Multifunctional ZrB2-rich Zr1-xCrxBy thin films with enhanced mechanical, oxidation, and corrosion properties

Refractory transition-metal (TM) diborides have high melting points, excellent hardness, and good chemical stability. However, these properties are not sufficient for applications involving extreme environments that require high mechanical strength as well as oxidation and corrosion resistance. Here, we study the effect of Cr addition on the properties of ZrB 2 -rich Zr 1-x Cr x B y thin films grown by hybrid high-power impulse and dc magnetron co-sputtering (Cr-HiPIMS/ZrB 2 -DCMS) with a 100-V Cr-metal-ion synchronized bias. Cr metal frac- tion, x = Cr/(Zr + Cr), is increased from 0.23 to 0.44 by decreasing the power P ZrB 2 applied to the DCMS ZrB 2 target from 4000 to 2000 W, while the average power, pulse width, and frequency applied to the HiPIMS Cr target are maintained constant. In addition, y decreases from 2.18 to 1.11 as a function of P ZrB 2 , as a result of supplying Cr to the growing film and preferential B resputtering caused by the pulsed Cr-ion flux. ZrB 2.18 , Zr 0 ⋅ 77 Cr 0 ⋅ 23 B 1.52 , Zr 0 ⋅ 71 Cr 0 ⋅ 29 B 1.42 , and Zr 0 ⋅ 68 Cr 0 ⋅ 32 B 1.38 films have hexagonal AlB 2 crystal structure with a columnar nanostructure, while Zr 0 ⋅ 64 Cr 0 ⋅ 36 B 1.30 and Zr 0 ⋅ 56 Cr 0 ⋅ 44 B 1.11 are amorphous. All films show hardness above 30 GPa. Zr 0.56 Cr 0.44 B 1.11 alloys exhibit much better toughness, wear, oxidation, and corrosion resistance than ZrB 2.18 . This combination of properties makes Zr 0 ⋅ 56 Cr 0 ⋅ 44 B 1.11 ideal candidates for numerous strategic applications.


Introduction
Transition-metal (TM) nitride coatings have many industrial applications from mechanical components in aerospace industry to cutting tools [1][2][3][4]. Metastable NaCl-structure Ti 1-x Al x N layers grown by magnetron sputtering are the most attractive group of TiN-based thin films, suitable as protective coatings for cutting tools, which show good hardness (typically ~30 GPa), high wear and oxidation resistance (depending on Al concentration), and self-hardening effects at elevated temperatures up to ~900 • C (resulting from spinodal decomposition) [3][4][5]. However, the ever-increasing demand from industry for enhanced coating properties motivates the search for alternatives.
One particularly promising group of materials are TM diborides, extensively studied in the recent years. TM diborides, which typically crystallize in a hexagonal AlB 2 structure (P6/mmm, SG-191) -where B atoms form graphite-like honeycomb sheets between hexagonal-closepacked TM layers [6,7], exhibit high melting points, excellent hardness, high thermal and chemical stability, and good conductivity [8]. This unique combination of properties originates from their dual ceramic/metallic nature where strong combined covalent/ionic bonding between TM and B atoms together with the covalent bonding within the honeycomb B sheets provide high melting point, hardness, and stiffness [9,10], while metallic bonding between TM atoms results in good thermal and electrical conductivities [6]. Hence, TM diborides are good candidates for a broad range of applications, particularly in extreme environments, such as hypersonic aerospace vehicles [11,12], rockets [12], nuclear reactors [8], optoelectronic and microelectronic components [13,14], solar power [15], and cutting tools [16][17][18][19][20].
However, compared to TiAlN, the industrial applications of sputterdeposited TM diboride thin films are very limited, primarily due to poor oxidation resistance [16] and high brittleness [17]. Bulk diborides, which are mostly synthesized by powder metallurgy processes [11,21], start to oxidize at temperatures below ~450 • C with oxidation products that are typically TMO 2 and glassy B 2 O 3 phases [22]. The B 2 O 3 phase tends to rapidly evaporate at temperatures above ~1000 • C which results in the formation of a porous oxide scale that does not passivate the surface against oxidation [22]. This issue is even worse for the TM diboride thin films that are overstoichiometric (B/TM ratios > 2), in which oxide scales formed at temperatures above ~400 • C are highly B deficient with no oxidation protection; causing a poor oxidation resistance [16]. Similar to TiN-based thin films, alloying TM diborides with Al enhances their oxidation properties [16].
