Interfacial stabilization for epitaxial CuCrO2 delafossites

ABO2 delafossites are fascinating materials that exhibit a wide range of physical properties, including giant Rashba spin splitting and anomalous Hall effects, because of their characteristic layered structures composed of noble metal A and strongly correlated BO2 sublayers. However, thin film synthesis is known to be extremely challenging owing to their low symmetry rhombohedral structures, which limit the selection of substrates for thin film epitaxy. Hexagonal lattices, such as those provided by Al2O3(0001) and (111) oriented cubic perovskites, are promising candidates for epitaxy of delafossites. However, the formation of twin domains and impurity phases is hard to suppress, and the nucleation and growth mechanisms thereon have not been studied for the growth of epitaxial delafossites. In this study, we report the epitaxial stabilization of a new interfacial phase formed during pulsed-laser epitaxy of (0001)-oriented CuCrO2 epitaxial thin films on Al2O3 substrates. Through a combined study using scanning transmission electron microscopy/electron-energy loss spectroscopy and density functional theory calculations, we report that the nucleation of a thermodynamically stable, atomically thick CuCr1−xAlxO2 interfacial layer is the critical element for the epitaxy of CuCrO2 delafossites on Al2O3 substrates. This finding provides key insights into the thermodynamic mechanism for the nucleation of intermixing-induced buffer layers that can be used for the growth of other noble-metal-based delafossites, which are known to be challenging due to the difficulty in initial nucleation.


Scientific Reports
| (2020) 10:11375 | https://doi.org/10.1038/s41598-020-68275-w www.nature.com/scientificreports/ high-quality CuCrO 2 thin films cannot be simply understood from conventional thin film growth mechanisms; therefore, a direct observation of the atomic and electronic structure of the CuCrO 2 /Al 2 O 3 interface is required to reveal the underlying reason for how the epitaxy of CuCrO 2 on Al 2 O 3 is accomplished despite the relatively large lattice mismatch.
In this study, we grew high-quality CuCrO 2 thin films by systematically tuning the growth conditions, including the growth temperature (T) and oxygen partial pressure (P O2 ). Our films revealed a high crystallinity, smooth surface, and reasonably high resistivity. Using these high-quality CuCrO 2 thin films, we studied the interface microstructure to understand the nucleation and growth behavior of CuCrO 2 using scanning transmission electron microscopy (STEM)/electron energy loss spectroscopy (EELS) and density functional theory (DFT) calculations. We have found that atomic-level interfacial intermixing between Al and Cr atoms within the atomiclayer thick substrate surface plays a critical role in stabilizing the nucleation of the CuCrO 2 delafossite phase. The initial intermixing-induced nucleation seems important to both reduce the most stable impurity phase, Cr 2 O 3 , and to stabilize the high-quality CuCrO 2 phase.
