Microstructure, mechanical, and corrosion properties of Zr 1-x Cr x B y diboride alloy thin films grown by hybrid high power impulse/DC magnetron co-sputtering

We study microstructure, mechanical, and corrosion properties of Zr 1-x Cr x B y coatings deposited by hybrid high-power impulse/DC magnetron co-sputtering (CrB 2 -HiPIMS/ZrB 2 -DCMS). Cr/(Zr + Cr) ratio, x, increases from 0.13 to 0.9, while B/(Zr + Cr) ratio, y, decreases from 2.92 to 1.81. As reference, ZrB 2.18 and CrB 1.81 layers are grown at 4000 W DCMS. ZrB 2.18 and CrB 1.81 columns are continual from near substrate toward the surface with open column boundaries. We find that the critical growth parameter to achieve dense films is the ratio of Cr + - dominated ion flux and the (Zr + B) neutral flux from the ZrB 2 target. Thus, the alloys are categorized in two groups: films with x < 0.32 (low Cr + /(Zr + B) ratios) that have continuous columnar growth, rough surfaces, and open column boundaries, and films with x ≥ 0.32 (high Cr + /(Zr + B) ratios) that Cr + -dominated ion fluxes are sufficient to interrupt continuous columns, resulting in smooth surface and dense fine-grain microstructure. The pulsed metal-ion irradiation is more effective in film densification than continuous Ar + bombardment. Dense Zr 0.46 Cr 0.54 B 2.40 and Zr 0.10 Cr 0.90 B 1.81 alloys are hard ( > 30 GPa) and almost stress-free with relative nano-indentation toughness of 1.3 MPa √ m and 2.3 MPa √ m, respectively, and remarkedly low corrosion rates (~1.0 × 10 -6 mA/cm 2 for Zr 0.46 Cr 0.54 B 2.40 and ~ 2.1 × 10 -6 mA/cm 2 for Zr 0.10 Cr 0.90 B 1.81 ).


Introduction
Zirconium diboride (ZrB 2 ) with high melting point (~3245 • C) [1] and hardness (ranging from 19.3 to 46.0 GPa) [2][3][4][5][6][7][8], classified as ultrahigh temperature ceramics [9], is a promising material for long-term service in extreme conditions including temperatures above 2000 • C, hydrostatic pressure, drastic chemical reactivity, mechanical stress, and very high levels of heat and radiation gradients [9][10][11].It typically crystallizes in a hexagonal AlB 2 -type structure (P6/mmm,  where B atoms form graphite-like honeycomb sheets sitting between hexagonal-close-packed Zr layers [12].The lattice parameters of ZrB 2 are 3.17 Å and 3.53 Å in the in-plane and out-of-plane directions, respectively [13].Strong covalent bonding between Zr and B atoms as well as within the honeycomb B sheets provide their high melting point, stiffness, and hardness [14].In addition, metallic bonding within the Zr layers results in good thermal and electrical conductivities [12].These ceramic and metallic properties make ZrB 2 a good candidate for many applications.For instance, they are suitable for aerospace applications such as in rocket components, atmospheric reentries, jet engine turbines, propulsion systems, and sharp leading edges in hypersonic vehicles [9,11,[15][16][17]. Refractory diboride coatings grown by magnetron sputtering have recently received increasing attention as a new class of hard ceramic coatings to replace transition-metal (TM) nitrides in many applications [4][5][6][18][19][20][21][22][23][24][25][26][27].However, the use of sputter-deposited TM diboride coatings are very restricted because of several critical issues.Contrary to TM nitrides, ZrB 2 are line-compound, and deviations from stoichiometry can lead to the formation of second phases [2].Sputter-deposited ZrB y thin films are typically overstoichiometric with B/Zr (y) ratios > 2 (containing excess B) that can influence their phase stability at elevated temperatures [4,6].This issue can be effectively resolved by employing high-power impulse magnetron sputtering (HiPIMS) that exploits differences in the ionization probabilities between TM and B atoms and gas rarefaction effects in order to guide ion fluxes toward substrates [21].
The other critical issue of ZrB 2 thin films is their inherent brittleness [28].We recently proposed a strategy to increase both hardness and toughness of ZrB 2 layers by alloying with Ta [4].The Zr 1-x Ta x B y alloy films were grown using a hybrid high power impulse/DC magnetron cosputtering (Ta-HiPIMS/ZrB 2 -DCMS) scheme.Hardness increases from ~ 35.0 GPa for ZrB 2.4 to ~ 42.0 GPa for Zr 0.7 Ta 0.3 B 1.5 , with an increase in nanoindentation toughness from ~ 4.0 to ~ 5.2 MPa√m.This simultaneous enhancement in mechanical properties is due to a transition in their nanostructure from B-rich to Ta-rich column boundaries and the formation of a self-organized columnar core/shell nanostructure [4,19].In addition, TMB 2 suffer from poor high-temperature oxidation resistance, which is a critical property for many aggressive environments [29,30].They oxidize at temperatures < ~450 • C with oxidation products consisting of TMO 2 and B 2 O 3 [31].Contrary to bulk diborides, the B 2 O 3 phase evaporates rapidly at temperatures > ~400 • C from overstoichiometric TMB 2 thin films.This leads to the formation of a highly porous oxide layer with no oxidation protection [20,32,33].As also proved for TiN-based coatings, alloying diboride thin films with Al significantly improves their oxidation resistance properties [20].
The ever-increasing demand for enhanced coating properties such as high hardness, good toughness, low wear rates, and increased oxidation and corrosion resistance motivates the search for multifunctional materials.We previously reported on the growth of ZrB 2 -rich Zr 1-x Cr x B y thin films using hybrid Cr-HiPIMS/ZrB 2 -DCMS co-sputtering.Decreasing the ZrB 2 target power operated at DCMS mode (P ZrB2 ) leads to an increase in the Cr metal fraction x, Cr/(Zr + Cr), up to x = 0.44, while a decrease in the B/(Zr + Cr) ratio y from 2.18 to 1.11 due to both Cr addition and preferential B resputtering.The alloy films with x = 0.23 and 0.29 showed the highest hardness values, but their cubecorner nanoindentation impressions revealed brittle behavior.Increasing the Cr concentration (x > 0.29) resulted in enhanced ductility, wear, and corrosion resistance.However, it also caused a significant decrease of ~ 12 GPa in hardness.In addition, the alloys with x ≥ 0.36 were amorphous [18].
To retain the diboride lattice structure and thus avoid deterioration of mechanical properties, both caused by B deficiency, here, we use a CrB 2 target instead of Cr to grow the Zr 1-x Cr x B y thin films by hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS co-sputtering with metal-ion-synchronized substrate potential, and study the effect of CrB y addition on their microstructure, mechanical, and corrosion properties.We demonstrate that the hybrid growth scheme enables to cover broad compositional range y = B/(Zr + Cr) from 2.90 to 1.81 for x = Cr/(Zr + Cr) varied from 0 to 1, which was hitherto not achievable by conventional DCMS processing.All alloy coatings show hexagonal crystal structure (solid solution) and have nanoindentation hardness values ranging from ~ 27.0 to ~ 32.5 GPa.Our results reveal that the hybrid co-sputtering technique that provides intense metal ion fluxes, resulting in dense layers with smooth surfaces, is favorable to DCMS, as long as the metal ion-flux density is sufficiently high with respect to the neutral flux from the DCMS-operated source.