Moreover, the applications of sputter-deposited TM diborides are restricted due to their inherent brittleness [17]. Although these films have high hardness ranging from 30 to 50 GPa [20,23,24], this alone is not sufficient for preventing failure in applications which involve high mechanical stresses. Hardness is usually accompanied by brittleness that causes crack formation and propagation at the presence of high stresses [25]. Hence, TM diboride films require to have a combination of high hardness and ductility (referred to as toughness [26]) in order to avoid brittle cracking. To accommodate this requirement, we recently showed that alloying ZrB 2 thin films with Ta can result in a simultaneous increase in both nanoindentation hardness and toughness [17]. Zr 1-x Ta x B y alloys with x ≥ 0.2 exhibit a self-organized columnar core/shell nanostructure in which crystalline Zr-rich stoichiometric Zr 1-x Ta x B 2 cores are surrounded by narrow dense, disordered Ta-rich (B-deficient) shells that have the structural characteristics of metallic-glass thin films; both high strength and toughness [18]. These layers also show a high thermal stability in which their hardness increases as a function of annealing temperature up to 800 • C. The age hardening observed in the Zr 1-x Ta x B y films with 0 ≤ x ≤ 0.3, which occurs without any phase separation or decomposition, can be explained by point-defect recovery that enhances the chemical bond density [20]. For temperatures above 800 • C, hardness decreases due to recrystallization, column coarsening, and stacking fault annihilation. All Zr 1-x Ta x B y films generally have hardness values H > 34 GPa up to 1200 • C [20].

Experimental
Zr 1-x Cr x B y thin films are grown in a CC800/9 CemeCon AG sputtering system equipped with rectangular 8.8 × 50 cm 2 stoichiometric ZrB 2 and elemental Cr targets. Al 2 O 3 (0001), Si(001), and WC-Co substrates, 1.5 × 1.5 cm 2 , are cleaned sequentially in acetone and isopropyl alcohol, and then mounted symmetrically with respect to the targets, which are tilted toward the substrates resulting in a 21 • angle between the substrate normal and the normal to each target. The target-to-substrate distance is 20 cm. The chamber is degassed before deposition by applying 8.8 kW to each of two resistive heaters for 2 h, which results in a temperature of ~475 • C at the substrate position. The system base pressure is 3.8 × 10 − 6 Torr (0.5 mPa). The film growth is carried out at ~475 • C and a total Ar pressure of 3 mTorr (0.4 Pa). Prior to deposition, the targets are sequentially DCMS sputter-cleaned in Ar at 2 kW for 60 s with shutters protecting the substrate table and the opposite target. A thin continuous Cr buffer layer, with a thickness of 4 ± 1 nm, is initially deposited on all substrates to improve adhesion and minimize their influence on the film morphological evolution.
ZrB y films are grown by DCMS with a target power of 4 kW and a negative dc substrate bias of 100 V. For growing the Zr 1-x Cr x B y films, a hybrid target-power scheme [32] (Cr-HiPIMS/ZrB 2 -DCMS) is employed in which the ZrB 2 target is continuously sputtered by DCMS, while the Cr magnetron is operated in HiPIMS mode to provide pulsed Cr ion fluxes. The Cr metal fraction, Cr/(Zr + Cr), is increased from 0.23 to 0.44 by decreasing the power P ZrB2 applied to the DCMS ZrB 2 target from 4000 to 2000 W in 500-W increments, while the average power, pulse width, and frequency applied to the HiPIMS Cr target are maintained constant at 700 W, 50 μs, and 100 Hz, respectively. This results in a constant peak Cr-target current density of ~0.73 A/cm 2 . A negative substrate bias of 100 V is applied in synchronous with the 100-μs metal-ion-rich portion of each HiPIMS pulse, starting 30 μs after the cathode HiPIMS pulse. The substrates are at a negative floating potential of 10 V at all other times. The film deposition rate is ~0.85 nm/s for ZrB 2.4 , while it increases from ~0.45 nm/s for P ZrB2 = 2000 W to ~0.78 nm/s for P ZrB2 = 4000 W for the Zr 1-x Cr x B y films.
