolution of the fcc-(Al-Cr-O) lattice parameter with
increasing Al content would be in fair agreement with a Vegard-like behaviour resulting in
extrapolated lattice parameters a = 0.409 and 0.394 nm for an Al content of 0 and 100 at.% on
the metallic sublattice, respectively. Several studies support the hypothesis of the existence of
a vacancy stabilised fcc-CrO monoxide (B1 NaCl structure) with a suggested lattice
parameter of a = 0.408 nm [50-52].
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3.4 Coating hardness and Young’s modulus
In Fig. 8 hardness and Young’s modulus of the investigated (AlxCr1-x)2O3 coatings with
x = 0.25, 0.5 and 0.7 are presented as a function of the applied substrate bias duty cycle. The
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single-phase corundum-type coatings with x = 0.25 exhibit an almost constant hardness at
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~26 GPa for duty cycle settings of 70, 90 and 95%. For the coatings with a dual-phase
microstructure (x = 0.5 and 0.7), an increasing hardness with increasing duty cycle is
noticeable. Both coating compositions show hardness values of ~20 and ~24.5 GPa for a
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substrate duty cycle of 70 and 90%, respectively. A further increase of the bias duty cycle up
to 95% results in a moderately increased hardness of ~25 GPa for the coating with x = 0.7,
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while a peak hardness of ~32 GPa was measured for the coating with x = 0.5. However,
changing from the bipolar pulsed to a negative DC substrate bias (duty cycle 100%) results in
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considerably decreased hardness values of ~18 GPa independent of the Al/(Al+Cr) ratio.
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This evolution of the hardness is in good accordance with the corresponding Young’s
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modulus (see Fig. 8b). While the modulus of the single-phase corundum coating with x = 0.25
is in the range of ~300 to ~310 GPa and therefore essentially unaffected by the applied bias
duty cycle, the coatings with higher Al contents show a distinct dependence of the modulus
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on the used duty cycle settings. Again, an increasing duty cycle within the bipolar pulsed bias
regime yields an increased modulus. Thereby, the coatings with x = 0.5 exhibit values
between ~280 and ~360 GPa. The indentation moduli for the coatings with x = 0.7 are slightly
lower, ranging from ~250 to ~275 GPa. Nevertheless, for all coating compositions a drop of
the Young’s modulus to ~250 GPa is observed after switching from the bipolar to the DC bias
mode.
Variations in the ion flux and decreasing electron bombardment for increasing the substrate
bias duty cycle may hinder out-annealing of growth defects within the coatings, and thus
promote the formation of compressive stress and the reduction of domain size [44]. These
structural changes (also including the described modifications in the phase composition)
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enable the observed increase of hardness and Young’s modulus for the coatings synthesised
with high bias duty cycles [19,53]. However, in absence of a periodic electron bombardment
(i.e., a DC negative bias) the increased amount of amorphous-like domains and the formation
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of the metastable fcc-(Al-Cr-O) phase in combination with a pronounced columnar coating
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morphology results in significant deteriorating mechanical properties for the coatings with
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x = 0.5 and 0.7. In general, it is assumed that there is a distinct correlation between the
mechanical properties and the volume fraction of the remaining amorphous-like regions and
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the fcc-(Al-Cr-O) phase (c.f. section 3.3.1). Again, especially the coatings with a pronounced
dual-phase microstructure and a concurrently well-developed corundum-type phase fraction
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(x = 0.5) seem to be sensitive to variations of the substrate bias duty cycle concerning their
hardness and Young’s modulus. Moreover, the predominant fcc structured coatings (x = 0.7)
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exhibit a distinct increase of hardness and modulus with increasing duty cycle. This is
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attributed to the starting crystallisation of the amorphous regions. Obviously, compressive
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stress formation (c.f. Fig. 6) together with the moderate growth of the corundum phase at the
expense of the amorphous fraction as indicated by GIXRD yields the slight hardening effect.
The latter coincides with findings on sputtered Cr2O3 after increasing the crystallinity of the
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coating [54]. However, the coatings with the single-phase corundum microstructure (x = 0.25)
show only minor effects of the duration of the ion bombardment on their mechanical
properties although a significant decrease of the domain size was measured.
3.5 Tribological properties
The results of the room temperature ball-on-disc tribometer tests using an Al2O3 sliding
partner revealed an average steady-state friction coefficient µ of ~0.55 for the low Al
containing coatings with x = 0.25 and 0.5. The coating with x = 0.7 showed a slightly higher
average value of ~0.65. As it turned out, the steady-state friction coefficient is independent of
the applied substrate bias duty cycle for all investigated coating compositions and the constant
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friction regime could be reached after a running-in period of about 70 m (dominated by
polishing of the sliding surfaces).