Material and methods
Zr 1-x Cr x B y thin films, in which x = Cr/(Zr + Cr) and y = B/(Zr + Cr), are grown in a CC800/9 CemeCon AG sputtering system (featuring integrated HiPIMS and DCMS power supplies) using cast rectangular 8.8 × 50 cm 2 stoichiometric ZrB 2 and CrB 2 targets.Al 2 O 3 (0001) and Si(001) are used as substrates.The Si(001) substrates are used for structural studies, while the Al 2 O 3 (0001) substrates are used for nanoindentation and residual stress measurements.The distance between the targets and substrates is 20 cm.The substrate temperature before starting the film growth is fixed at ~ 475 • C. The sputtering-system base pressure is 3.0 × 10 -6 Torr (0.4 mPa).The films are deposited with an Ar pressure P Ar = 3 mTorr (0.4 Pa).
The Zr 1-x Cr x B y alloy films are grown using a hybrid CrB 2 -HiPIMS/ ZrB 2 -DCMS scheme with metal-ion-synchronized substrate potential [34].The ZrB 2 target is operated in DCMS mode, while the CrB 2 target is operated in HiPIMS mode, providing effective ion subplantation into DCMS-grown layers [35].In order to cover a wide compositional range on the metal lattice, two series of Zr 1-x Cr x B y alloy films are grown.First, at a constant CrB 2 -HiPIMS-target average power (here, "average" refers to the entire cycle, i.e., including the pulse off time), frequency, and pulse width (950 W, 100 Hz, and 50 µs), the Cr metal fraction, x, increases from 0.13 to 0.17 to 0.24 by decreasing P ZrB2 from 4000 W to 3000 W to 2000 W, respectively.In addition, while maintaining the average power, pulse width, and frequency applied to the HiPIMS CrB target constant at 3400 W, 50 µs, and 300 Hz, respectively, the Cr metal fraction (y) increases from 0.32 to 0.90 by decreasing P ZrB2 from 4000 W to 400 W. For both HiPIMS deposition conditions, the average peak CrB 2 -target current density is ~ 1.0 A/cm 2 to provide high ionization of the sputtered flux [36].A substrate potential of − 100 V synchronized to the metal-rich portion of each HiPIMS pulse is employed (a 100-µs bias pulse with a 35-µs delay with respect to each HiPIMS pulse) [37].In addition, ZrB y and CrB y reference thin films are deposited by DCMS with a target power of 4000 W at two different substrate bias conditions; (i) a floating substrate potential (-10 V) and (ii) a DC substrate potential (-100 V).Monte Carlo TRIM simulations are performed for better understanding the effects of metal ion irradiation.
Film elemental compositions are determined using time-of-flight elastic recoil detection analyses (ToF-ERDA).The ToF-ERDA measurements are carried out in a tandem accelerator with a 36 MeV 127 I 8+ probe beam incident at 67.5 • with respect to the sample surface normal.Recoils are detected at 45 • .A Philips XPert X-ray diffractometer operated with a Cu-K α source (λ = 0.15406 nm) is used to conduct X-ray diffraction (XRD) θ-2θ scans to obtain the crystal structure as well as orientations.A Zeiss LEO1550 scanning electron microscope (SEM) is employed to determine the layers' cross-sectional morphologies and thicknesses.Cross-sectional and plan-view transmission electron microscopy (TEM) analyses are carried out in a FEI Tecnai G 2 TF 20 UT instrument operated at 200 kV.
The residual stresses σ f of the films are determined from substrate wafer curvatures using the modified Stoney equation (more details are given in References [38,39]).Nanoindentation analyses are carried out in an Ultra-Micro Indentation System equipped with a Berkovich diamond tip.The values of hardness H and elastic modulus E are determined by indenting the layers grown on Al 2 O 3 (0001) using a load of mN.E values are calculated from reduced elastic moduli using the diamond indenter's elastic modulus (1141 GPa) and Poisson's ratio ν = 0.07.The ν values of Zr 1-x Cr x B y required for obtaining E are unknown and estimated based upon a linear interpolation between the Poisson ratio of ZrB 2 (0.13 [40,41]) and that of CrB 2 (0.21 [41]).
The films grown on Al 2 O 3 (0001) are also indented by a diamond cube-corner tip in order to determine their relative nanoindentation toughness K c .The cube-corner tip is a three-sided indenter similar to Berkovich tips, but sharper.Thus, it provides much higher local stresses that can significantly reduce the cracking threshold; suitable for inducing well-developed redial cracks at low loads, particularly in brittle materials [42].The K c values are estimated by measuring the average lengths of radial cracks around sample indents induced with the diamond cube-corner tip over a load range from 10 to 50 mN.Five indents are made at each load.
Electrochemical behavior of the thin films is assessed with a Bio-logic VSP-300 potentiostat/galvanostat system using an Ag/AgCl electrode as reference electrode, a Pt mesh as counter electrode, and the films as working electrode (a three-electrode corrosion setup).All measurements are carried out in a stationary 1.0 M NaCl corrosive medium without applying external heating.The open circuit potential E ocp values are monitored versus the potential of the Ag/AgCl reference electrode.After 1-hour immersion, electrochemical impedance spectroscopy (EIS) measurements are performed in the frequency range of 100 kHz to 0.1 Hz with an amplitude of 10 mV at E ocp .The potentiodynamic polarizations are then performed from − 50 mV to 1200 mV with respect to E ocp using a sweeping rate of 1 mV/s.The corrosion potentials E corr and current densities i corr are calculated from Tafel extrapolations.

