Effect of Ni content on the microstructure and mechanical behaviour of CrAlNiN coatings deposited by closed field unbalanced magnetron sputtering

Surface & Coatings Technology 357 (2019) 445–455

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Effect of Ni content on the microstructure and mechanical behaviour of CrAlNiN coatings deposited by closed field unbalanced magnetron sputtering

T

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Chuhan Shaa, , Paul Munroea, Zhifeng Zhoub, Zonghan Xiec,d a

School of Materials Science and Engineering, UNSW Sydney, NSW 2052, Australia Department of Mechanical and Biomedical Engineering, City University of Hong Kong, Kowloon, Hong Kong, China c School of Mechanical Engineering, University of Adelaide, SA 5005, Australia d School of Engineering, Edith Cowan University, WA 6027, Australia b

A R T I C LE I N FO

A B S T R A C T

Keywords: Chromium aluminium nickel nitride Nickel content Magnetron sputtering Structural evolution Mechanical properties

Ni incorporation is known to impose an intriguing influence on the composition, microstructure and mechanical properties of nitride-based coatings prepared by physical vapour deposition. This work examined the effect of varying Ni concentrations, controlled by NiCr alloy target current (INiCr), on the microstructures and properties of CrAlNiN coatings deposited on to M2 tool steel substrates. These coatings were deposited by closed field unbalanced magnetron sputtering (CFUMS). The composition and structure of the as-deposited coatings were investigated by XPS, XRD, FIB, and TEM. Residual stresses were determined using the XRD-sin2ψ method. Nanoindentation tests were performed to assess mechanical properties of the CrAlNiN outer layer. At relatively low INiCr values, below 3 A, the grains within CrAlNiN layer underwent refinement, yet maintained a columnar structure. At higher INiCr values, from 3 to 5 A, the columnar grains transitioned into more equiaxed grains. Hardness values of 25–28 GPa at low INiCr values were associated with the high residual compressive stress, solid solution hardening and grain refinement.

1. Introduction Chromium nitride (CrN) coatings, among several common families of transition metal nitride (TMN) coatings, have been extensively studied and used, for example, in protecting tool steels from hostile machining conditions such as high temperatures, extreme loads, or with the use of minimal lubricants [1–5]. Studies are ongoing with respect to tailoring the mechanical and tribological properties of CrN coatings to be more adaptive in response to more demanding machining operations (e.g., high speed, dry). To date, alloying CrN with one or two elements to form ternary or quaternary compound films has demonstrated to be a promising approach in designing CrN-based multicomponent or multilayer structures with improved properties [6–11]. A series of effects, such as higher dislocation densities, alteration of lattice structure, or solid solution effects, generated through the incorporation of other elements, are believed to contribute to the higher hardness, improved fracture toughness, wear and oxidation resistance of these ternary or quaternary films. Despite these positive outcomes, there have been a number of

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studies investigating the addition of metals such as Al, Ni, V, Nb, Zr to CrN [12–15]. It has been shown that the introduction of Al into CrN increases hardness, wear and corrosion resistance due to the densification of microstructure and by promotion of covalent bonding. Barshilia et al. found that pores and surface defects contained within CrN films deposited by direct current magnetron sputtering were removed when Al was added, thus, making the substrate less susceptible to damage caused by corrosion and external loads (with a concomitant hardness increase from ~7.5 GPa to ~22.5 GPa [16]). Li et al. reported that Al additions affected the preferred orientation and compressive residual stress in nitride coatings. In some studies, the formation of AlN amorphous phases occurred at CrN grain boundaries. This led to the formation of self-toughened columnar structures that increased the strength and produced a super high hardness of 42.5 GPa in addition to excellent damage tolerance [17]. It is worth pointing out these promising properties are, however, achieved only for the metastable cubic structure (B1), while Al-rich phases prefer the hexagonal structure (wurtzite B4) which displays poorer mechanical properties. Thus, knowledge of the maximum AlN solubility in the cubic phase is

Corresponding author. E-mail address: [email protected] (C. Sha).

https://doi.org/10.1016/j.surfcoat.2018.10.052 Received 28 April 2018; Received in revised form 24 September 2018; Accepted 18 October 2018 Available online 22 October 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.

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Table 1 Deposition parameters during the preparation of the CrAlNiN coatings.

Cr adhesive layer CrN buffer layer CrAlNiN graded layer CrAlNiN top layer

Bias voltage

OEM setting

Target current

−80 V −80 V −80 V −80 V

−100% 50% 45% 45%

Cr(2) = 4.0 Cr(2) = 4.0 Cr(2) = 4.0 Cr(2) = 4.0

A A A; Al(1) = 0 → 8.0 A; NiCr(1) = 0 → setpoint (0, 0.5, 1, 2, 3, 4, 5 A) A; Al(1) = 8.0 A; NiCr(1) = setpoint (0, 0.5, 1, 2, 3, 4, 5 A)