Cross-sectional scanning electron microscopy (XSEM) analyses are conducted in a Zeiss LEO 1550 electron microscope to obtain the thicknesses and cross-sectional morphologies of the films. θ-2θ X-ray diffraction (XRD) scans are carried out using a Philips X'Pert X-ray diffractometer with a Cu K α source (λ = 0.15406 nm) to determine crystal structure and orientations of the layers. Film compositions are obtained from time-of-flight elastic recoil detection analyses (ToF-ERDA) 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 • . Chemical bonding in the films is evaluated by X-ray photoelectron spectroscopy (XPS) using a Kratos Axis Ultra DLD instrument employing monochromatic Al K α radiation (hν = 1486.6 eV). All surfaces are sputter-etched for 120 s with a 4-keV Ar + ion beam incident at 70 • with respect to the sample normal. Then, the Ar + ion energy is reduced to 0.5 keV for 600 s to minimize surface damage. The analyzed area, which is located in the center of a 3 × 3 mm 2 ion-etched region, is 0.3 × 0.7 mm 2 . The core level spectra are referenced to the Fermi edge cut-off to avoid problems caused by the referencing method based on the C 1s peak from adventitious carbon [33].
Cross-sectional transmission electron microscopy (XTEM) analyses are carried out in a monochromated and double-corrected FEI Titan 3 60-300 electron microscope operated at 300 kV. Images are acquired using bright-field (BF) and dark-field (DF) TEM imaging modes. TEM specimens are prepared by mechanical polishing, followed by Ar + ion milling at 5 keV, with a 3 • incidence angle, on both sides of each sample during rotation, in a Gatan precision ion miller. The specimens are finally sputter-cleaned using an ion energy of 0.5 keV without changing the angle of incident Ar + ions.
The in-plane residual stresses of Zr 1-x Cr x B y thin films are obtained using the modified Stoney equation by determining the substrate wafer curvature from XRD rocking-curve measurements. More details are provided in reference 17. The nanoindentation analyses of the layers are performed in an Ultra-Micro Indentation System with a sharp Berkovich diamond tip calibrated using a fused-silica standard. For hardness H and elastic modulus E measurements, the layers are indented using a fixed load of 12 mN, while indention depths are maintained below 10% of the film thickness. Reported values are the average of 35 indentations. The results are analyzed using the Oliver and Pharr method [34]. The films are also indented by a diamond cube-corner tip with a load of 200 mN to measure the average lengths of induced radial cracks. The average crack length, which is an indication of nanoindentation toughness [35], is obtained from four cube-corner indents for each film.
To evaluate the films' adhesion and toughness, the Revescratch tests are performed using an Anton-Paar-TriTec UNHT 3 system equipped with a Rockwell-C diamond indenter with a tip radius of 100 μm. A progressive loading regime is used in which the load is linearly increased from 1 to 80 N with a rate of 158 N/min. The scratch length is 3 mm with a speed of 6 mm/min. The same equipment with a ball-on-disc tribometer with a 3-mm-diameter GCR15 steel ball is also used to investigate the friction and wear properties of the layers at room temperature. A load of 2 N with 0.1 m/s sliding speed (2000 laps) is applied during the wear tests. The wear track profiles are measured by a confocal laserscanning microscope. The wear rates are obtained using the following equation [36]: where V is the volume loss by wear (mm 3 ), F is the applied load (N), and s is the sliding distance (m). ZrB 2.18 and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films are annealed at 700 • C in air for times t a ranging from 1.0 to 5.0 h using a high-temperature furnace from MTI Corporation (GSL-1100 × -S). The heating rate is constant at 10 • C/min, and the specimens are cooled down to room temperature, while the furnace is turned off.