Composition and microstructure
Fig. 1 shows B/(Zr + Cr) and Cr/(Zr + Cr) ratios in Zr 1-x Cr x B y thin films.The B/(Zr + Cr) ratios, y, in the Zr 1-x Cr x B y alloy films deposited using hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS co-sputtering with the HiPIMS CrB 2 -target average power of 950 W (pulse frequency of 100 Hz) increases from 2.81 for P ZrB2 = 2000 W to 2.92 for P ZrB2 = 4000 W. For the alloys grown with the HiPIMS CrB 2 -target average power of 3400 W (pulse frequency of 300 Hz), the y value gradually increases from 1.81 for P ZrB2 = 400 W to 2.80 for P ZrB2 = 4000 W. Overall, the B contents in the films deposited by the hybrid technique varies in the range from 2.92 to 1.81.In contrast, the y-values, in the ZrB y and CrB y films grown using DCMS at 4000 W and a DC substrate potential of − 100 V, Fig. 1(a), span a narrower range of 2.18 and 1.8.
The Cr metal fraction, x, in the alloys deposited with the HiPIMS-CrB 2 -target average power of 950 W (pulse frequency of 100 Hz) decreases from 0.24 for P ZrB2 = 2000 W to 0.13 for P ZrB2 = 4000 W, Fig. 1  (b).For the Zr 1-x Cr x B y alloys grown with the HiPIMS-CrB 2 -target average power of 3400 W (pulse frequency of 300 Hz), the x value significantly decreases from 0.90 for P ZrB2 = 400 W to 0.32 for P ZrB2 = 4000 W.
The B/(Zr + Cr) ratio, y, is plotted versus the Cr metal fraction, x, in Fig. 1(c).For Zr 1-x Cr x B y alloy films grown in a hybrid scheme, y shows a gradual decrease from 2.90 to 1.81 as the fraction of DCMS-sputtered flux from the ZrB 2 target (arriving at the self-bias potential) in the total material flux to the substrate decreases.Thus, the hybrid growth scheme with metal-ion-synchronized substrate potential offers the ability to tune the film composition over very wide range, including the possibility for stoichiometric diboride around x = ~0.8.For comparison, we deposited reference DCMS ZrB y and CrB y thin films using a DC substrate potential of − 100 V, data points shown in Fig. 1(c), which indicate that the B/metal ratio can be varied in a much narrower range, from 2.18 to 1.8.
In order to assess the effect of substrate potential on the B/metal ratios, the ZrB y and CrB y thin films are also deposited using 4000-W DCMS with a floating substrate potential of − 10 V, which results in y = 3.2 and 1.7, respectively.Thus, we find a very strong effect of energetic bombardment on the B concentration in the case of the Zrcontaining film, with only minor effect in the case of the CrB 2 layer.This interesting result can be explained by the large mass-mismatch between the ions, predominantly Ar in this case (M Ar = 40 amu) and Zr (M Zr = 92.1 amu) that leads to the larger probability of backscattering of the ions (TRIM [43] simulations estimates the backscattering yield of ~ 6.8% with an energy up to ~ 15 eV).Energetic backscattered Ar may sputter B on the way back.For comparison, 100-eV Ar backscattering yield from Cr is only 0.07%, with a maximum energy that is below ~ 2 eV.This observation explains the large effect of DC substrate potential on the B concentration in the case of ZrB y films and the minor effect for CrB y layers.When we switch to predominantly Cr ion irradiation during the deposition of the alloy films in hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS configuration, the maximum backscattered energy of 100 eV Cr + from Zr drops to ~ 7.5 eV, which together with the fact that the energetic flux constitutes a small portion of the HiPIMS pulse period (1-3 %), leads to a limited decrease in the B concentration in the Zr-rich alloy films, ~ 0.1 ≤ x ≤ ~0.2.As the alloy film becomes more Cr-rich, they retain less B and eventually reach the CrB y value of y = 1.81.
Measuring the oxygen concentration after exposing the samples to air is an indirect approach that in addition to direct methods, like TEM, can be employed for evaluating the film densification.Upon exposure to ambient atmosphere, oxygen penetrates into the open column boundaries and contributes to a high oxygen concentration in porous films.If the films are dense, the oxygen concentrations measured ex-situ are indicative of the oxygen incorporation during film growth under high vacuum conditions.This approach has been applied for assessing the effect of metal ion irradiation on increasing the densification of TM nitride thin films obtained by hybrid HiPIMS/DCMS co-sputtering [44][45][46].Fig. 2 shows the ex-situ ToF-ERDA-measured oxygen concentration in Zr 1-x Cr x B y alloys grown by CrB 2 -HiPIMS/ZrB 2 -DCMS plotted as a function of Cr/(Zr + Cr) ratio, x.For x ≤ 0.24, the oxygen concentration is relatively high, from 1.5 at.% to 2.1 at.%, indicative of residual porosity (oxygen incorporation during film growth under present conditions accounts for < 0.2 at.% [4,5,18,47]).This occurs since these layers are grown at lower HiPIMS frequency of 100 Hz (necessary to reach low x), hence with a very low substrate potential duty cycle of only 1%.Consequently, majority of the film-forming species (the ZrB 2 flux) is deposited at the self-bias potential condition, which does not provide enough adatom mobility, necessary for densification.The oxygen content, however, decreases abruptly to ≤ 0.2 at.% for the films grown with 300 Hz, and we find a range of compositions, 0.32 < x < 0.9, where dense alloys are obtained, as evidenced by the low oxygen concentration.The reference ZrB 2.18 and CrB 1.81 films deposited with − 100 V DC substrate potentials show an increased level of oxygen; 1.5 and 0.8 at.%, respectively.The oxygen concentrations in the ZrB 3.2 and CrB 1.7 layers grown with self-bias potentials are even higher; 5.1 and 1.5 at.%, respectively.Thus, fully-dense layers are only obtained with the hybrid growth scenario, provided that the substrate potential duty cycle is sufficiently high.The concentration of Ar is ≤ 0.3 at.% in all coatings grown with − 100 V substrate potential.θ-2θ scans of Zr 1-x Cr x B y thin films are shown in Fig. 3.All Zr 1-x Cr x B y alloy films are crystalline as evidenced by strong XRD reflections from the crystalline hexagonal structure (solid solution).The positions of (000l) reflections move toward higher 2θ values that corresponds to a decrease in the out-of-plane (c) lattice parameter due mainly to the smaller size of Cr atoms incorporated in the diboride and the variation of residual stress levels.This lattice parameter gradually decreases from 3.535 Å for ZrB 2.18 to 3.045 Å for CrB 1.81 thin films as a function of the Cr concentration.The (1010) reflection disappears by alloying the ZrB 2.18 film with Cr.In addition, the (1011) peak intensity decreases by increasing Cr concentration up to x = 0.32, and this reflection disappears for the alloys with x > 0.32.
The increase of Cr metal fraction, x = Cr/(Zr + Cr), on the cation sublattice also significantly increases the full-width-at-half-maximum (FWHM) values of the XRD peaks that can be primarily attributed to a smaller size of coherently diffracting domains and an increase in microstrains.For example, the FWHM value of the (0001) reflection increases from 0.387 • for x = 0 to 0.852 • for x = 0.70, whereas it decreases to 0.444 • for x = 0.90 and 0.658 • for x = 1.0.While the (1011) reflections are dominant peaks in the XRD patterns of CrB 1.81 and ZrB 2.18 , the dominant peaks are (000l) for the Zr 1-x Cr x B y alloy films.[13] and CrB 2 [48], respectively.The peak at 32.8 • is a forbidden 002-substrate reflection [49].
where substrates face targets.The change observed in the columnar growth direction results from employing Cr + ion irradiation during the deposition, where the energetic ions steered by the electric potential near the substrate arrive at the surface along the substrate normal and enhance atomic displacement [50].In similarity to ZrB 2.18 , the columns of the CrB 1.81 film are also inclined (α = ~7• ) with respect to the substrate normal, but in opposite direction -toward the CrB 2 target.
The plan-view SEM images in Fig. 4 Thus, the Zr 1-x Cr x B y alloy films can be categorized in two groups: (i) the layers with x < 0.3 that have uninterrupted columnar growth, rough surfaces, and open column boundaries, which retain oxygen upon air exposure, and (ii) the layers with x > 0.3 that have dense microstructures, smooth surfaces, and low oxygen contents.Hence, we further study the nanostructure of representative sample from each group together with reference ZrB 2.18 and CrB 1.81 films using TEM.
Cross-sectional bright-field and dark-field TEM (BF-and DF-XTEM) images of ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , Zr 0.46 Cr 0.54 B 2.40 , Zr 0.10 Cr 0.90 B 1.81 , and CrB 1.81 thin films, with corresponding selected area electron diffraction (SAED) patterns, are compared in Fig. 5.The BF-and DF-XTEM images of ZrB 2.18 in Fig. 5(a) and 5(f) show that the DCMSgrown ZrB 2.18 thin film has a columnar nanostructure with porosities, which are marked with white arrows in Fig. 5(a).The ZrB 2.18 columns are continual from close to the substrate to the surface and inclined (α = ~7• ).This is attributed to the conservation of incident flux momentum component parallel to the film surface (the angle between the substrate and the ZrB 2 target is 21 • ) in combination with limited surface diffusion.The corresponding SAED pattern of ZrB 2.18 in Fig. 5(k) has diffraction rings with (0001), (1010), and (1011) components.The (0001) component has the weakest intensity, while (1011) is the strongest one, in agreement with the XRD θ-2θ pattern of ZrB 2.18 shown in Fig. 3(a).
The BF-and DF-XTEM images of the Zr 0.87 Cr 0.13 B 2.92 alloy film, Fig. 5(b) and 5(g), show that the concurrent metal ion irradiation achieved by the hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS co-sputtering in this compositional range, x < 0.3, is not sufficient to interrupt the continuous columnar growth.As a result, the films exhibit open porosities, marked with arrows, which develop due to the surface roughness evident in the SEM images in Fig. 4(b) and 4(h), and correspondently have high oxygen concentrations (~2.1 at.%).The SAED pattern of Zr 0.87 Cr 0.13 B 2.92 , Fig. 5 (l), consists of diffraction arcs with (0001), (1010), and (1011) components, in which the (1010) reflection has the weakest intensity.This SAED pattern also shows a change in the crystal orientation from predominant (1011) for ZrB 2.18 to (000l) for Zr 0.87 Cr 0.13 B 2.92 .
In contrast, the Zr 1-x Cr x B y alloy films with x > 0.3, represented by Zr 0.46 Cr 0.54 B 2.40 have a dense nanostructure with very fine columns.These columns do not extend throughout the whole layer (the continuous columnar growth is interrupted), Fig. 5(c) and 5(h).Thus, in the range of x values for which the films are dense, the ion fluxes relative to the ZrB x deposition rate are sufficiently high to cause recoil density that it eliminates porosity and results in re-nucleation in the column growth, which in turn maintains the film's surface smooth [51,52].Its SAED pattern is composed of (0001), (1010), and (1011) diffraction spots with a pronounced 000l fiber texture.The BF-and DF-XTEM images of Zr 0.10 Cr 0.90 B 1.81 in Fig. 5(d) and 5(i) confirm that this alloy has a dense crystalline columnar nanostructure with columns that appear wider than those of Zr 0.46 Cr 0.54 B 2.40 .The reflections in the Zr 0.10 Cr 0.90 B 1.81 SAED pattern arise from the (0001), (1010), and (1011) planes with pronounced 000l orientation, Fig. 5(n), in agreement with its XRD θ-2θ signature in Fig. 3(i).
Fig. 6 shows the plan-view BF-and DF-TEM images of ZrB  6(c) and 6(d), do not show bright column boundaries and instead, exhibit dense nanostructures with point-defect clusters induced due to ion irradiation during film growth.The CrB 1.81 thin film has a porous nanostructure, indicated with arrows in Fig. 6(e).In agreement with the XTEM micrographs in Fig. 4, the   The H(x) plot shown in Fig. 7 is the effect of two factors.First, in the range x < 0.32, the films are underdense, as evidenced by higher oxygen content (Fig. 2) and their cross-sectional and plan-view TEM images (Figs. 5 and 6), which results in lower H values.The layers with 0.32 ≤ x ≤ 0.9 are dense (deposited at higher substrate potential duty cycle of 3%), and the decreasing trend in H is due to that the hardness of CrB 1.81 (25.5 ± 2.0 GPa) is lower than that of ZrB 2.18 (30.0 ± 1.0 GPa).Hence, the maximum hardness obtained for Zr 0.68 Cr 0.32 B 2.80 can be attributed to its dense nanostructure, highest Zr content of all dense alloys, solid solution hardening [53], and narrow column widths (based on Hall-Petch effect [54,55]).