were used in preparing these CrAlNiN coatings. The system was equipped with a rotating substrate holder with a diameter of 300 mm, which was set at a spacing of 170 mm from the targets and rotated at the speed of 10 rpm to obtain uniform deposition. The substrates were pre-treated and polished to an average surface roughness of 0.01 μm, degreased, ultrasonically cleaned and dried in nitrogen gas. Prior to deposition, the base pressure was reduced to below 2.67 × 10−4 Pa. The substrate was pre-heated to 550 °C and biased with a pulsed DC signal at a frequency of 250 kHz. An Ar plasma was used to sputter clean the substrate surface at a bias of −450 V for 30 min to ensure the removal of surface contaminants. During deposition, an Ar working gas (purity 99.999%) was introduced into the chamber at a constant flow rate of 50 sccm and kept at around 0.17 Pa. N2 gas flow (purity 99.999%) was dynamically controlled by an optical emission monitor (OEM) set at 45%. By such means, the total gas pressure was maintained at around 0.25 Pa. The bias voltage was decreased to −80 V when deposition commenced. A Cr adhesive layer (~0.2 μm) was first deposited onto the substrate. Then, a CrN buffer layer (~0.3 μm) was deposited. For each sample, the current of the two Cr targets (ICr) was set at 4.0 A and this remained unchanged for the entire coating process. Subsequently, a CrAlNiN graded layer was formed by progressively increasing the concentrations of Al and Ni via varying both the Al target current (IAl) from 0 to 8.0 A and the NiCr alloy target current (INiCr) from 0 to a defined setpoint (0, 0.5, 1, 2, 3, 4, 5 A). Finally, a CrAlNiN outer layer (1.5–1.7 μm in thickness) was deposited with IAl = 8.0 A and INiCr = setpoint (0, 0.5, 1, 2, 3, 4, 5 A). Differentiated by their Ni content, and controlled by the NiCr alloy target current, these coatings were designated as CrAlNiN-0 A, CrAlNiN-0.5 A, CrAlNiN-1 A, CrAlNiN-2 A, CrAlNiN-3 A, CrAlNiN-4 A, and CrAlNiN-5 A. Details of the deposition conditions are presented in Table 1.

important. Makino predicted the maximum mole fraction of AlN, while maintaining the cubic structure as 0.77 in Cr1-xAlxN system using a semi-empirical band parameter method [18]. Mayrhofer et al. calculated formation energies of cubic B1 and wurtzite B4 phases using a fully ab initio approach and determined the solubility limit value of 0.66 for Cr1-xAlxN [19]. The solubility limit concluded from experimental observations on the maximum AlN fraction in the cubic Cr1xAlxN system varied in the range of 0.6 and 0.8, which reasonably agrees with the simulated calculations [20,21]. In comparison with the hardness enhancements arising from Al additions, less-hard, but tougher CrN-based coatings through Ni additions have also received attention. CrN-Ni, can be represented in the generic form M1N-M2 where M1N represents the TMN, which is hard, yet relatively brittle, and M2 refers to a soft metal which enhances fracture toughness (e.g. M1 = Ti, Cr, W, Zr, V…; M2 = Ni, Cu, Fe, Y, Ag, Co, Mo) [22]. Such a combination of hard and ductile phases provides greater potential for coatings with both good strength and damagetolerance. Musil et al. first synthesized hard, yet tough, CrNiN coatings and obtained a maximum hardness of 40 GPa, where all the grains were arranged in a single orientation [23]. Similarly, Wo et al. deposited CrNiN coatings via closed field unbalanced magnetron sputtering and found that additions of 15.7–46.7 at.% Ni refined the columnar CrN grains and transformed the coating's primary deformation behaviour from intercolumnar shear sliding into deformation mediated by geometrically necessary dislocations, thus a higher load was required to induce crack formation [24]. It was found that there is a critical content for Ni above which CrNiN starts to soften significantly. Tan et al. prepared CrNiN coatings using RF reactive magnetron sputtering and found that the coating hardness first increased when the Ni content reached 2.92 at.% and then decreased when it exceeded 8.79 at.% [25]. Wang et al. added Ni into CrAlN/SiNx nanocomposites via magnetron sputtering and improved the scratch toughness by around 200% only sacrificing 18% in hardness with a 4.2 at.% Ni addition [26]. Despite all this, the pursuit of high hardness and improved toughness is still a priority in coating development. To our knowledge, little work has been performed on studying the influence of Ni additions on the microstructural evolution and mechanical properties of CrAlN coatings deposited on AISI M2 tool steel substrates by closed field unbalanced magnetron sputtering technique. The Ni addition was varied by regulating the sputtering current of a NiCr alloy target and the resultant CrAlNiN coatings were investigated in terms of their microstructure, chemical and phase compositions, hardness and reduced elastic modulus using a variety of surface characterization methods. Mechanisms interpreting how these microstructures relate to mechanical properties are discussed in relation to prior studies in this field.

2.2. Microstructural analysis and mechanical tests The surface elemental composition and chemical bonding states of the phases present were measured by X-ray photoelectron spectroscopy (XPS) (ESCALAB250Xi, Thermo Scientific, UK). The X-ray source used in the measurement was monochromated Al Kα with an energy of 1486.68 eV and power of 120 W (accelerating voltage: 13.8 kV; emission current: 8.7 mA). The pass energy used to perform a survey scan was set to 100 eV, while for region scans the range of interest was set as 20 eV. Before XPS spectral acquisition, the sample surface was ion beam cleaned with Ar+ for 1500 s to remove any surface contaminants. Thermo Scientific Avantage software was utilized to interpret the collected spectra. The spectrometer calibration was performed on the binding energies of Au 4f7 (83.96 eV), Ag 3d5 (368.21 eV), and Cu2p3 (932.62 eV). A C 1 s (284.8 eV) peak from adventitious hydrocarbons was employed as a reference to calibrate all the spectra. A Shirley background correction was applied for spectral analysis of survey scans and regions scans of Cr 2p, Al 2p, Ni 2p, and N 1s. The crystal structures and phase compositions were analysed by a PANalytical Empyrean Thin-Film X-ray diffractometer (XRD), using CuKα radiation (λ = 1.54060 Å) with an operating voltage of 45 kV and current of 40 mA. The XRD patterns were acquired over a range of 2θ values from 20° to 99° at a grazing incident angle of 3°. The residual stresses in the coatings were measured using an X'pert MRD-Philips diffractometer and analysed by the XRD-sin2ψ method [27,28].