Open circuit potential (E ocp ) with superimposed linear polarization resistance (R p ) followed by potentiodynamic polarization measurements are employed to study the corrosion resistance of ZrB 2.18 and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films. All measurements are carried out in an aqueous 0.1 M NaCl corrosive medium, at room temperature and without agitation, using a Bio-logic VSP-300 potentiostat/galvanostat system. A standard three-electrode system is used with a silver/silver chloride electrode (Ag/AgCl) as the reference electrode, a platinum mesh as the counter electrode, and the films as the working electrode. The E ocp values of the films immersed in the corrosive medium are monitored for 25 h and reported versus the Ag/AgCl reference electrode potential, unless mentioned differently. The R p measurement is performed after 0.25, 0.5, 1.0, 2.0, 4.0, 8.0, 16.0, and 24.0 h of immersion by a sweeping potential of ±10 mV versus E ocp with a scanning rate of 0.167 mV/s. The R p values are obtained from the inverse of the slopes of current-potential plots at the corrosion potential (E corr ). Immediately afterward, the potentiodynamic polarization is performed with a sweeping rate of 0.167 mV/s from − 160 to +1260 mV with respect to E ocp . The corrosion potentials (E corr ) and current densities (i corr ) are calculated according to the Tafel extrapolation [37,38]. Table 1 gives the elemental compositions of as-deposited Zr 1-x Cr x B y thin films obtained from ToF-ERDA measurements. The as-deposited ZrB y films grown using DCMS at P ZrB2 = 4000 W are overstoichiometric with the B/(Zr + Cr) ratio y of 2.18. The Cr/(Zr + Cr) ratio, x, in the alloys deposited by hybrid Cr-HiPIMS/ZrB 2 -DCMS cosputtering increases from 0.23 for P ZrB2 = 4000 W, to 0.29 for P ZrB2 = 3500 W, 0.32 for P ZrB2 = 3000 W, 0.36 for P ZrB2 = 2500 W, and 0.44 for P ZrB2 = 2000 W, while the B/(Zr + Cr) ratio, y, gradually decreases from 1.52 to 1.42, 1.38, 1.30, and 1.11 with decreasing P ZrB2 . The total concentration of carbon, nitrogen, and oxygen is ≤ 1.6 at. %, and the Ar concentration is ≤ 0.5 at. % in all films. Alloying ZrB y with Cr using a flux of energetic Cr ions bombarding the growing film not only adds Cr atoms, but it also affects the B content via preferential resputtering.

Elemental compositions and microstructure
XRD θ-2θ scans of as-deposited Zr 1-x Cr x B y thin films grown on Si (001) substrates are shown in Fig. 1. Vertical solid and dashed lines correspond to reference powder-diffraction peak positions for ZrB 2 [39] and CrB 2 [40], respectively. All reflections in the XRD patterns of ZrB 2.18 , Zr 0⋅77 Cr 0⋅23 B 1.52 , Zr 0⋅71 Cr 0⋅29 B 1.42 , and Zr 0⋅68 Cr 0⋅32 B 1.38 films originate from the crystalline hexagonal AlB 2 -type structure (solid-solution), while the patterns of Zr 0⋅64 Cr 0⋅36 B 1. 30 and Zr 0⋅56 Cr 0⋅44 B 1.11 films show very low intensity (notice the logarithmic scale), broad 0001 and 1010 X-ray reflections; indicating that they are X-ray amorphous. The (1010) reflection disappears for Zr 0⋅71 Cr 0⋅29 B 1.42 and Zr 0⋅68 Cr 0⋅32 B 1.38 alloys. The formation of X-ray amorphous Zr 0⋅64 Cr 0⋅36 B 1.30 and Zr 0⋅56 Cr 0⋅44 B 1.11 films can be attributed to the collapse of the hexagonal AlB 2 -structure that results from the lack of B between the hexagonal-close-packed Zr 1-x Cr x layers (y ≤ 1.30) as well as the difference between the crystal structures of Zr and Cr (Zr has a hexagonal close-packed structure, while Cr has a body-centered-cubic structure [41]).
While the position of (1011) reflections does not change with increasing the Cr concentration, the positions of (0001) and (0002) reflections shift toward higher 2θ values, corresponding to a decrease in the out-of-plane c lattice parameter from 0.352 nm for ZrB 2.18 to 0.349 nm for Zr 0⋅68 Cr 0⋅32 B 1.38 . This is mainly due to the smaller covalent radius of Cr atoms incorporated in the diboride structure, the corresponding lower B concentrations, and a change in the film's residual stress level. The incorporation of Cr atoms also results in a significant increase in the Table 1 Concentrations of primary elements, B, Cr, and Zr in as-deposited Zr 1-x Cr x B y thin films grown on Si(001) substrates, obtained from ToF-ERDA, as a function of ZrB 2 target power P ZrB2 . The total concentration of contaminants (not included in the  (001) substrate, which appears due to multiple scattering events [42].