The wear behaviour of the different coatings after a sliding distance of 500 m is summarised
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in Fig. 9 where the 2D wear track profiles for the coatings synthesised with a substrate bias
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duty cycle of 95% are shown. It is obvious that a predominant adhesive wear characteristic is
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present for all investigated coatings and abrasive loss of materials is only detectable for the
coatings with x = 0.7. Thereby, the amount of the adherent transfer material as well as the
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width of the wear track continuously increases with increasing Al content in the coatings. This
indicates an enhanced formation of wear debris either generated from the ball or the coating
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itself (e.g. ripped out macro-particles [55]) and is, thus, in good agreement with the increasing
friction coefficient for the coating with x = 0.7 since the wear particles also enhance
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third-body abrasion which can cause higher friction [56]. As a direct consequence, the
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formation of distinct wear grooves owing to ploughing, increased adhesion of transfer
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material next to the wear scar as well as an increased wear on the Al2O3 sliding partner can be
observed for the high Al containing coatings. However, all investigated coatings were not
worn through during the tribological tests or showed other coating failure. In particular, the
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coating with x = 0.25 was almost unaffected by the applied tribological load.
In contrast to the friction coefficient, the wear of the coating as well as the wear of the Al2O3
sliding partner is noticeably influenced by the adjusted substrate bias duty cycle. The wear
track width on the coating and also the wear scar width on the Al2O3 ball is increased with
increasing duty cycle for all investigated coating compositions. However, the lowest wear of
coating and ball was observed for the coatings synthesised in the DC bias mode. Light optical
micrographs of the worn Al2O3 balls used against the coatings with x = 0.5 for all duty cycles
are presented in Fig. 10. The diameter of the wear scar continuously grows from 170 to 185,
200 and finally 220 µm when changing the bias duty cycle from 100 to 70, 90 and 95%,
respectively. In turn, the adhesive wear of the coatings increases in the same order. The
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highest amount of transferred wear debris can be found in the wear track of the coatings
deposited with a substrate bias duty cycle of 90 and 95%. It is assumed that for coatings with
identical chemical composition higher coating hardness promotes abrasion of the alumina ball
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counterpart during the sliding contact. Thereby, surface asperities on the coating may act as
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abrasive edges for the ball. However, it should be considered that also changes in the surface
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morphology of the different coatings and modifications in the phase composition due to
4 Summary and conclusions
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variations in the substrate bias duty cycle will influence the wear behaviour.
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Within this study, a series of (AlxCr1-x)2O3 coatings with x = 0.25, 0.5 and 0.7 were
synthesised by reactive cathodic arc evaporation using DC and pulsed substrate bias voltage
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with different duty cycle settings. With increasing Al content, a change in the coatings’
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microstructure from a single-phase corundum-type (AlxCr1-x)2O3 to a dual-phase
corundum-type (AlxCr1-x)2O3 and fcc-(Al-Cr-O) and finally to a mainly amorphous-like
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structure with traces of the fcc and corundum-type phases was detected.
For increasing duty cycles, a decrease in domain size, the formation of compressive stress and
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improved mechanical properties were observed. This was accompanied by changes in the
microstructure as the growth of the corundum-type (AlxCr1-x)2O3 phase fraction was found to
be favoured for high substrate bias duty cycle settings. However, the application of a negative
DC substrate bias corresponding to a duty cycle of 100% resulted in the growth of a very fine
grained, amorphous-like structure and inferior mechanical properties. The evolution of
coating structure and properties could be attributed to variations of the energetic particle flux
onto the growing film with changing the substrate bias duty cycle. In particular, a strong
increase of the electron current and a significant reduction of the ion bombardment with
decreasing duty cycle was found to cause the growth of larger grains with the consequences of
reduced coating stress and hardness.
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In addition, the impact of the substrate bias duty cycle on the coating properties was noticed
to depend significantly on the Al content x and, hence, on the initial phase composition. The
most pronounced changes in structure and mechanical properties upon variation of the
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substrate bias duty cycle were found for the coating with x = 0.5 showing a
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corundum-dominated dual-phase structure.
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In summary, the presented results reveal that the choice of the pulsed substrate bias setting
has a distinct impact on the morphology as well as structural and mechanical properties of arc
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evaporated (AlxCr1-x)2O3 coatings. Especially an appropriate selection of the ratio between ion
and electron attraction during the growth of the investigated non-conducting coatings is of
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crucial importance for tuning the material properties with respect to the desired application.