Mechanical properties
The elastic moduli of the ZrB 2.18 and CrB 1.81 thin films grown using − 100 V substrate potential are 485 ± 20 GPa and 345 ± 19 GPa, respectively, while the ZrB 3.2 and CrB 1.7 films deposited at floating substrate potential are 345 ± 18 GPa and 300 ± 12 GPa, respectively.The alloy films show lower E values than the reference layers fluctuating between 320 GPa and 345 GPa.Analogous to their H values, the Zr 1- x Cr x B y alloy films exhibit close elastic moduli.
The in-plane residual stress σ f of the films decreases from + 1.68 ± 0.07 GPa for ZrB 2.18 , to + 1.65 ± 0.08 GPa for Zr 0.87 Cr 0.13 B 2.92 , to + 0.11 ± 0.06 GPa for Zr 0.68 Cr 0.32 B 2.80 , to + 0.05 ± 0.07 GPa for Zr 0.46 Cr 0.54 B 2.40 , to − 0.26 ± 0.15 GPa for Zr 0.10 Cr 0.90 B 1.81 (see Fig. 8).Then, it slightly increases to + 0.16 ± 0.07 GPa for the CrB 1.81 layer.The σ f values are corrected for thermal stresses σ th due to cooling the films  from deposition temperature to room temperature, ΔT = 450 K, using the equation [56]; in which α f and α s are thermal expansion coefficients of film and substrate, respectively.The thermal expansion coefficient α s of Al 2 O 3 is 8.1 × 10 -6 K − 1 [57].For the Zr 1-x Cr x B y alloy films, α f is estimated from a linear interpolation between the thermal expansion coefficients of ZrB 2 (5.9 × 10 -6 K − 1 ) and CrB 2 (10.5 × 10 -6 K − 1 ) [58], and using elastic moduli E determined by nanoindentation.σ th increases from − 0.562 GPa for ZrB  [59,60] and high surface roughness [61].However, for the alloys with x ≥ 0.32 several effects simultaneously result in their low residual stresses (almost stress free layers) such as the formation of dense nanostructure [62], ion-induced lattice damages [63,64], the crystal orientation change [65] as shown in the XRD data in Fig. 3, the incorporation of Cr atoms, which have smaller covalent radius than Zr atoms [66], in the diboride structure, and the decrease in the column widths [67].
The relative cube-corner nanoindentation toughness K c of ZrB in which α is the indenter geometry coefficient, 0.0319 [69]; C m is the average length of radial cracks around a cube-corner indent; and P is the applied load.This equation was initially derived for bulk materials with infinite thicknesses [68,70].Thus, for ceramic coatings, the coating thickness and substrate effects must be accounted for.One approach is to measure K c at different loads, plot the K c values as a function of maximum indentation penetration depth h max , and then extrapolate the results to h max = 0 [71].Fig. 8 shows the K c values of the Zr 1-x Cr x B y thin films, determined using this approach, as a function of x.A range of loads from 10 to 50 mN is used for determining different K c .The K c value decreases from 1.7 ± 0. alloy films, which proves their higher ductility.11(a), which is composed of a constant phase element (CPE), polarization resistance (R p ), and solution resistance (R s ).CPE indicates a double layer capacitance in parallel with R p .In general, CPE is a non-ideal capacitance that is extensively employed to mimic the defective dielectric behavior.The CPE values are calculated by the equation [72];