2. Experimental 2.1. Sample synthesis CrAlNiN coatings were deposited on AISI M2 high speed tool steel substrates using a closed field unbalanced magnetron sputtering ion plating system (CFUMSIP) (Teer UDP 650/4 system, Teer Coatings Ltd., UK) configured with four targets. Two Cr targets, one Al target, and one NiCr alloy paramagnetic target (80:20 at.%), all of which were 380 mm × 175 mm × 8 mm in size and with a purity above 99.9%, 446

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Cr 2p1/2 and Cr 2p3/2 bonding states. The Cr 2p3/2 spectra located at 575.80 eV presents an asymmetric feature, which can be deconvoluted into two binding energies of 575.2 eV and 576.7 eV, representing CrN and Cr2O3, respectively [30]. Fig. 1(b) shows the Al 2p spectra with a dominant peak located at around 74 eV, indicating the presence of AleN bonding [31]. Within these spectra, peaks from a higher binding energy (74.4 eV) and lower energy (66.7 eV) corresponding to Cr 3s and Ni 3p, respectively were noted. Notably, the Ni 3p peak was not clearly evident until INiCr reached 3 A. The Ni 2p spectrum shown in Fig. 1(c) comprised Ni 2p1/2 (870.15 eV) and Ni 2p3/2 (852.9 eV), both attributed to the Ni0 state, as well as their own satellite peaks, Ni 2p1/2 (875.5 eV) and Ni 2p3/2 (859.4 eV) [26]. Intensities of those characteristic peaks from the Ni 2p spectra increase with increasing INiCr. Fig. 1(d) shows a primary peak from the N 1s spectra centred in a range of 396.6 eV and 397.2 eV, suggesting the presence of bonding of N with Cr or Al. With increasing the NiCr target current, the peak shifts towards a higher binding energy, implying a decreasing bond ratio between NeCr and NeAl. The XRD patterns of the CrAlNiN coatings deposited under different INiCr values are given in Fig. 2(a). A fcc B1 type crystalline structure can be observed in all the coatings confirmed by a set of typical diffraction peaks located at 2θ values of 37.6°, 43.7°, 63.5°, 76.2°, 80.3°, and 96.8°, correspondingly representing the (111), (200), (220), (311), (222) and (400) crystal planes for fcc CrN (ICDD 01-076-2494). As diffraction peaks from the fcc AlN structure (ICDD 00-025-1495), which are located at 37.8°, 43.9°, 63.8°, 76.7°, 80.8°, and 96.9° coincide with those from the fcc CrN structure, it is speculated that a fcc (Cr,Al)N solid solution is formed. Small peaks at 52°, ascribed to the Ni (200) peak, can be identified from coatings deposited under an INiCr of > 3 A. In addition, peak shifts can be observed as a function of INiCr. When INiCr increases from 0 to 1.0 A, all the diffraction peaks shift towards lower

The microstructures of the coatings were determined via a focused ion beam microscope (FIB, FEI xP200, USA) under an imaging current of 11 pA. Higher resolution images were obtained using a field emission transmission electron microscope (TEM, Philips CM200, Netherlands). The elemental composition and distribution in the coatings were identified using an energy dispersive X-ray spectroscopy (EDS) system (Bruker QUANTAX) interfaced to the TEM. TEM samples were sectioned by a dual beam FIB (FEI xT Nova Nanolab 200, USA) to a thickness of around 100 nm. Mechanical properties, specifically hardness, H, and elastic modulus, E, of these coatings were obtained from nanoindentation tests conducted using a Hysitron Triboindenter (TI900, USA) configured with a Berkovich indenter. Prior to nanoindentation, the instrument and the indenter tip were calibrated on a standard fused quartz sample. The load function during nanoindentation comprised linear loading from 0 to 8 mN over 10 s, a hold for 5 s at 8 mN, and then unloading to 0 mN over 10 s. A peak load of 8 mN was set to ensure the maximum penetration depth was always less than one tenth of the coating thickness to ensure no interference from the substrate in the measured hardness and modulus values. These were calculated by the OliverPharr method [29]. For each sample, at least 27 indents were made from different sites across the sample to reflect the average properties of the coatings. 3. Results 3.1. Surface chemical composition and phase analysis The regional XPS spectra for Cr 2p, Al 2p, Ni2p, and N1s from the CrAlNiN coatings deposited under various INiCr values are displayed in Fig. 1. Fig. 1(a) shows the spin doublets of the Cr 2p spectra, namely the

Fig. 1. High-resolution XPS core level spectra for (a) Cr 2p, (b) Al 2p, (c) Ni 2p, and (d) N 1s peaks. 447

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Fig. 2. (a) XRD patterns from the CrAlNiN coatings deposited with varying INiCr; (b) Texture coefficients of (111), (200), and (220) crystalline orientations in CrAlNiN coatings deposited with varying INiCr; Enlarged 2θ diffraction peaks of (c) (111), (d) (200), and (e) (220) from (a). Table 2 Crystallite size of the CrAlNiN coatings in (111) reflection.