full-width at half-maximum values of the XRD reflections; e.g. from 0.18 • for ZrB 2.18 to 0.4 • for Zr 0⋅68 Cr 0⋅32 B 1.38 for the (0001) reflection. B 1s, Zr 3d, and Cr 2p XPS core-level spectra acquired from the asdeposited Zr 1-x Cr x B y thin films grown on Si(001) substrates are plotted in Fig. 2. The B 1s and Zr 3d spectra shown in Fig. 2(a) are normalized to the intensity of the Zr 3d 5/2 peaks. The Zr 3d 3/2 and 3d 5/2 peaks appear at 181.3 and 178.9 eV, respectively, with no detectable change in their positions or shapes as a function of Cr concentration. The position of the B 1s peaks does not change noticeably for ZrB 2.18 , Zr 0⋅77 Cr 0⋅23 B 1.52 , Zr 0⋅71 Cr 0⋅29 B 1.42 , and Zr 0⋅68 Cr 0⋅32 B 1.38 ; the peaks appear at ~188.0 eV. This indicates that the incorporation of Cr, which has a slightly higher electronegativity than Zr (1.66 for Cr and 1.33 for Zr, based on the Pauling scale [43]), does not change the effective valence-charge density residing on the B atoms. However, there is a slight shift in the position of the B 1s peaks toward lower binding energies for the Zr 0⋅64 Cr 0⋅36 B 1. 30 and Zr 0⋅56 Cr 0⋅44 B 1.11 alloys (~187.7 eV). In addition, the width of B 1s peaks for these alloys is larger than the other films. These slight changes correlate to the apparent loss of crystalline structure (see Fig. 1) and, hence, cannot be directly related to the change in the bonding configuration, as likely the layer electrical properties are modified, which may have a direct effect on the screening ability [44]. Fig. 2(b) shows that increasing Cr concentration does not have an obvious effect on the positions and shapes of the Cr 2p peaks (the Cr 2p spectra in Fig. 2 18 consists of discernable porosities, marked by black arrows in the micrograph. The ZrB 2.18 columns with a width of 10.1 ± 2 nm near the film's surface are continual from close to the substrate toward the surface. The columns are inclined at an angle of 7 • with respect to the substrate normal, due to the 21 • angle between the substrate and the ZrB 2 target. The corresponding SAED pattern, the inset in Fig. 3(e), is composed of diffraction arcs with (0001), (1010), and (1011) components in which the (0001) signal in the growth direction is the weakest one, in agreement with the XRD result in Fig. 1(a).
The BF-and DF-XTEM images of Zr 0⋅77 Cr 0⋅23 B 1.52 and Zr 0⋅68 Cr 0⋅32 B 1.38 , Figs. 3(f), (j), 3(g), and 3(k), show that alloying with Cr interrupts the continuous columnar growth and produces dense nanostructure. The column length of Zr 1-x Cr x B y alloys decreases as a function of Cr concentration up to x = 0.32. Moreover, adding Cr leads to a decrease in the column width; the nanostructure of Zr 0⋅68 Cr 0⋅32 B 1. 38 consists of very fine columns that do not extend throughout the whole film, see Figs. 3(g) and 3(k). The corresponding SAED patterns of Zr 0⋅77 Cr 0⋅23 B 1.52 and Zr 0⋅68 Cr 0⋅32 B 1.38 alloys, the insets in Fig. 3(f) and 3 (g), indicate the presence of (0001), (1010), and (1011) diffraction arcs with a decrease in the crystallinity by increasing the Cr concentration. The BF-and DF-XTEM micrographs of Zr 0⋅56 Cr 0⋅44 B 1.11 Fig. 3(h), confirm that this alloy has an amorphous nanostructure, which is consistent with its XRD θ-2θ result in Fig. 1(f).