To grow CuCrO 2 thin films with high crystallinity, we mapped out the optimal growth condition for CuCrO 2 thin films on Al 2 O 3 (0001) substrates by varying the temperature and oxygen partial pressure using a single-phase CuCrO 2 target (see Figure S1 in Supplementary Information). Figures S1a,b show X-ray diffraction (XRD) 2θ-θ scans for CuCrO 2 films grown under different T and P O2 conditions. The CuCrO 2 phase could be stabilized under a wide range of growth conditions, but an impurity phase was observed under both low P O2 (< 0.01 mTorr) and high T (> 800 °C) growth conditions. Figure 1a summarizes results for CuCrO 2 films grown at different T and P O2 . The contour plot indicates rocking curve full width at half maximum (FWHM) values of the 0006 CuCrO 2 peak, and the symbols indicate whether the film is single-phase (black circles) or has impurity phases (blue stars). This result indicates that the growth window for the epitaxy of CuCrO 2 films is relatively wide (500 < T < 800 °C and 0.01 < P O2 < 500 mT). We found the best quality films were grown at T = 650 °C and P O2 = 10 mTorr. Figure 1b shows X-ray reflectivity (XRR) and XRD patterns of CuCrO 2 thin films with different thicknesses (d = 3.1-44 nm) grown on Al 2 O 3 (0001) at the optimum condition (T = 650 °C and P O2 = 10 mTorr). XRR results of all films show clear interference fringes, indicating smooth surfaces of CuCrO 2 thin films. All of the film peaks in the XRD patterns correspond to the delafossite 0003n peaks. The width of the CuCrO 2 0003n peaks become broader with decreasing film thickness, as expected from the Laue function. As shown in Fig. 1c, the rocking curve FWHM values for the 0006 peak of CuCrO 2 films are ~ 0.1°, which is smaller than those in previous reports. Figure 1d shows an atomic-force microscopy (AFM) image of a CuCrO 2 thin film surface, showing a triangular shaped grain boundary (we note that such a grain boundary can act as a scattering center in the carrier relaxation process, yielding a higher resistivity than materials without such disorder). The root mean www.nature.com/scientificreports/ square (RSM) roughness of our film was estimated to be 1.58 nm over a 3 × 3 μm 2 range. The RSM roughness of a structural domain was only 0.25 nm, which is much smaller than that from spin-coated 12,13 and MBE-grown films 14 (RSM = ~ 3-50 nm). The RSM value is even smaller than that from previously PLE grown CuCrO 2 thin films (RSM = ~ 1 nm) 15 . Figure 2a shows the temperature dependence of 4-probe dc resistivity for a CuCrO 2 thin film (11.5 nm), which exhibits semiconducting behavior (dρ/dT < 0). The thermal activation energy of charge carriers was ~ 97 meV as shown in Fig. 2b, which is consistent with previous reports from thin film samples [16][17][18] . Overall, the epitaxial growth of high-quality CuCrO 2 thin films is particularly notable if we consider that the growth of delafossite thin films without impurity phases is a big challenge in many other compounds, e.g. PdCrO 2 , PdCoO 2 , and PtCoO 2 . Figure 3a shows a high-angle annular dark field (HAADF) STEM image of a CuCrO 2 thin film grown on an Al 2 O 3 (0001) substrate seen along the [ 1100 ] zone axis. The brightest and second brightest features in this image indicate Cu and Cr atomic columns, respectively. Note the HAADF STEM provides scattering intensities that are approximately proportional to the square of the atomic number. Thus, the lightest element, O, is not visible in this HAADF STEM image. The HAADF image confirms that the CuCrO 2 thin film is epitaxially grown on the Al 2 O 3 (0001) substrate. The atomic structure at the interface, shown in Fig. 3a,b, exhibits several interesting aspects. First, the growth of the CuCrO 2 thin film initiates with the CrO 2 sublayer, followed by the Cu sublayer. Second, the top one or two layers of the Al 2 O 3 substrate exhibit a brighter intensity than the bulk layers of the substrate. In this study, a monolayer (ML) of each material was defined to satisfy stoichiometry. That is, the MLs for CuCrO 2 thin film were composed of a set of Cu and CrO 2 sublayers (thickness: 0.57 nm) and for Al 2 O 3 substrate in the (0001) direction as a single Al 2 O 3 layer (thickness: 0.22 nm), which are illustrated in Fig. 4b. Thus the increased intensity implies that some atomic-level interfacial intermixing occurred during the initial stage of film growth. Although it was not identified previously, these interfacial features were also similarly observed in a PdCrO 2 thin film grown on a CuCrO 2 -buffered Al 2 O 3 (0001) substrate 9 . In addition, the HAADF STEM image shows that stacking faults exist in the delafossite CuCrO 2 thin film, denoted by black lines on the left side of Fig. 3a. Such stacking faults were frequently observed in previous PLE or MBE grown delafossite films 9,10 . The presence of these stacking faults indicates that their formation energies are relatively small in this delafossite film. Further studies will be required to understand the influence of stacking faults on the optical and transport properties of delafossite materials.