Thkl =

INiCr (A)

0

0.5

1

2

3

4

5

Crystallite size (nm)

275

260

46.2

32

14.7

12.2

10.3

Cλ βsample cosθ

n

∑1 [Im (hkl)/ I0 (hkl)]

(2)

where Im(hkl) is the measured relative intensity of the (hkl) plane reflection, I0(hkl) is that from the identical plane retrieved from a standard reference sample (ICDD 01-076-2494), and n is the total number of reflection peaks from the samples. Based on the relationships of those crystal orientations under investigation with variations in INiCr from Fig. 2(b), the preferred orientation in CrAlNiN coatings varies from (220) when the INiCr is 0.5 A, to (111) when the current reaches 1 A, and ultimately to (200) when the current reaches 3 A. The change of preferred orientation can also be supported by the broadening or sharpening of the corresponding diffraction peaks from Fig. 2(c), (d), and (e). For the (111) peak, it appears the sharpest peak is at INiCr of 1.0 A and begins to broaden when INiCr increases. The (200) peak starts to broaden from an INiCr of 0 to 1.0 A, yet transforms to sharpen above an INiCr of 2.0 A, with apparent further sharpening at an INiCr of 3.0 and 4.0 A. The shape change of the (220) peak is similar to the (111) peak, except that it appears to be the sharpest at INiCr of 0.5 A, above which it starts to broaden.

angles, whereas when the current increases from 1.0 A to 5.0 A, those peaks shift towards higher values. This is clearly evident in the enlarged XRD spectra (Fig. 2(c), (d), and (e)). Possible reasons contributing to this will be discussed later. Average crystallite sizes of the CrAlNiN coatings (Table 2) were estimated using the full-width-at-half-maximum (FWHM) of (111) reflection peaks according to the Scherrer formula [32]:

D=

Im (hkl)/ I0 (hkl) 1 n

(1)

where C is shape factor, equal to 0.89, λ = 1.5406 Å is the wavelength of the CuKα radiation, βsample = βobs 2 − βins 2 , βobs is the FWHM in the observed (111) reflection while βins is the correction factor for instrument broadening, and θ is the Bragg diffraction angle for (111) peak. Standard relationships for Gaussian and Lorentzian line shapes were used to determine the broadening associated with the micro-strain and crystallite size. In brief, the variables of the Fourier transforms of the (111) diffraction line were computed and subjected to a Warren-Averbach analysis [33]. Table 2 shows that the crystallite size of CrAlNiN coatings decrease notably with increasing INiCr, from 200 to 300 nm at low INiCr (0–0.5 A), to 30–50 nm at medium INiCr (1–2 A), and to tens of nm at high INiCr (3–5 A). The texture coefficients Thkl for three major crystalline orientations, (111), (200), and (220), from fcc CrN were calculated to determine the change of preferred orientations according to the equation [34]:

3.2. Microstructural analysis FIB cross-sectional images of the CrAlNiN coatings deposited under various NiCr target currents are given in Fig. 3. Three distinct layers are visible in these coatings, which from the bottom up are the Cr adhesive layer, the CrN buffer layer, and the CrAlNiN outer layer, respectively. The obtained graded structures are consistent with the coating design presented in Table 1, except that the graded CrAlNiN layer is too thin to be clearly resolved via FIB imaging. A noticeable transition of grain refinement is apparent within the outer CrAlNiN layer with increasing INiCr, from clearly delineated columnar grains at the lowest INiCr, to finer, yet well-defined grains, under currents of 0.5, 1, 2, 3 A, to more equiaxed grains at currents of 4 and 5 A. This agrees with the crystallite 448

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Fig. 3. FIB cross-sectional images of CrAlNiN coatings deposited with (a) INiCr = 0 A; (b) INiCr = 0.5 A; (c) INiCr = 1 A; (d) INiCr = 2 A; (e) INiCr = 3 A; (f) INiCr = 4 A; (g) INiCr = 5 A. scale bar: 1 μm.

thin films deposited by the magnetron sputtering, which display nonequilibrium phase compositions [35,36]. Fig. 6 shows the elemental distributions of Cr, Al, Ni, N, and Fe obtained by EDS elemental mapping the coating deposited under an INiCr of 4 A. The graded microstructure of three distinct layers is clearly evident. The Cr adhesive layer is confirmed by its highest Cr content with no other elements detected. The CrN buffer layer exhibits relatively higher Cr and N contents than the outer layer. Al and Ni are only present within the CrAlNiN outer layer, which appears to be less rich in Cr and N. Similar mapping data were obtained from the other samples, but are not presented here for the sake of brevity. In short, the Ni content within the CrAlNiN outer layer increases, as expected, with higher INiCr values.