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Acknowledgements
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Financial support by the Christian Doppler Research Association is highly acknowledged.
The authors are grateful to Sarah Ploberger and Nadine Raidl for the experimental work as
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well as Gerhard Hawranek (Department of Physical Metallurgy and Materials Testing,
Montanuniversität Leoben) for performing the SEM Investigations. This work has been
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supported by the European Community as an Integrating Activity “Support of Public and
Industrial Research Using Ion Beam Technology (SPIRIT)” under EC contract no. 227012.
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References
[1]
J. Ramm, M. Ante, T. Bachmann, B. Widrig, H. Brändle, M. Döbeli, Surf. Coat.
[2]
IP
T
Technol. 202 (2007) 876-883.
M. Witthaut, R. Cremer, K. Reichert, D. Neuschütz, Mikrochim. Acta 133 (2000) 191-
[3]
SC
R
196.
V. Edlmayr, M. Pohler, I. Letofsky-Papst, C. Mitterer, Thin Solid Films 534 (2013)
K. Pedersen, J. Bøttiger, M. Sridharan, M. Sillassen, P. Eklund, Thin Solid Films 518
MA
[4]
NU
373-379.
(2010) 4294-4298.
I. Levin, D. Brandon, J. Am. Ceram. Soc. 81 (1998) 1995-2012.
[6]
W. Sitte, Mater. Sci. Monog. 28A (1985) 451-456.
[7]
S. Ruppi, J. Phys. IV 11 (2001) Pr3847-Pr3859.
[8]
O. Zywitzki, G. Hoetzsch, Surf. Coat. Technol. 76-77 (1995) 754-762.
[9]
H.D. Liu, Q. Wan, B.Z. Lin, L.L. Wang, X.F. Yang, R.Y. Wang, D.Q. Gong, Y.B.
AC
CE
P
TE
D
[5]
Wang, F. Ren, Y.M. Chen, X.D. Cheng, B. Yang, Sol. Energy Mater. Sol. Cells 122
(2014) 226-232.
[10]
D. Levchuk, H. Bolt, M. Döbeli, S. Eggenberger, B. Widrig, J. Ramm, Surf. Coat.
Technol. 202 (2008) 5043-5047.
[11]
D. Diechle, M. Stueber, H. Leiste, S. Ulrich, V. Schier, Surf. Coat. Technol. 204
(2010) 3258-3264.
- 20 -
ACCEPTED MANUSCRIPT
[12]
M. Pohler, R. Franz, J. Ramm, P. Polcik, C. Mitterer, Thin Solid Films 550 (2014) 95104.
[13]
H. Najafi, A. Karimi, P. Dessarzin, M. Morstein, Surf. Coat. Technol. 214 (2013) 46-
IP
A. Khatibi, J. Palisaitis, C. Höglund, A. Eriksson, P.O.Å. Persson, J. Jensen, J. Birch,
SC
R
[14]
T
52.
P. Eklund, L. Hultman, Thin Solid Films 519 (2011) 2426-2429.
A. Khatibi, J. Sjölen, G. Greczynski, J. Jensen, P. Eklund, L. Hultman, Acta Mater. 60
NU
[15]
[16]
MA
(2012) 6494-6507.
J. Paulitsch, R. Rachbauer, J. Ramm, P. Polcik, P.H. Mayrhofer, Vacuum 100 (2014)
D
29-32.
A. Anders, Thin Solid Films 518 (2010) 4087-4090.
[18]
D. Mattox, J. Vac. Sci. Technol. A 7 (1989) 1105.
[19]
I. Petrov, P.B. Barna, L. Hultman, J.E. Greene, J. Vac. Sci. Technol. A 21 (2003) 117-
AC
128.
CE
P
TE
[17]
[20]
R.L. Boxman, S. Goldsmith, Surf. Coat. Technol. 52 (1992) 39-50.
[21]
G. Lin, Y. Zhao, H. Guo, D. Wang, C. Dong, R. Huang, L. Wen, J. Vac. Sci. Technol.
A 22 (2004) 1218-1222.
[22]
A. Anders, Cathodic Arcs, 1st ed., Springer, New York, 2008.
[23]
A. Bendavid, P.J. Martin, E.W. Preston, Thin Solid Films 517 (2008) 494-499.
[24]
P.J. Martin, A. Bendavid, Thin Solid Films 518 (2010) 5078-5082.
- 21 -
ACCEPTED MANUSCRIPT
[25]
Z.L. Dai, Y.N. Wang, J. Appl. Phys. 92 (2002) 6428-6433.