Mechanical properties
The residual stress of as-deposited Zr 1-x Cr x B y thin films grown on Al 2 O 3 (0001) substrates changes from +0.91 ± 0.04 GPa for ZrB 2.18 , to − 0.83 ± 0.23 GPa for Zr 0⋅77 Cr 0⋅23 B 1.52 , − 1.27 ± 0.15 GPa for Zr 0⋅71 Cr 0⋅29 B 1.42 , +0.15 ± 0.02 GPa for Zr 0⋅68 Cr 0⋅32 B 1.38 , +0.04 ± 0.04 GPa for Zr 0⋅64 Cr 0⋅36 B 1.30 , and − 0.53 ± 0.02 GPa for Zr 0⋅56 Cr 0⋅44 B 1.11 . Fig. 4 shows the nanoindentation hardnesses H and elastic moduli E of as-deposited layers grown on Al 2 O 3 (0001) substrates as a function of x. The hardness of ZrB 2.18 is 31.8 ± 1.0 GPa, and increases to 41.7 ± 1.2 GPa for Zr 0⋅77 Cr 0⋅23 B 1.52 and 41.6 ± 0.9 GPa for Zr 0⋅71 Cr 0⋅29 B 1.42 , which is primarily due to their high compressive stress, solid-solution hardening [45], and narrow column widths (Hall-Petch effect [46,47] 1.38 , and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films grown on WC-Co substrates are evaluated by Revescratch tests. Optical microscope images from the scratch tracks, together with corresponding SEM micrographs acquired from the regions indicated by dashed and solid boxes in the optical microscope images, are exhibited in Fig. 6. The optical microscope images show that all films follow a similar scratch-failure mode; starting with chips spallation on the side of tracks, then wedge spallation, and eventually substrates exposure, which are the common failure modes observed for hard thin films [50,51]. The minimum load at which peeling and spallation occurs, referred to as the critical load (L c2 ) [52], is considered as the representative of adhesive failure, i.e. film delamination and spallation. The ZrB 2.18 film exhibits a poor adhesion together with severe chipping and buckling along its scratch track, with L c2 =~29 N, due to its high brittleness. Although the L c2 value of the Zr 0⋅77 Cr 0⋅23 B 1.52 alloys (~28 N) is almost similar to that of ZrB 2.18 , it increases significantly to ~42 and ~49 N for Zr 0⋅68 Cr 0⋅32 B 1.38 and Zr 0⋅56 Cr 0⋅44 B 1.11 , respectively. The SEM images from the regions indicated with dashed boxes, at distances between ~1.0 mm and ~1.1 mm (~30 N), in the optical microscope images of ZrB 2.18 and Zr 0⋅68 Cr 0⋅32 B 1.38 reveal angular cracks (indicated with black arrows) appeared close to the scratch tracks, which form due to their tensile stresses [51,53]   Zr 0⋅56 Cr 0⋅44 B 1.11 alloys obtained from their L c2 regions, indicated by blue solid boxes in their optical microscope images, consist of chipping debris and transverse semicircular cracks appeared in the scratch tracks, an indication of plastic deformation [51].
The broad wear track of ZrB 2.18 , Fig.7(a), shows a higher material loss (i.e. higher wear rate) occurring during the wear test with a typical wear caused by plastic deformation. There is a decrease in the width of wear tracks as a function of Cr concentration in the alloys. Compared to ZrB 2.18 , the adhesive wear is the primary wear mechanism of asdeposited Zr 1-x Cr x B y thin films, Figs. 7(b), (c), and 7(d). The adhesion of the alloys to the GCr15 steel produces a high material loss of the friction ball, instead of the alloys, which reduces the wear rate. The  0.522 ± 0.007 0.6 ± 0.1 enhanced wear resistance can be attributed to the combination of high hardness (>30 GPa) and increased toughness. Out of all compositions investigated, Zr 0⋅56 Cr 0⋅44 B 1.11 alloys are chosen for further oxidation and corrosion studies as they have metallicglass structure, relatively low residual stress, good hardness and toughness, and the highest wear resistance. Fig. 8 compares the XSEM images of ~2800-nm ZrB 2.18 and ~2100-nm Zr 0⋅56 Cr 0⋅44 B 1.11 thin films annealed in air at 700 • C for the time t a of 1, 3, and 5 h. The thickness of the oxide scale on ZrB 2.18 increases from 830 ± 50 nm for t a = 1 h, to 2620 ± 80 nm for t a = 3 h, and 3460 ± 90 nm for t a = 5 h. The oxide-scale thickness changes linearly as a function of oxidation time (d ox = 708⋅t a + 135). However, the oxide scales formed on the Zr 0⋅56 Cr 0⋅44 B 1.11 alloys are significantly thinner than those on ZrB 2.18 over the entire t a range. The thickness of the oxide scale on Zr 0⋅56 Cr 0⋅44 B 1.11 increases from 350 ± 30 nm for t a = 1 h, to 