sizes measured by the Scherrer formula. The microstructures of the CrN buffer layers and Cr adhesive layers remain almost unchanged across this range of samples as coarse columnar grains on account of the identical deposition parameters employed for all the coatings. An interesting observation is that grains within the Cr adhesive layer exhibit similar texture with, and appear connected to, the grains in the M2 steel substrate. It is speculated there is good bonding between the coating and the substrate. TEM images of three selected CrAlNiN coatings deposited under various NiCr target currents are presented in Fig. 4. In addition to the three distinct layers with clearly defined interfaces, the grain size decreases within the outer CrAlNiN layer at higher INiCr values. This is evidently demonstrated from the bright-field and dark-field images and is in agreement with the FIB observations. The outer layer thickness values which are broadly equivalent to the lengths of CrAlNiN columnar grains for those coatings remain nearly the same (around 1.6 μm) as the deposition time for the outer layers was reduced for the higher target currents, so as to keep the overall coating thicknesses approximately constant. The grain refinement is mainly reflected by the decreasing width of those grains from an average size of around 200 nm in the coating deposited under an INiCr of 0 A (Fig. 4(a)), to around 100 nm in that for 1 A (Fig. 4(d)), and then to ~70 nm where INiCr is 3 A (Fig. 4(h)). The size reduction broadly agrees with the results measured by the Scherrer formula. Moreover, the uniformity of the grain sizes is more apparent for the finer grains at higher currents. Selected area electron diffraction (SAED) patterns acquired from the outer layer area marked by yellow dashed circles from Fig. 4(a), (d), and (g) are given in Fig. 4(c), (f), and (i). An enhanced continuity in those diffraction rings with higher INiCr values can be observed, confirming the observation of grain refinement and an overall reduction in texture. In addition, observable rings from all three patterns comprise diffraction reflections from the (111), (200), (220), (311), and (222) crystal planes for CrN, which is consistent with the XRD results. Fig. 5 gives compositional data, in atomic percent, for the relative concentrations of Cr, Al, Ni, and N obtained from the CrAlNiN outer layers from TEM-EDS point analysis. The data presented is the average of data from three points. With increasing INiCr from 0 to 5.0 A, there is an increase in the atomic percent of Ni from 0 to 31 at.%, as expected, and decreases in Al from 8 at.% to 4 at.%, Cr from 45 at.% to 31 at.%, and N from 47 at.% to 34 at.%. These changes in composition are consistent with the presence of a Ni-rich phase consistent with the XRD data. Assuming the nickel hardly forms any nitrides, the non-stoichiometry of CrAlN solid solution, presented by an excess of Cr and Al in comparison to the anionic lattice of nitrides, is normally observed in

3.3. Residual stress and mechanical properties The residual stresses in the direction of ϕ = 90° were determined using the sin2ψ method. The angles of ϕ and ψ are used to identify the direction of the diffraction vector in the residual stress measurement, where ϕ refers to the angle between a reference axis in the plane of the sample and the projected lattice plane normal while ψ refers to the inclination angle of the sample surface normal with respect to the diffraction vector [37,38]. The CrN (111) diffraction peak at 2θ = 37.4° was selected for stress measurement, due to its lack of overlap with diffraction peaks from the Fe substrate. The samples were tilted to 10 different ψ values, while at each tilt the profile of (111) peak was recorded. Fig. 7(a) displays a Gaussian fitted diffraction peak of CrN (111) at ψ = 53° from the CrAlNiN coating deposited with an INiCr of 0.5 A. Fig. 7(b) displays the least square line plotted from the raw data, from which the line slope is negative, indicating the residual stress in CrAlNiN-0.5 A is compressive. The elastic constants used to derive the stress values are set as those of pure CrN, which are S1 [10−6 MPa−1] = −1, ½ S2 [10−6 MPa−1] = 6. The residual stresses in coatings deposited with various INiCr are shown in Fig. 7(c) as well as in Table 3. CrAlNiN-0 A has the lowest compressive stress of −2.82 GPa. When INiCr increases to 0.5 A and 1.0 A, the compressive stress increases to −4.67 GPa and −4.99 GPa, respectively. As INiCr increases to 5.0 A, the compressive stress decreases steadily. The sin2(ψ0) value corresponding to the stress-free direction is 0.4. Considering these coatings are deposited under energetic conditions, the lattice parameter obtained at sin2(ψ0) = 0.4 corresponds to an expanded lattice parameter due to the defect-containing layer, so a triaxial stress mode is employed to obtain a true ‘stress-free’ and ‘defect-free’ lattice parameter in each coating, as shown in Fig. 7(d) [38,39]. The stress-free lattice parameter 449

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Fig. 4. TEM bright-field, dark-field images and SAED pattern from the outer layer of the coating deposited with INiCr = 0 A: (a)–(c); INiCr = 1 A: (d)–(f); and INiCr = 3 A: (g)–(i).

first decreased slightly from 0.4139 nm in CrAlNiN-0 A to 0.4130 nm in CrAlNiN-2 A, and then more sharply from 0.4111 nm in CrAlNiN-3 A to 0.4056 nm in CrAlNiN-5 A. Representative load-displacement curves for each coating obtained from nanoindentation testing at a peak load of 8 mN are given in Fig. 8. The maximum penetration depths, hmax, are all below one tenth of the thickness of CrAlNiN outer layers, so exclude any influence from the buffer and adhesive layers or the steel substrate. Broadly, the maximum penetration depth of each coating represents its resistance against deformation when an external load is applied. In such cases, the resistance is quantified by the hardness value, so there is an inverse relationship between the penetration depth and the hardness [40]. Fig. 9(a) demonstrates the calculated hardness, H, and reduced elastic modulus, Er, with respective error bars describing the standard deviation from the mean value, of those coatings as a function of the applied NiCr current. In general, the hardness and elastic modulus of CrAlNiN coatings increased slightly at low Ni content, but reduced significantly when the Ni content exceeded 7 at.%. The mechanical properties of each CrAlNiN coating deposited under various INiCr are summarized in Table 3. When there is no Ni, the CrAlN coating exhibits a good hardness of 25.6 GPa. When INiCr increases to 0.5 A, 1 A, and 2 A, the hardness increases by 4%, 9%, and 6%, to 26.7 GPa, 28.0 GPa, and 27.2 GPa, respectively. However, when INiCr reaches 3 A and above, the hardness decreases sharply to 20.7 GPa, and below 15 GPa when INiCr is above 4 A. There is

Fig. 5. Elemental compositions of Cr, Al, Ni, and N obtained from the CrAlNiN outer layers detected by TEM-EDS point analysis.