To systematically understand the interface structure and the nucleation of epitaxial growth of CuCrO 2 thin films, EELS spectrum imaging was performed across the CuCrO 2 /Al 2 O 3 interface as shown in Fig. 3c-g. Previous studies have shown that the direct interpretation of interface structure can be achieved only under very thin specimen conditions (ideally less than approximately 20 nm) 19,20 . Otherwise, complex propagation effects, which include beam broadening, cross-talk, and dechanneling, result in the potential misinterpretation of interface structure 21,22 . For this study, we selected a thin region of a specimen for atomic-scale quantification of the interface structure. Thickness measurements across the interface using low-loss EELS spectra are shown in Figure S2 of the Supplementary Information. We confirmed by the 2D spectrum line profile that core-loss excitations are sensitive to the individual atomic planes (see Figure S3 in Supplementary Information). The O-K and Cr-L 2,3 edge signals dropped to almost zero at the Cu sublayers, as did the Cu-L 2,3 edge at the CrO 2 sublayers. Figure 3c shows a HAADF STEM image of the CuCrO 2 /Al 2 O 3 interface, which is simultaneously acquired during EELS To minimize specimen damage due to electron beam irradiation, EELS spectrum imaging was performed using a low-current electron probe with a short exposure time (0.01 s per pixel). Even in the thinnest specimen section, both Cr and Al elements are clearly detected across the interface as shown in Fig. 3d,e, suggesting interlayer mixing occurred at the interface during the film growth. We note that the plume energy in PLE growth is quite high (1-100 eV) 23 , exceeding the surface bonding energy of substrate materials (typically on the order of 1 eV). Thus, the growth species can easily penetrate into the substrate surface 23,24 , resulting in intermixing of elements especially under high vacuum conditions. Interestingly, the concentration of Cr was significantly reduced only for the first ML of the CuCrO 2 thin film. The stoichiometry of the CuCrO 2 layer appears to be fully recovered from the second ML of the film, indicating that the interfacial inter-layer-mixing in the thin film largely occurred within one ML. The Cr-L 2,3 edge spectra in Fig. 3f further confirmed that Cr atoms penetrated up to two MLs below the interface, and the oxidation state of Cr ions was maintained as Cr 3+ even in the Al 2 O 3 substrate side (the shape and energy of the Cr-L 3 edge do not change across the interface 25,26 ). . It is worth noting that no discernible vacancy-related features could be detected from the integrated line-profile spectra.
Scientific Reports | (2020) 10:11375 | https://doi.org/10.1038/s41598-020-68275-w www.nature.com/scientificreports/ Meanwhile, the Cu-L 2,3 edge was not detected on the substrate side, verifying that Al atoms were intermixed with the B-site Cr atoms not Cu atoms. In general, the O-K edge in transition metal oxides has been used for the investigation of the individual electronic structure of materials, since the 1s core states have a relatively small exchange interaction with the final states, resulting in no visible multiplet effects [27][28][29] . The O-K edge spectra profile across the interface confirms that one ML above the interface and two MLs below the interface exhibits the characteristic electronic structures from those of the CuCrO 2 thin film and Al 2 O 3 substrate, respectively. More interestingly, the O-K edge of the first ML of the CuCrO 2 thin film (spectrum #6) revealed the same spectral signature as that of the CuAlO 2 delafossite, which is shown in Fig. 3g. The overall shape of the O-K edge of the first ML is similar to that of CuAlO 2 30 , and, in particular, the peak indicated by the black arrow, which cannot be explained by the O-K edge spectra of the existing CuCrO 2 and Al 2 O 3 , only appears in the O-K edge of the first ML. The in-depth analysis of EELS spectral images strongly suggests that the interfacial intermixing, which occurred during the initial stage of the epitaxial growth, creates an atomically thin ML of CuCr 1−x Al x O 2 alloy.