550 ± 50 nm for t a = 3 h, and 665 ± 55 nm for t a = 5 h, following a d ox = 352.4⋅t 0.4 a relationship. The enhanced oxidation resistance observed for the Zr 0⋅56 Cr 0⋅44 B 1.11 alloys is attributed to their elemental composition and nanostructure. The TMB 2 oxidation, which is mainly influenced by the evaporation rate of B 2 O 3 (g) phase, largely depends on the oxygen partial pressure, annealing temperature, and B concentration [11,54]. At constant oxygen partial pressure and annealing temperature, the vapor pressure of the B 2 O 3 (g) phase increases as a function of B concentration that results in decreasing the oxidation resistance [54]. We recently showed that sputter-deposited columnar TiB 2.4 thin films, in which the excess B segregates to the column boundaries, are highly prone to continuous vigorous oxidation in air [16]. The B 2 O 3 (g) phase preferentially forms at the column boundaries, which are B-rich, during annealing at temperatures above 400 • C. The evaporation of this phase, together with the coarsening of TiO 2 (s), lead to the formation of large gaps between the TiO 2 (s) columns that act as wide channels for oxygen to readily access the unoxidized regions; consequently, causing a continuous oxidation [16]. Hence, the higher oxidation resistance of Zr 0⋅56 Cr 0⋅44 B 1.11 can be explained by its very-low B concentration and amorphous structure, where the alloy does not have the B-rich column boundaries that are susceptible to preferential oxidation.

Oxidation properties
In addition, Lee et al. [28] showed that the oxidation resistance of Ti 1-x Cr x N films, isothermally annealed from 700 to 1000 • C in air, increases as a function of Cr concentration. The oxide scales formed on these alloys mainly consist of TiO 2 (s) and Cr 2 O 3 (s) phases [27,28]. As the Cr 2 O 3 (s) phase has a significantly lower coarsening rate than TiO 2 (s) [55], alloying TiN films with Cr decreases the coarsening of the oxide scale, which leads to suppressing the porosity formation and hence, decreasing the oxygen diffusion through the scale. This results in enhancing the oxidation resistance of Ti 1-x Cr x N films. Similar effect can be expected for the Zr 0⋅56 Cr 0⋅44 B 1.11 alloys.

Corrosion properties
The corrosion properties of as-deposited ZrB 2.18 and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films grown on WC-Co substrates are obtained by electrochemical measurements during the immersion of the layers in the 0.1 M NaCl corrosive medium for 25 h, at room temperature and without agitation. The open circuit potential (E ocp ), linear polarization resistance (R p ), and potentiodynamic polarization curves of ZrB 2.18 and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films are shown in Fig. 9. The electrochemical data determined from the polarization curves are summarized in Tables 3 and 4. Fig. 9(a) [56]. Fig. 9(b) exhibits the R p values of as-deposited ZrB 2.18 and Zr 0⋅56 Cr 0⋅44 B 1.11 thin films determined at times ranging from 0.25 to 24.0 h. The alloys have significantly higher R p values than the ZrB 2.18 films (see Table 3); the R p value of Zr 0⋅56 Cr 0⋅44 B 1.11 obtained after a 24.0h immersion in the corrosive medium is about twelve times higher than that of ZrB 2.18 (2.2 ± 0.3 MΩ cm 2 for ZrB 2.18 and 27.3 ± 2.9 MΩ cm 2 for Zr 0⋅56 Cr 0⋅44 B 1.11 ). Comparing the potentiodynamic polarization curves acquired after 25.0 h, Fig. 9(c), indicates that the Zr 0⋅56 Cr 0⋅44 B 1.11 alloys have lower corrosion current densities (i corr ) than the reference ZrB 2.18 films. The i corr value decreases from (4.6 ± 2.0) × 10 − 6 mA/cm 2 for ZrB 2.18 to (0.54 ± 0.1) × 10 − 6 mA/cm 2 for Zr 0⋅56 Cr 0⋅44 B 1.11 at the E corr values of − 250 ± 49 mV and − 180 ± 59 mV, respectively. This indicates that the corrosion rate of Zr 0⋅56 Cr 0⋅44 B 1.11 is almost nine times lower than that of ZrB 2.18 . The anodic polarization curves of both films consist of passive regions (from ~450 to ~1050 mV), where the one for Zr 0⋅56 Cr 0⋅44 B 1.11 is to some extent metastable, but at lower anodic current densities compared to that for ZrB 2.18 . The passive current density i pass is (1.3 ± 0.3) × 10 − 3 mA/cm 2 for ZrB 2.18 and (0.6 ± 0.1) × 10 − 3 mA/cm 2 for Zr 0⋅56 Cr 0⋅44 B 1.11 .