[26]
E.V. Barnat, T.M. Lu, Phys. Rev. E: Stat. Nonlinear Soft Matter Phys. 66 (2002)
P. Kudlacek, R.F. Rumphorst, M.C.M. Van De Sanden, J. Appl. Phys. 106 (2009)
IP
[27]
T
056401.
[28]
SC
R
073303.
P. Polcik, Neue Horizoonte in der Pulvermetallurgie - Werkstoff, Produkte und
NU
Verfahren, in: H. Danninger, H. Kestler, H. Kolaska (Eds.) Pulvermetallurgie in
[29]
MA
Wissenschaft und Praxis, Hagen, Germany, 2014, 181-203 (in German).
N.P. Barradas, S. Parascandola, B.J. Sealy, R. Grötzschel, U. Kreissig, 161 (2000) 308-
Powder Diffraction File (Card 00-004-0787 for Al, Card 00-006-0694 for bcc-Cr, Card
TE
[30]
D
313.
00-046-1212 for -Al2O3, Card 00-029-0015 for Al8Cr5, Card 00-040-1242 for Al4Cr,
CE
P
Card 00-002-1202 for Al2Cr), International Centre for Powder Diffraction Data, ICDD
- JCPDS, 2007.
N.D. Chatterjee, H. Leistner, L. Terhart, K. Abraham, R. Klaska, Am. Mineral. 67
AC
[31]
(1982) 725-735.
[32]
P.H. Mayrhofer, C. Mitterer, Surf. Coat. Technol. 133-134 (2000) 131-137.
[33]
W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564-1580.
[34]
R. Franz, P. Polcik, A. Anders, IEEE Trans. Plasma Sci. 41 (2013) 1929-1937.
[35]
J.M. Schneider, A. Anders, I.G. Brown, B. Hjörvarsson, L. Hultman, Appl. Phys. Lett.
75 (1999) 612-614.
- 22 -
ACCEPTED MANUSCRIPT
[36]
J. Rosén, S. Mráz, U. Kreissig, D. Music, J.M. Schneider, Plasma Chem. Plasma
Process. 25 (2005) 303-317.
R. Franz, P. Polcik, A. Anders, Surf. Coat. Technol. 272 (2015) 309-321.
[38]
M. Pohler, R. Franz, J. Ramm, P. Polcik, C. Mitterer, Surf. Coat. Technol. 206 (2011)
IP
T
[37]
SC
R
1454-1460.
Y. Wei, C. Gong, Appl. Surf. Sci. 257 (2011) 7881-7886.
[40]
M. Kathrein, C. Michotte, M. Penoy, P. Polcik, C. Mitterer, Surf. Coat. Technol. 200
NU
[39]
[41]
MA
(2005) 1867-1871.
J. Ramm, A. Neels, B. Widrig, M. Döbeli, L. de Abreu Vieira, A. Dommann, H.
TE
D
Rudigier, Surf. Coat. Technol. 205 (2010) 1356-1361.
O. Kyrylov, D. Kurapov, J.M. Schneider, Appl. Phys. A: Mater. 80 (2005) 1657-1660.
[43]
R. Brill, F. Koch, J. Mazurelle, D. Levchuk, M. Balden, Y. Yamada-Takamura, H.
CE
P
[42]
AC
Maier, H. Bolt, Surf. Coat. Technol. 174-175 (2003) 606-610.
[44]
S. Grasser, R. Daniel, C. Mitterer, Surf. Coat. Technol. 206 (2012) 4666-4671.
[45]
B. Chapman, Glow Discharge Processes, Wiley Interscience, New York, 1980.
[46]
A.G. Fedorus, E.V. Klimenko, A.G. Naumovets, E.M. Zasimovich, I.N. Zasimovich,
Nuclear Inst. and Methods in Physics Research, B 101 (1995) 207-215.
[47]
R. Daniel, K.J. Martinschitz, J. Keckes, C. Mitterer, Acta Mater. 58 (2010) 2621-2633.
[48]
B. Alling, A. Khatibi, S.I. Simak, P. Eklund, L. Hultman, J. Vac. Sci. Technol. A 31
(2013) 030602.
- 23 -
ACCEPTED MANUSCRIPT
[49]
A. Khatibi, J. Lu, J. Jensen, P. Eklund, L. Hultman, Surf. Coat. Technol. 206 (2012)
3216-3222.
[50]
M. Schmid, G. Leonardelli, M. Sporn, E. Platzgummer, W. Hebenstreit, M. Pinczolits,
IP
L. Castaldi, D. Kurapov, A. Reiter, V. Shklover, P. Schwaller, J. Patscheider, Surf.
SC
R
[51]
Coat. Technol. 203 (2008) 545-549.