Corrosion properties
where P is the magnitude of CPE and n is the deviation parameter.The equivalent circuit reveals that all layers follow a charge-transfercontrolled process in 1.0 M NaCl, while there are no other elements associated with a diffusion-controlled process.Both reference DCMS ZrB 2.18 and CrB 1.81 thin films exhibit comparably high corrosion resistance.In addition, alloying ZrB y layers with CrB y using the hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS co-sputtering does not have a considerable influence on the corrosion resistance of the Zr 1- x Cr x B y thin films with x ≤ 0.13.Plan-view and cross-sectional SEM and TEM images in Figs. 4, 5, and 6 show that the reference layers and the alloy films with x ≤ 0.13 have rough surfaces and porous columnar structures in which the columns are extended along the films.Such structural properties may make these thin films susceptible to corrosion attacks.The potentiodynamic polarization and EIS results summarized in Tables 1 and 2 prove that ZrB 2.18 , CrB 1.81 , and Zr 0.87 Cr 0.13 B 2.92 have higher corrosion rates than the other alloys.
However, the Zr  1 and 2. These impressive corrosion properties can be mainly attributed to their non-columnar growth, smooth surfaces, and fully-dense fine-grain microstructures.The other factor that also effectively improves the corrosion resistance of the alloys with ≥ 0.32 is their chemistry.These layers contain more Cr, but less B. As Zr and Cr are more noble elements with lower dissolving rates; thus, it can enhance the corrosion properties of these alloys.Zr 0.46 Cr 0.54 B 2.40 has the lowest corrosion rate (i corr = ~1.0× 10 -6 mA/cm 2 ) over the entire compositional range studied here; however, this rate is almost twice as high as the corrosion rate of Zr 0.56 Cr 0.44 B 1.11 (i corr = ~0.54× 10 -6 mA/cm 2 ) that has an amorphous nanostructure and very-low B concentration [18].It is a good comparison that reveals the influence of   both chemistry and structure on the corrosion properties of diboride alloy thin films.The slight decrease in the corrosion resistance together with the absence of the formation of the passive region in the anodic polarization curve of the Zr 0.10 Cr 0.90 B 1.81 alloy film, Fig. 10, can be primarily due to the intensive ion bombardment during the film growth that induces a large local distortions and residual strains in its crystal lattice, see Fig. 8.These regions with locally ion-induced strains can be highly prone to corrosion attacks and does not allow the passive-layer formation.On the other hand, the formation of the passive regions in the layers with x ≤ 0.54 can be mostly attributed to the presence of diffusion-controlled zones, proved by the EIS measurements.