Alloying directly influences on the corrosion properties of materials by changing their nobility [56][57][58]. The lower E ocp value obtained for Zr 0⋅56 Cr 0⋅44 B 1.11 demonstrates that the electrochemical stability of these alloys is higher than that of ZrB 2.18 as Cr is a more noble element. The other factors that effectively change the corrosion resistance are the column boundaries and their density [59,60]. The column boundaries of ZrB 2.18 are more prone to corrosion attack than inside the columns, due to heterogeneity in their structure and chemistry (e.g. the B-rich phase). Thus, the absence of column boundaries for amorphous Zr 0⋅56 Cr 0⋅44 B 1.11 may contribute to a better corrosion resistance compared to that for polycrystalline ZrB 2.18 [61].

Conclusions
We demonstrate control of the composition, nanostructure, and properties of ZrB 2 -rich Zr 1-x Cr x B y films grown by hybrid Cr-HiPIMS/ ZrB 2 -DCMS co-sputtering. The reference ZrB 2.18 layers are deposited by DCMS with a negative dc substrate bias of 100 V. For the Zr 1-x Cr x B y alloy growth, the ZrB 2 target is continuously sputtered by DCMS, while the Cr magnetron is operated in HiPIMS mode providing pulsed Cr-ion fluxes. The Cr metal fraction, Cr/(Zr + Cr), is increased from x = 0.23 to x = 0.44 by decreasing the power P ZrB2 applied to the DCMS ZrB 2 target from 4000 to 2000 W in 500-W increments, while the average power, pulse width, and frequency applied to the HiPIMS Cr target are maintained constant at 700 W, 50 μs, and 100 Hz, respectively. Concurrently, y decreases from 2.18 to 1.11 as a function of P ZrB2 , due both to the addition of Cr (primarily) and preferential B resputtering. The energetic Cr-ion bombardment increases the density of the alloys and causes renucleation of the column growth. As a result, there is a refinement of the columnar structure with increasing the Cr concentration accompanied by increasing hardness to ~42 GPa for Zr 0⋅77 Cr 0⋅23 B 1.52 and Zr 0⋅71 Cr 0⋅29 B 1.42 . However, the further increase of Cr concentration leads to a significant B deficiency that results in the collapse of the hexagonal AlB 2 -structure into amorphous dense alloys, as revealed by XRD, TEM, and SAED patterns, with hardness values above 30 GPa.
The changes in the composition and nanostructure result in enhanced toughness and wear properties. The Zr 0⋅56 Cr 0⋅44 B 1.11 alloys, with the highest Cr concentration, exhibit considerably better toughness and wear resistance compared to ZrB 2.18 . The wear rate decreases from   − 180 ± 59 0.54 ± 0.1 0.6 ± 0.1 7.9 × 10 − 16 m 3 /(Nm) for ZrB 2.18 to ~0.6 × 10 − 16 m 3 /(Nm) for Zr 0⋅56 Cr 0⋅44 B 1.11 . In addition, these alloy films exhibit significantly higher oxidation and corrosion resistance. The thickness of oxide scale formed after air-annealing at 700 • C for 5.0 h markedly decreases from ~3460 nm for ZrB 2.18 to ~665 nm for Zr 0⋅56 Cr 0⋅44 B 1.11 . The corrosion rate of Zr 0⋅56 Cr 0⋅44 B 1.11 is about nine times lower than ZrB 2.18 . The Zr 0⋅56 Cr 0⋅44 B 1.11 alloys with the structural characteristics of metallicglass thin films show simultaneously several enhanced properties, which are essential for many strategic applications.

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.