450

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Fig. 6. TEM/EDS elemental mapping of the coating deposited with INiCr = 4.0 A: distribution of Cr, Al, Ni, N and Fe; scale bar: 700 nm.

maximum values were obtained for CrAlNiN-1 A. Taken together, CrAlNiN-1 A has the highest hardness and reduced elastic modulus of 28.0 GPa and 302.6 GPa, and thus likely to be most resistant to plastic deformation and wear.

a similar trend for values of the elastic modulus. The highest hardness and elastic modulus values were achieved in CrAlNiN-1A. Although a high hardness is normally viewed as the primary property in defining high wear and damage resistance, an integration of high hardness and low elastic modulus can be applied to characterize wear behaviour. Fig. 9(b) gives values of H/Er and H3/Er2 for these coatings, indicating the elastic strain to failure and ability to resist plastic deformation [41,42]. The effects of H/Er and H3/Er2 values for various NiCr target currents mirror those values for H and Er, in which

Fig. 7. (a) Gaussian fitted CrN (111) diffraction peak of CrAlNiN-0.5 A at ψ = 53°; (b) Least square lines plotted to calculate residual stress from slope; (c) Residual stress in each coating with various INiCr; (d) Stress-free lattice parameter in each coating as a function of INiCr. 451

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Table 3 Mechanical properties of the CrAlNiN coatings deposited under different INiCr. Sample

Residual stress (GPa)

H (GPa)

CrAlNiN-0 A CrAlNiN-0.5 A CrAlNiN-1 A CrAlNiN-2 A CrAlNiN-3 A CrAlNiN-4 A CrAlNiN-5 A

−2.82 −4.67 −4.99 −4.35 −4.05 −3.36 −3.13

25.6 26.7 28.0 27.2 20.7 14.8 14.3

± ± ± ± ± ± ±

0.23 0.22 0.30 0.36 0.25 0.27 0.34

± ± ± ± ± ± ±

Er (GPa) 2.9 4.4 4.0 3.1 4.1 2.1 1.7

297.0 304.2 302.6 291.5 267.2 242.9 241.0

± ± ± ± ± ± ±

H3/Er2 (GPa)

H/Er 22.8 31.2 22.7 20.5 33.3 16.7 15.9

0.0863 0.0877 0.0924 0.0934 0.0776 0.0607 0.0592

± ± ± ± ± ± ±

0.0032 0.0039 0.0049 0.0045 0.0029 0.0023 0.0030

0.191 0.205 0.239 0.238 0.125 0.054 0.050

± ± ± ± ± ± ±

0.022 0.025 0.027 0.027 0.013 0.008 0.005

4.2. Phase transformation and texture coefficient of CrAlNiN coatings The XRD patterns of CrAlNiN coatings confirm the presence of fcc B1 NaCl type crystalline structure in all the coatings independent of Ni content. It is known that in the Cr-Al-N system, N tends to bond with Al (ΔH° = −319.9 kJ mol−1) more strongly than with Cr (ΔH° = −216 kJ mol−1) due to its lower formation enthalpy [43,44]. The formation of a discrete AlN phase is possible in CrAlNiN coatings. It is well-known that a B4-structured wurtzite type (hexagonal ZnS) AlN is thermodynamically more stable than the metastable B1 cubic-phase [45]. However, no peaks consistent with the presence of wurtzite-AlN could be detected in the XRD pattern, indicating that only a single fcc CrAlN solid solution is formed. Despite that Cr1-xAlxN is prone to transform into a B4-structured wurtize type phase which has an inferior hardness, such a transition does not occur when x value is below 0.6 [46]. In the present work, the atomic ratios of Al/(Cr + Al) among all the samples are consistent with such a condition (x ≤ 15 at. %), which accounts for the absence of B4 wurtize AlN and the formation of single fcc CrAlN phase in these coatings. As an increasing amount of Ni is present in the Cr-Al-N system, as determined by higher values of INiCr, a slight peak at 2θ = 52°was noted only when the current exceeded 4 A. This was attributed to the (200) plane for fcc Ni. The formation of a metallic Ni-based phase within the CreN system, instead of NieN bonding in these Ni-rich coatings, is because Ni has a poor affinity to nitrogen, making it hard to form stable nitrides [47]. When INiCr is below 3 A, no diffraction peaks from nickelrich phases were present in any of these CrAlNiN coatings, presumably because of either the amorphous nature of the Ni-rich phases, as suggested by Tan et al., or low volume fractions of Ni that are below the sensitivity for XRD measurement [25]. When either Ni or Al is incorporated, at low concentrations, into the CreN system, these elements substitute for Cr and cause a change in the size of the unit cell. For either case, an overall lattice shrinkage is expected within the crystalline structure as both Ni (Ra = 0.1246 nm) and Al (Ra = 0.1210 nm) atoms are smaller than Cr (Ra = 0.1280 nm). The

Fig. 8. Representative load-displacement curves for coatings deposited at various INiCr.