A. Picone, G. Fratesi, M. Riva, G. Bussetti, A. Calloni, A. Brambilla, M.I. Trioni, L.
NU
[52]
T
P. Varga, Phys. Rev. Lett. 82 (1999) 355-358.
Duò, F. Ciccacci, M. Finazzi, Physical Review B - Condensed Matter and Materials
[53]
MA
Physics 87 (2013) 085403.
R. Daniel, J. Keckes, I. Matko, M. Burghammer, C. Mitterer, Acta Mater. 61 (2013)
[54]
TE
D
6255-6266.
X. Pang, K. Gao, F. Luo, Y. Emirov, A.A. Levin, A.A. Volinsky, Thin Solid Films 517
[55]
CE
P
(2009) 1922-1927.
M. Tkadletz, C. Mitterer, B. Sartory, I. Letofsky-Papst, C. Czettl, C. Michotte, Surf.
[56]
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Coat. Technol. 257 (2014) 95-101.
N.P. Suh, H.C. Sin, Wear 69 (1981) 91-114.
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Table captions
Table 1: Chemical composition as measured by ERDA together with the calculated metal
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ratio [%] for coatings synthesised from different Al-Cr cathodes and with different substrate
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bias duty cycles.
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Table 1:
Al
Cr
O
metal ratio
Al/(Al+Cr)
Al0.25Cr0.75
90
11
29.5
59.5
27.2
Al0.5Cr0.5
90
22
19.8
58.2
52.6
Al0.5Cr0.5
70
21.3
20.2
58.5
51.3
Al0.7Cr0.3
90
30
11.6
58.4
72.1
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duty cycle
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coating composition [at.%]
cathode
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Figure captions
Fig. 1:
Nominal and measured substrate voltage and current obtained for an arc plasma of
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two Al0.5Cr0.5 cathodes operated with an arc current of 180 A in reactive oxygen atmosphere.
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The applied bipolar pulsed substrate bias has an amplitude of 60 V, a frequency of 25 kHz
SEM top view of the (AlxCr1-x)2O3 coatings with x = 0.5 synthesised with a substrate
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Fig. 2:
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and a duty cycle 90%.
Fig. 3:
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bias duty cycle of (a) 70%, (b) 90%, (c) 95% and (d) 100%.
SEM cross-section of the (AlxCr1-x)2O3 coatings with x = 0.5 synthesised with a
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substrate bias duty cycle of (a) 70%, (b) 90%, (c) 95% and (d) 100%. The dashed line
GIXRD patterns of the (AlxCr1-x)2O3 coatings with x = 0.5 synthesised with a
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Fig. 4:
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indicates the substrate–coating interface.
substrate bias duty cycle of 70%, 90%, 95% and 100%. The peak at ~56° 2θ in the 90%
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pattern was manually cut off due to pronounced overlapping with a reflection from the Si
substrate.
Fig. 5:
GIXRD patterns of the (AlxCr1-x)2O3 coatings with x = 0.25 and x = 0.7 synthesised
with a substrate bias duty cycle of 70 and 95%.
Fig. 6:
Total stress of the (AlxCr1-x)2O3 coatings with x = 0.25, 0.5 and 0.7 in the
as-deposited state on Si substrates as a function of the substrate bias duty cycle.
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Fig. 7:
Domain sizes of the corundum-type (AlxCrx-1)2O3 phase for the coatings with
x = 0.25 and 0.5 calculated from the corresponding XRD patterns.
Fig. 9:
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and 0.5 as a function of the substrate bias duty cycle.
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(a) Hardness and (b) Young’s modulus for the (AlxCr1-x)2O3 coatings with x = 0.25
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Fig. 8:
Two dimensional plot of the wear track for the (AlxCr1-x)2O3 coatings with (a)
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x = 0.25, (b) x = 0.5 and (c) x = 0.7 synthesised with a substrate duty cycle of 95% after
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ball-on-disc tests.
Fig. 10: Light optical micrographs of the wear scars on the alumina balls used in ball on disc
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testing against the (AlxCr1-x)2O3 coatings with x = 0.5 synthesised with a substrate bias duty
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Research highlights:
Industrial-scale cathodic arc deposition of corundum-type (AlxCr1-x)2O3 coatings

Duty cycle and Al content dependent phase composition of (AlxCr1-x)2O3 coatings

Variation of pulsed bias parameters to influence the growth conditions

High hardness and Young’s modulus at high duty cycle settings

Pulsed bias settings yield improved coating properties compared to DC bias voltage
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