Conclusions
The use of the hybrid CrB 2 -HiPIMS/ZrB 2 -DCMS co-sputtering with Cr + ion bombardment changes the predominant crystallographic orientation from (1011) for the DC-sputter-deposited reference ZrB 2.18 and CrB 1.81 thin films to (000l) for the Zr 1-x Cr x B y alloys.The columns of DCMS-grown ZrB 2.18 and CrB 1.81 are continual from the substrate to the surface with porous structure.We find that to eliminate pores and achieve dense alloy films, a critical intensity of the Cr + -dominated ion flux, relative to the deposition rate from the ZrB 2 target, is required.Correspondingly, the films can be categorized in two group: (i) alloys with x < 0.32 that have continuous columnar growth, rough surfaces, and open column boundaries, and (ii) alloys with x ≥ 0.32, where the intensity of Cr + -dominated ion fluxes are sufficient to interrupt continuous column growth, which exhibits a smooth surface and dense finegrain microstructure.The oxygen content in the underdense alloy films with x < 0.32 is in the range of 1.5-2 at.%, while it drops below 0.2 at.% for the alloy films with x ≥ 0.32.The later number is significantly lower than the value of 1.5 at.% for the reference films deposited under continuous 100 eV Ar + bombardment in film densification.Compared to the porous films, the dense alloy films show higher hardness, in the range of 29.5-32.5 GPa.All porous films show tensile stress, while the dense alloy films are almost stress-free.In addition, they exhibit good average relative nanoindentation toughness of 1.3 MPa√m to 2.3 MPa√m, and remarkedly low corrosion rates (i corr = ~1.0× 10 -6 mA/cm 2 for Zr 0.46 Cr 0.54 B 2.40 and i corr = ~2.1 × 10 -6 mA/cm 2 for Zr 0.10 Cr 0.90 B 1.81 ).
These results reveal that the hybrid HiPIMS/DCMS co-sputtering technique provides energetic fluxes of metal ions resulting in dense TM diboride layers with smooth surfaces and a combination of favorable properties, similar to results previously reported for ternary TM nitrides.The hybrid technique is superior to DCMS.We find that a threshold ion flux density is required to achieve dense Zr 1-x Cr x B y diboride thin films.