4. Discussion 4.1. Elemental composition of CrAlNiN coatings The elemental compositions of Cr, Al, Ni, and N in the CrAlNiN outer layers quantified from TEM/EDS (Fig. 5) show that increasing NiCr target current, INiCr, increases the Ni concentration, but reduces the Al and Cr concentrations. The atomic ratios between Cr and Al, however, increase with increasing INiCr. Note, the N content also decreases since the N bonds with Cr or Al rather than bond with Ni which increases in its concentration. Decreasing intensities of Cr 2p, Al 2p, and N1 s peaks (Fig. 1(a), (b), and (d)) with increasing INiCr indicating decreasing CreN and AleN bond fractions, are thus ascribed to lesser Cr and Al concentrations, and greater Ni contents.

Fig. 9. Mechanical properties of the CrAlNiN coatings as a function of INiCr: (a) hardness, H and reduced elastic modulus, Er; (b) H/Er and H3/Er2. 452

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of loose, voided structures (zone 1), while favouring highly dense (zone 3) and dense columnar (zone 2) structures. Thornton introduced the zone T as a broad transition region between zone 1 and 2, which is composed of V-shaped grains [56]. The CrN buffer layers among different coatings present a similar zone-T type structure. The CrAlNiN graded layer was deposited on the CrN buffer layer via gradually increasing both the Al and NiCr target currents, IAl from zero to 8 A, and INiCr from zero to the set value used for depositing CrAlNiN outer layer, which formed a barely visible intermediate zone made up of very fine grains with graded Al and Ni concentrations. When INiCr reaches and is maintained at its set value (0, 0.5, 1, 2, 3, 4, 5 A), IAl is kept constant at 8 A for deposition of the CrAlNiN outer layer. The graded layers with multiple interfaces can help structurally strengthen the coating and the adhesion to the substrate. The grains within the CrAlNiN outer layers, with increasing Ni contents, present a transformation from a zone-2 type structure (INiCr = 0, 0.5, 1, 2 A) to a zone-3 type structure (INiCr = 3, 4, 5 A). Musil and Vlcek suggested impurities and additives incorporated in single phase films interfere with grain growth and activate grain re-nucleation, in which case the zone-2 type columnar structure is transformed into a zone-3 type dense fine-grained one [57]. In this work, it is suggested that Ni disturbs the columnar growth of CrN grains which undergo structural densification due to the high mobility and immiscibility of Ni. Such microstructural evolution induced by the Ni addition might also work on the reduction of compressive stress as shown in Fig. 7(c). According to the elemental composition from Fig. 5, the Al concentrations remain as 8 at.%, while the Ni concentrations are below 3 at.% in coatings deposited with INiCr below 1 A. It is speculated that the lattice distortion induced by the CrAlN solid solution prevails in these coatings, which causes insufficient release of strain energy during coating growth. After INiCr exceeds 1 A, the Ni concentration increases significantly from 7 to 31 at.%, while the Al concentration drops from 7 to 4 at.%. It is believed the volume fraction of CrAlN solid solution decreases, while the surrounding Ni metallic phase increases. As suggested above, this composition could be achieved through repeated nucleation of CrAlN crystallites prompted by the Ni segregation through an atom-by-atom building process, which could further release the strain energy due to the discontinuous distribution of CrAlN crystals, and thus reduce the compressive stress. A similar model was also presented for TiN-Ni nanocomposite coatings deposited by the ion beam sputtering [58].

stress-free lattice parameter as a function of the Ni concentration given in Fig. 7(d) decreases continually for this reason. This would shift 2θ diffraction angles towards higher values. Moreover, as those coatings are deposited using closed field unbalanced magnetron sputtering technology, high levels of residual compressive stress are commonly induced during film growth caused by accumulative vacancies, interstitial atoms, and dislocations arising from energetic ion bombardment [48,49]. It is generally agreed that the presence of compressive stresses in PVD-processed films increases d-spacing values, which leads to 2θ diffraction angles shifting to lower values [50]. Lin et al. compared the structure and properties of CrN coatings deposited via various magnetron sputtering techniques and confirmed that CrN deposited via pulsed DC magnetron sputtering had the highest compressive residual stress, whose diffraction peaks via XRD measurements also shifted to the lower angles [51]. In the current work, both effects, namely smaller atom substitution and compressive residual stress introduction, need to be considered together to interpret the relative shifts in the diffraction peaks. According to Fig. 2(a), (c), and (d), diffraction peaks of the CrAlNiN coating with INiCr of 1 A are located at the lowest 2θ angles, which suggests that residual compressive stress effects are more dominant. It is also anticipated that diffraction peaks may shift towards higher 2θ angles under increasing INiCr from 1 A due to the greater role played by lattice parameter reduction when increasing concentrations of Ni atoms are substituted into the CrN lattice. This agrees with the result of residual stress measurement, that the CrAlNiN coatings deposited with INiCr of 1.0 A contains the highest compressive stress of −4.99 GPa, and the stress is reduced with higher INiCr values. The variation of texture coefficient with higher INiCr from Fig. 2(b) maintains consistency with the contrasting brightness of diffraction rings from Fig. 4(c), (f), and (i). When there is no Ni addition, CrAlN exhibits a randomly arranged grain structure with almost identical texture coefficients for (111), (200), and (220). With higher Ni concentrations, more textured grains were achieved with varying preferred orientations, that is (220) for an INiCr of 0.5, (111) in INiCr of 1 and 2 A, and (200) in INiCr above 2 A. The INiCr values corresponding to the sharpest (111), (200), and (220) diffraction peak are 1.0 A, 4.0 A, and 0.5 A, respectively, as given in Fig. 2(c), (d), and (e), which can be used to indicate the preferred orientation in any individual coating. Similar relationships between the preferred orientation and NiCr target current values have also been seen in other studies [24,52]. Zhang et al. suggested that the replacement of Cr by Ni might promote adatom mobility, which is mainly associated with competitive growth mechanism, on the (200) crystal plane, or inhibition of growth on the (111) plane [52]. Li et al. reported that Al incorporation in CrAlN coatings, with a decrease in the (Cr/Al) atomic ratios, weakened (200) preferred orientation and promoted (111) and (220) peak [17]. According to the variation in elemental compositions from Fig. 5, increasing INiCr leads to both higher Ni concentrations and (Cr/Al) atomic ratios, which based on the aforementioned reports will concurrently promote crystal growth on the (200) plane and suppress it on the (111) plane. As expected, the data shown in Fig. 2(b) suggests that when INiCr exceeds 2 A, (200) becomes the preferred orientation while (111) has the lowest texture coefficient.

