Characterization of CrN-Based Hard Coating Materials with Addition of GaN
2018 Volume 59 Issue 10 Pages 1574-1577
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2018 Volume 59 Issue 10 Pages 1574-1577
Cr1−xGaxN thin films with various GaN concentrations were prepared on Si(100) substrates by pulsed laser deposition in order to clarify the effects of the GaN content on the material characteristics. The compositions of these films were determined by Rutherford backscattering spectroscopy, while crystal structures were elucidated using Fourier transform infrared spectroscopy and X-ray diffraction, and hardness values were measured by nanoindentation. Analyses determined that x was in the range of 0 to 0.51 and, at x ≤ 0.31, a single B1-(Cr,Ga)N phase was present. In those films for which x ≥ 0.38, a secondary phase based on B4-GaN also appeared. The hardness increased with increases in x up to 0.31 as the thin films maintained a single B1-(Cr,Ga)N phase, and a maximum hardness of 29.4 GPa was obtained at x = 0.31.
CrN-based hard coating materials have been widely applied to cutting tools to improve their useable lifespans and cutting performances.1–4) There have been many reports regarding improvements in the hardness and other characteristics of these materials following the addition of a third element to CrN. Typical examples include Cr–Si–N5–7) and Cr–Al–N8–10) coatings. It has been determined that Cr–Si–N exhibits hardening as a result of the formation of a nanocomposite structure11) in which nanocrystalline(nc)-CrN is surrounded by a-SiNx. In contrast, Cr–Al–N forms (Cr,Al)N as a metastable phase via the partial substitution of CrN having a B1(NaCl type) structure with AlN having a B1 structure.12) AlN undergoes a phase transition from a B4 (würtzite type) structure to a B1 structure at high pressures (approximately 23 GPa).13,14) The B1-AlN dissolves in the B1-CrN through a non-equilibrium process, thus improving the hardness by solution hardening.9,10) Our own group has investigated new hard coating materials based on solution hardening, focusing on Ga added as a third element to CrN. Like Al, Ga is a group 13 element and its nitride, GaN, has characteristics similar to AlN.15,16) GaN also shows a phase transition from a B4 to B1 structure at approximately 53 GPa.16–18) In our previous work, we confirmed that B1-GaN can be dissolved to form B1-CrN as a result of the epitaxial growth of a Cr–Ga–N thin film having a single B1 type phase on a MgO(100) substrate by pulsed laser deposition (PLD) using a Cr-10 mol% Ga alloy as a target material. Consequently, the GaN level in the thin film was only 10 mol%.19) In the case of Cr–Al–N thin films, the AlN content has a significant effect on their characteristics.20,21) The hardness of Cr–Al–N thin films has been found to increase with increasing AlN solid solution amounts up to a level of approximately 70 mol%. Above this concentration, the films consist of two phases, B1-(Cr,Al)N and B4-AlN, so that their hardness decreases drastically. However, in the case of Cr–Ga–N thin films, the effect of the GaN concentration on the mechanical properties has not been clarified. Therefore, in the present study, Cr1−xGaxN thin films with various GaN levels (as indicated by x) were prepared on Si(100) substrates by PLD and the impact of varying the amount of GaN was investigated.
Cr1−xGaxN thin films were prepared on Si(100) substrates by PLD, using a Cr–50 mol% Ga alloy attached to a Cr plate as the target. The substrate and the target were positioned 50 mm apart in a chamber that was evacuated to a pressure of less than 1.0 × 10−5 Pa. A nitrogen plasma was generated by ionizing a 1 sccm flow of N2 gas using a 400 W radio frequency radical source. During deposition, substrates were maintained at a temperature of 473 K using an infrared lamp. An ablation plasma was produced by applying a Nd:YAG laser (λ = 355 nm) in the form of intense pulses over brief durations of 7 ns at a repetition rate of 10 Hz to the rotating target. The GaN content in the thin films was tuned by varying the area of the Cr plate covered by the target (SR) from 0% to 100%. The deposition was carried out for 12 h, and the resulting film thicknesses were in the range of 200 to 300 nm. The composition of each film was assessed using Rutherford backscattering spectroscopy (RBS). During these analyses, the film was irradiated with a He2+ ion beam at an acceleration voltage of 2 MeV. The elemental proportions in the films were determined based on simulations performed with the RUMP software package that applied fittings to the RBS spectra. The bonding states of the films were evaluated by Fourier transform infrared spectroscopy (FT-IR; JASCO FT/IR-4000), acquiring spectra from 375 to 1425 cm−1 at a resolution of 4 cm−1. Crystal structures were studied by X-ray diffraction (XRD; Rigaku RINT 2500HF+/PC) in the Bragg-Brentano configuration with Cu Kα radiation (λ = 0.154 nm). Indentation hardness values were obtained by nanoindentation testing with a Berkovich indenter (Agilent Technologies G200), utilizing the continuous stiffness measurement (CSM) technique. Using this method, indentation depth profiles of both hardness and Young’s modulus could be obtained with a single indentation cycle.22) For each thin film, 15 points were tested to an indentation depth of 100 nm.