Declaration of Competing Interest
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

B
.Bakhit et al.

Fig. 4 Fig. 2 .
Fig. 2. Oxygen concentration in Zr 1-x Cr x B y thin films as a function of Cr/(Zr + Cr) ratio, x.
(g)-4(l) show that Zr 0.68 Cr 0.32 B 2.80 , Zr 0.46 Cr 0.54 B 2.40 , and Zr 0.10 Cr 0.90 B 1.81 have significantly smoother surfaces than the ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , and CrB 1.81 films.Both ZrB 2.18 and Zr 0.87 Cr 0.13 B 2.92 layers have feather-like surface morphologies, where the features are inclined.There is also an increase in the feature sizes from ZrB 2.18 to the Zr 0.87 Cr 0.13 B 2.92 alloy.Opposite to these layers, the surface morphology of the CrB 1.81 film appears rough, consisting of granular features with different sizes and shapes.Fig. 4(i), 4(j), and 4(k) exhibit that the Zr 0.68 Cr 0.32 B 2.80 , Zr 0.46 Cr 0.54 B 2.40 , and Zr 0.10 Cr 0.90 B 1.81 alloy films have featureless, smooth surface morphologies.

Fig. 7 Fig. 5 .
Fig.7shows the H and E values of the Zr 1-x Cr x B y thin films as a function of x.The H value of ZrB 2.18 grown by DCMS using − 100 V B 2.90 , to 29.0 ± 2.0 GPa for Zr 0.76 Cr 0.24 B 2.81 , to 32.5 ± 1.5 GPa for Zr 0.68 Cr 0.32 B 2.80 .However, adding more CrB y results in a slight decrease in H to 31.5 ± 1.9 GPa for Zr 0.57 Cr 0.43 B 2.50 , 30.5 ± 2.0 GPa for Zr 0.46 Cr 0.54 B 2.40 , 30.0 ± 1.4 GPa for Zr 0.30 Cr 0.70 B 2.14 , and 29.0 ± 1.5 GPa for Zr 0.10 Cr 0.90 B 1.81 .The CrB 1.81 thin film deposited by DCMS using − 100 V substrate potential has an H value of 25.5 ± 2.0 GPa, whereas the hardness of CrB 1.7 grown by DCMS at floating substrate potential is 20.5 ± 1.2 GPa.