To evaluate the thermodynamic stability of delafossites compared to the Cr 2 O 3 impurity phase, the formation enthalpy (ΔH f ) was calculated using the DFT-based fitted elemental-phase reference energies (FERE) method 31 . As shown in Fig. 4a, ΔH f of Cr-based delafossites, including CuCrO 2 and PdCrO 2 , were far greater than that of Cr 2 O 3 , accounting for the formation of Cr 2 O 3 impurities. Interestingly, CuAlO 2 was found to be much more stable than CuCrO 2 and even more stable than Cr 2 O 3 ; 0.08 eV/atom lower than that for Cr 2 O 3 . These numerical results suggest that the substitution of Al atoms for Cr atoms in CuCrO 2 will lower its formation enthalpy. This thermodynamic consideration explains the formation of the CuCr 1−x Al x O 2 interfacial ML we observed in our STEM-EELS investigations. In addition, the interfacial mixing helps to reduce the epitaxial strain as the lattice constant of CuCr 1−x Al x O 2 is closer to the Al 2 O 3 substrate than for pure CuCrO 2 . It is worth noting that the formation of PdAlO 2 has not been experimentally reported so far, suggesting that the substitution of Al atoms for Cr atoms to form the equivalent interfacial PdCr 1−x Al x O 2 for the growth of epitaxial PdCrO 2 on Al 2 O 3 is unlikely.
The STEM-EELS and DFT results provide direct insights into the nucleation of epitaxial growth of CuCrO 2 thin films and the critical role of a CuCrO 2 buffer layer for the epitaxy of delafossites. First, inter-layer-mixing was observed at the interface, but not further into the film, indicating that it happens at the initial nucleation step of the epitaxial growth. At the initial nucleation stage, Cr atoms will penetrate into the sub-surface-layers of the Al 2 O 3 substrate and, as an exchange, Al atoms will out-diffuse to the surface. Second, the free Al atoms at the surface will act to stabilize the nucleation of CuCrO 2 delafossite thin films. Without these Al atoms, the Cr 2 O 3 will be the most stable phase at the nucleation step, disturbing the formation of the delafossite phase. Third, the homogenous and stable nucleation, with the delafossite symmetry provided by CuCr 1−x Al x O 2 , will enable the high-quality growth of CuCrO 2 thin films. The growth process understood here is summarized in Fig. 4b. In a previous study, we reported that the use of a one-ML-thick CuCrO 2 buffer layer significantly suppressed the formation of Cr 2 O 3 impurities in the epitaxy of PdCrO 2 thin films 9 . It was quite puzzling because the lattice mismatch of the CuCrO 2 /Al 2 O 3 interface (δ = 7.2%) is much larger than that of the PdCrO 2 /Al 2 O 3 interface (δ = 5.9%), which is counter intuitive. Our discovery of the formation of the CuCr 1−x Al x O 2 delafossite at the CuCrO 2 /Al 2 O 3 interface now explains the role of the CuCrO 2 buffer layer in the growth of the PdCrO 2 thin films. The deposition of CuCrO 2 buffer layer will induce the homogenous and stable nucleation with delafossite symmetry, which cannot be achieved from the direct deposition of PdCrO 2 layers. Note again that the preferential nucleation by Al substitution only occurs for Cu-based delafossites, not for Pd-based delafossites; the substitution of Al atoms for Cr atoms acts to decompose PdCrO 2 delafossite films. www.nature.com/scientificreports/ In summary, we have grown high quality CuCrO 2 thin films by pulsed laser epitaxy. Compared with CuCrO 2 thin films grown by other methods, PLE grown films show better quality in terms of crystallinity and surface roughness. The successful growth was possible owing to the non-equilibrium energetic process of PLE growth. The intermixing-induced alloying of the Al and Cr atoms was found to play a crucial role in stabilizing the nucleation of CuCrO 2 delafossite phase and in reducing the most stable impurity phase Cr 2 O 3 often found in other delafossites. We believe a similar consideration can be also applied to Co-based delafossites as the formation of the Co 3 O 4 spinel structure has been also a challenge for, e.g., PdCoO 2 7,8,10,11 . Our results suggest that the key to achieving the layer-by-layer growth of CuCrO 2 delafossite films is the nucleation of the structurally similar CuCr 1−x Al x O 2 buffer layer at the interface when grown on structurally dissimilar substrates, e.g., conventional Al 2 O 3 substrates. Thus, this discovery may provide a critical strategy for the epitaxial growth of other delafossites with the new CuCr 1−x Al x O 2 buffer layer to accelerate innovations in future electronic and spintronic quantum devices made from delafossites.