4.4. Mechanical properties of CrAlNiN coatings The hardness (H) and reduced elastic modulus (Er) values of CrAlNiN coatings with varying Ni contents are summarized in Fig. 9(a). As noted in the results section, the overall trend is that small Ni additions lead to large increases in both H and Er reaching a peak at INiCr = 1 A, but higher Ni additions lead to significant softening. As the sole processing variable in the present work is INiCr, the phase compositions, microstructures, and residual stress formed under various INiCr are significant in determining H and Er. From Fig. 9(a), coatings deposited under INiCr no higher than 2 A all exhibit high hardness, possibly associated with solid solution hardening by the presence of the fcc CrAlN phase, and Hall-Petch effects derived from grain refinement within CrAlNiN layers, however maintaining zone 2 dense columnar structures [24,59–62]. In particular, the highest hardness was measured for the coating deposited under INiCr of 1 A. This was largely attributed to its high compressive residual stress according to the XRD and residual stress measurement results [15]. When INiCr is > 2 A, the hardness decrease is believed to result from the excess Ni atoms accumulating as a cluster of Ni-based metallic phase within the CrN matrix, as well as a decreasing level of compressive residual stress, which together softened the coatings. Although under such conditions the grain size is further decreased and a solid solution hardening is expected to derive from the fcc CrAlN phase, they might not be sufficient to counteract the other effects that are more dominant. From Fig. 9(b), similar trends in

4.3. Microstructures of CrAlNiN coatings The microstructures of the graded layers were characterized by FIB (Fig. 3), TEM (Fig. 4) and EDS elemental mapping (Fig. 6). The Cr adhesive layer at the interface between the steel substrate and the outer coating was deposited in order to lower the thermal mismatch while enhance the adhesion on the substrate, which is common in many CrNbased coatings [53]. A structure zone model (SZM) relating to deposition via closed-field unbalanced magnetron sputtering system (CFUMS) was developed by Kelly and Arnell based on the classic structure zone model (MD) by Movchan and Demchishin [54,55]. According to this model, it is proposed that the CFUMS process suppresses the formation 453

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H/Er and H3/Er2 as a function of INiCr with that of H and E are expected, as high coating hardness normally provide more resistance to plastic deformation, which predicts desired wear resistance for CrAlNiN coatings deposited under INiCr < 3 A.

[12]

[13]

5. Conclusions [14]

In order to study the effect of Ni additions, CrAlNiN coatings were deposited onto AISI M2 steel substrates in a Cr/CrN/CrAlNiN graded structure using a closed field unbalanced magnetron sputtering ion plating technique with various NiCr alloy target currents (INiCr). The relationships between INiCr and the phase composition, microstructures, residual stress, and mechanical properties of CrAlNiN outer layers were investigated via XRD, the XRD-sin2ψ method, FIB, TEM, and nanoindentation tests. TEM/EDS results showed that increasing INiCr increased Ni concentration and decreased Al and Cr concentrations, consistent with the XRD results. The XRD results revealed that CrAlNiN coatings were composed of (Cr,Al)N fcc single phase and a Ni-based metallic phase. No AlN exhibiting a wurtzite structure or nickel nitride phase was detected in the coatings. FIB and TEM analysis showed that there was a grain refinement effect in the CrAlNiN outer layer with increasing INiCr values transitioning the layer from zone 2-type to zone 3-type structure. Compressive residual stresses were detected in these coatings. Nanoindentation tests demonstrated that high hardness (25–28 GPa) obtained at low INiCr (0–2 A) values was due to high level of compressive stress, solid solution hardening from the presence of the fcc CrAlN phase, and a refined columnar grain size, while the relatively low hardness (14–21 GPa) at high INiCr (3–5 A) values was attributed to decreased compressive stress, and the presence of a soft Ni metallic phase.

[15]

[16]

[17]

[18] [19] [20]

[21]

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[25] [26]

Acknowledgement

[27]

This work is financially supported under the Australian Research Council Discovery Project (DP150102417). The authors appreciate assistance from Dr. Bill Joe of the School of Materials Science and Engineering during the nanoindentation tests. Thanks also go to Dr. Yu Wang and Dr. Bill Gong of the SSEA Unit, UNSW for their support with the XRD and XPS measurement, along with Dr. Charlie Kong and the staff at EMU, UNSW for their support with the microscopy characterization.

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