Table 1 summarizes the compositions of the Cr1−xGaxN thin films as determined by RBS. Each film contained approximately 50 mol% N, in good agreement with the stoichiometric expectation for CrN and GaN. The x values (representing the Ga/(Cr + Ga) ratio) ranged from 0 to 0.51 in conjunction with SR values from 0% to 100%. Figure 1 shows the FT-IR spectra obtained from the Cr1−xGaxN thin films. It has been reported that B1-CrN and B4-GaN generate broad peaks at approximately 400 cm−1 (due to Cr–N bonding) and 570 cm−1 (due to Ga–N bonding), respectively.23–25) The spectra of the films for which x ≤ 0.31 exhibit a broad peak at around 400 cm−1 due to the presence of B1-CrN. In the case of the film having x = 0.31, this peak is shifted to a lower wavenumber as compared to the film with a lower x. This would be expected based on changes in the bond length with dissolution of GaN in the CrN lattice. In contrast, the Cr1−xGaxN thin films for which x ≥ 0.38 generated an additional peak at 570 cm−1 as a result of the appearance of B4-GaN, indicating that the GaN was not fully dissolved in the CrN. Figure 2 provides the XRD patterns of films with various x values, along with the patterns for B1-CrN and B4-GaN from the International Centre for Diffraction Data (ICDD). The thin films with x ≤ 0.31 produced peaks attributable to B1-CrN with orientation along the (200). In contrast, those films with x ≥ 0.38 also generated peaks at 2θ values of 34.3° and 36.3° due to the appearance of a secondary phase. Although these peak positions do not exactly coincide with the B4-GaN (002) and (101) peaks, it was concluded they can still be ascribed to the phase based on B4-GaN, considering the FT-IR results. The appearance of the secondary phase was caused by GaN, which could not dissolve in the B1 phase, adopting the B4 phase as the stabilized phase for GaN. The (200) peaks also broadened and shifted to lower 2θ values with increasing x. Figure 3 plots the crystallite size in the B1 phase in the thin film (as calculated from the (200) peaks using Scherrer’s equation) as a function of x. The crystallite growth of the B1 phase was prevented by the increasing of GaN content in the thin films. This was probably caused by forcibly dissolving of B1-GaN, which is a high-pressure phase, in CrN. The crystallite size for the thin films with x ≥ 0.31 decreased drastically. Since the thin film with x = 0.38 has the secondary phase based on B4-GaN, x = 0.31 was close to maximum solubility of B1-GaN into CrN. Therefore, it was considered that the crystallites of the B1 phase were hardly to grow for thin films with x ≥ 0.31. The lattice constant for the B1 phase, as calculated from the (200) peak positions in XRD patterns, is plotted as a function of x in Fig. 4. This constant increased by approximately 1.4% as x increased from 0 to 0.31, as a result of the dissolution of GaN in the CrN to form a B1-(Cr,Ga)N phase. As x was increased to 0.31 or more, the lattice constant plateaued. We believe that those films for which x ≥ 0.38 consisted of B1-(Cr,Ga)N supersaturated with GaN and the secondary phase based on B4-GaN. There is a possibility that Cr dissolved to the secondary phase, forming B4-(Ga,Cr)N.
FT-IR spectra of Cr1−xGaxN thin films having various x values.
XRD patterns of Cr1−xGaxN thin films having various x values. FWHM indicates full width at half maximum of the (200) peaks.
Crystallite size in the B1 phase (as calculated using Scherrer’s equation) as a function of the relative proportion of Ga.
Lattice constant of the B1 phase, as calculated from the (200) peak position in each XRD pattern, as a function of the relative proportion of Ga.
Typical indentation hardness and Young’s modulus depth profiles obtained during a single indentation test with the CSM technique are shown in Fig. 5. At depths ranging from 0 to 40 nm, the hardness initially increased rapidly and then became constant, after which it decreased slightly with increasing depth beyond 40 nm. This decrease in the hardness could have resulted from the Si substrates because the film thickness were all from 200 to 300 nm. Herein, the intrinsic hardness and Young’s modulus values were calculated by averaging over indentation depths from 30 to 40 nm. Figure 6 plots the intrinsic hardness and Young’ modulus values obtained by averaging data from 15 indentations for each sample. Both hardness and modulus increased with increases in x up to 0.31 as the thin films maintained a single B1-(Cr,Ga)N phase. At x = 0.31, the maximum hardness of 29.4 GPa was obtained. This trend could have been caused by solution hardening of B1-GaN to give B1-CrN and by a reduction in the crystallite size, as evident in Fig. 3. In addition, increases in the Young’s modulus could also contribute to the hardening. Simultaneous with the appearance of B4 type phase in the thin films, the hardness and Young’ modulus both decreased. There were some reports that Vickers and indentation hardness value of the B4-GaN showed approximately 12 GPa26) and 20 GPa,27) respectively. These values were lower than that of CrN fabricated in this study. Similar to Cr–Al–N thin films, the presence of a B4 type phase in the thin films prevented hardening.
Typical depth profiles for the indentation hardness and Young’s modulus values obtained from Cr1−xGaxN thin film with x = 0.31.
Indentation hardness and Young’s modulus values of Cr1−xGaxN thin films as functions of the relative proportion of Ga.
Cr1−xGaxN thin films for which x ≤ 0.51 were successfully prepared on Si(100) substrates by PLD. The films with x ≤ 0.31 had a single phase with a B1 structure, and the lattice constant of these films was found to increase with increasing x. From these results, it was determined that GaN was dissolved in the CrN lattice, forming a B1-(Cr,Ga)N phase. Conversely, films for which x = 0.38 or 0.51 consisted of two phases: B1-(Cr,Ga)N supersaturated with GaN and a secondary phase based on B4-GaN. The nanoindentation hardness and Young’s modulus values of the Cr1−xGaxN thin films increased with increases in x up to 0.31, giving a maximum hardness of 29 GPa. As the B4-GaN appeared in the thin films, the indentation hardness decreased. Based on the above results, GaN addition up to approximately 30 mol% can contribute to the hardening of CrN.