Methods
Thin film growth. High quality CuCrO 2 thin films were grown on c-plane Al 2 O 3 substrates by PLD using a polycrystalline CuCrO 2 target. The polycrystalline CuCrO 2 is prepared by sintering the mixture of Cu 2 O and Cr 2 O 3 at 1,100 °C for 10 h in air. The obtained pure polycrystalline CuCrO 2 were pelletized and annealed at 800 °C. Before the thin film growth, commercially available Al 2 O 3 (0001) substrates (CrysTec, Germany) were annealed at 1,100 °C for 1 h to achieve atomically flat surfaces with step-terrace structure. For the CuCrO 2 film growth, the growth conditions were widely varied (T = 400-800 °C, P O2 = 0.01-500 mTorr), whereas the repetition rate and energy of the KrF excimer laser (λ = 248 nm) were fixed at 5 Hz and 1.5 J/cm 2 , respectively. The best CuCrO 2 epitaxial thin films were obtained at optimal growth conditions of T = 650 °C, P O2 = 10 mTorr. After the growth, the samples were cooled to room temperature in P O2 = 100 Torr.
Characterization. The crystal structure was characterized by X-ray diffraction (XRD) using a four-circle high-resolution X-ray diffractometer (X'Pert Pro, PANalytical; Cu Kα 1 radiation), and the thickness of the film (d) was calibrated using X-ray reflectivity (XRR). The surface morphology measurements were made with atomic force microscopy (Veeco Dimension 3100). Cross-sectional TEM specimens were prepared using low-energy ion milling at LN 2 temperature after mechanical polishing. HAADF STEM measurements were performed on Nion UltraSTEM200 operated at 200 kV. The microscope is equipped with a cold field emission gun and a corrector of third-and fifth-order aberration for sub-Angstrom resolution. The convergence half-angle of 30 mrad was used and the inner angle of the HAADF STEM was approximately 65 mrad. To minimize the electronirradiation damage during EELS mapping, EELS spectra were measured at 0.01 s collection time.
Estimation of the enthalpy of formation (ΔH f ) using ab-initio DFT calculations. Ab initio DFT calculations were performed using the Vienna ab initio simulation package (VASP) code 32 , and the enthalpy of compound formation is estimated using the fitted elemental-phase reference energies (FERE) method 31 . The Perdew-Burke-Ernzerhof plus Hubbard correction (PBE + U) was used for the exchange-correlation functional 33 , in which the double-counting interactions were corrected using the full localized limit (FLL) 34 . The Hubbard U parameter of 3 eV (U = 3 eV) is used for all transition metals except Cu for U = 5 eV following previous work 31 . A plane wave basis set at a cutoff energy of 600 eV was used to expand the electronic wave functions, and the valence electrons were described using the projector-augmented wave potentials. The Γ-centered 9 × 9 × 9 Monkhorst-Pack K-point grid was used for sampling the Brillouin zone. Input structures are obtained from the Inorganic Crystal Structure Database (ICSD) and all cells and atomic positions in our calculations were relaxed with the force criteria of 0.01 eV/Å. In the FERE approach, the enthalpy of formation (ΔH f ) of a chemical compound A n1 B n2 … is expressed by the following equation: where E DFT tot is the total energy per formula unit of a given compound, µ DFT i are the total energies per atom of the elements in their elemental reference phase, and δµ FERE i are the FERE correction energies of the elements. The δµ FERE i for 50 chemical elements are tabulated in the paper describing the FERE approach 31 .