Fig. 7 .
Fig. 7. (a) Nanoindentation H and (b) E values of Zr 1-x Cr x B y thin films as a function of x ranging from 0 to 1.0.

Fig. 10
Fig. 10 compares the potentiodynamic polarization curves of ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , Zr 0.46 Cr 0.54 B 2.40 , Zr 0.10 Cr 0.90 B 1.81 , and CrB 1.81 thin films in a 1.0 M NaCl corrosive aqueous medium.The corrosion potentials (E corr ), corrosion current densities (i corr ), and passive current densities (i pass ) determined from the polarization curves are summarized in Table1.All layers exhibit high corrosion resistance with a very low dissolving rate across their anodic polarization curves.The i corr value, which is considered as an indicator for corrosion rate, significantly decreases from (3.0 ± 0.4) × 10 -6 mA/cm 2 for ZrB 2.18 and (8.2 ± 0.7) × 10 -6 mA/cm 2 for Zr 0.87 Cr 0.13 B 2.92 to (1.0 ± 0.3) × 10 -6 mA/cm 2 for Zr 0.46 Cr 0.54 B 2.40 .However, it increases to (2.1 ± 0.3) × 10 -6 mA/cm 2 for Zr 0.10 Cr 0.90 B 1.81 and (4.3 ± 0.5) × 10 -6 mA/cm 2 for CrB 1.81 .The anodic polarization curves of ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , and Zr 0.46 Cr 0.54 B 2.40 consist of passive regions formed at ~ 210 mV for ZrB 2.18 and Zr 0.87 Cr 0.13 B 2.92 and at ~ 450 mV for Zr 0.46 Cr 0.54 B 2.40 .The passive current density i pass is (2.3 ± 0.2) × 10 -3 mA/cm 2 for ZrB 2.18 and Zr 0.87 Cr 0.13 B 2.92 , while i pass is (1.3 ± 0.3) × 10 -3 mA/cm 2 for Zr 0.46 Cr 0.54 B 2.40 .In contrast, both Zr 0.10 Cr 0.90 B 1.81 and CrB 1.81 do not show any passivity up to 1200 mV and follow dominant Tafel behavior.The electrochemical behavior of the coatings is further studied using EIS measurements in the 1.0 M NaCl corrosive medium.The Nyquist and Bode curves of ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , Zr 0.46 Cr 0.54 B 2.40 , Zr 0.10 Cr 0.90 B 1.81 , and CrB 1.81 thin films are plotted in Fig. 11.All Nyquist plots in Fig. 11(a) show single imperfect semicircles with the smallest radius for Zr 0.87 Cr 0.13 B 2.92 , and the largest radii for Zr 0.46 Cr 0.54 B 2.40 and Zr 0.10 Cr 0.90 B 1.81 .These plots indicate that there is no other capacitive loops or linear tails in low frequency range.The impedance spectra are fitted to a simple equivalent circuit, inset in Fig. 11(a), which is composed of a constant phase element (CPE), polarization resistance (R p ), and solution resistance (R s ).CPE indicates a double layer capacitance in parallel with R p .In general, CPE is a non-ideal capacitance that is extensively employed to mimic the defective dielectric behavior.The CPE values are calculated by the equation [72]; Fig. 11(b)  shows that the phase angles of all layers are close to 90 • due to their high polarization resistance.All films exhibit

Fig. 11 .
Fig. 11.(a) Nyquist and Bode (b) phase change and (c) impedance plots of ZrB 2.18 , Zr 0.87 Cr 0.13 B 2.92 , Zr 0.46 Cr 0.54 B 2.40 , Zr 0.10 Cr 0.90 B 1.81 , and CrB 1.81 thin films.The inset in (a) is the equivalent circuit model used to fit the EIS spectra.
1-x Cr x B y alloys with ≥ 0.32 exhibit considerably better corrosion resistance.The corrosion resistance of Zr 0.46 Cr 0.54 B 2.40 and Zr 0.10 Cr 0.90 B 1.81 thin films is above seven times higher than ZrB 2.18 , CrB 1.81 , and Zr 0.87 Cr 0.13 B 2.92 , see Figs. 10 and 11 together with their polarization and EIS data in Tables

Table 2
Impedance parameters extracted from fitting the EIS data using the equivalent circuit model.Uppsala University), and 2021-00357, Swedish Energy Agency (Grant No. 51201-1), Swedish for Strategic Research (SSF), Swedish National Graduate School in Neutron Scattering (SwedNess), and Swedish Government Strategic Research Area in Materials Science on Advanced Functional Materials (AFM) at Linköping University (Faculty Grant SFO Mat LiU No.).