Structural modifications in the reduced phase

Under the same conditions that lead to fully oxidized LaCrO 3 (ref. 7), the deposition of epitaxial P-SCO on (001)-oriented perovskites, such as LaAlO 3 (LAO) and (LaAlO 3 ) 0.3 –(SrAl 0.5 Ta 0.5 O 3 ) 0.7 (LSAT), by means of molecular beam epitaxy, results in R-SCO. Conversion of R-SCO to P-SCO occurs during mild heating in air. Evidence for this phase transition is given by X-ray diffraction (XRD), as seen in Fig. 1a and Supplementary Fig. 1a. All as-grown films exhibit a single Bragg peak at 2θ≈46° near the substrate (002) reflection, revealing epitaxy. These Bragg angles increase substantially upon annealing in air, revealing a sizeable drop in c lattice parameter. The changes in lattice constants (based on reciprocal space maps—see Supplementary Figs 1 and 2) are summarized in Fig. 1b. The in-plane lattice parameters for the as-grown films (3.82 Å on LAO and 3.86 Å on LSAT) are nearly the same as those of their respective substrates, whereas the out-of-plane lattice constant (c) is ~3.92 Å for both films on both substrates. The larger value of c in the as-grown films compared with that of bulk SrCrO 3 (3.819 Å) implicates the presence of V O through a phenomenon called chemical expansivity8,9. The drop in c upon air annealing is consistent with oxidation of R-SCO to P-SCO.

Figure 1: Topotactic phase transition of strontium chromite from rhombohedral SrCrO 2.8 (R-SCO) to cubic perovskite SrCrO 3-δ . (a) XRD θ–2θ scans of the as-grown R-SCO film on LAO(001) (red curve) that transforms into cubic perovskite SrCrO 3-δ after annealing in air at 250 °C for 2 h (blue curve). (b) Evolution of the in-plane and out-of-plane lattice constants for the epitaxial films obtained from reciprocal space maps around the (103) Bragg reflection of the substrates; circles and squares represent films grown on LAO and LSAT, respectively. (c) Cross-sectional high-angle annular dark field (HAADF)-scanning transmission electron microscopy (STEM) image as viewed along [110] for as-grown R-SCO film on LAO(001), showing (111)-oriented planes of ordered oxygen vacancy planes every five unit cells along [111], both in the image (dark stripes) and by the superlattice peaks in the FT (inset). (d) After annealing in air at 250 °C, the ordered vacancies are substantially eliminated, resulting in P-SCO (upper right inset) and the corresponding disappearance of the superlattice reflections (lower left inset). (e) HAADF-STEM image (averaged over many short exposures) revealing an ordered oxygen-deficient structure that is well matched to a 15R unit cell with periodic (111)-planes of Cr4+ tetrahedra, overlaid on a calculated structural model structure (see text for details). Full size image

The R → P conversion is clearly seen by cross-sectional, high-angle annular dark field scanning transmission electron microscopy (TEM). Domains with {111}-ordered planar superlattice structures are observed, seen as dark stripes repeating every five unit cells in Fig. 1c. The ordered {111}-plane domains generate a set of superlattice spots in the Fourier transforms (FT; inset in Fig. 1c). Figure 1e shows a more magnified lattice image, along with the structural model discussed below. The positions of visible Cr and Sr atomic columns are consistent with {111}-oriented SrO 2 layers separated by four SrO 3 layers. This ordered structure (a 15R polymorph) is the same as that of R-SCO that forms upon reduction of P-SCO powder10. As in ref. 10, we find that Cr adjacent to the SrO 2 planes is tetrahedral Cr4+, whereas the other three Cr layers located between SrO 2 planes contain Cr with an average charge of +3.33; these ions retain the octahedral coordination found in P-SCO. The distribution of Cr charge states throughout the structure was imaged by electron energy loss spectroscopy (EELS). EELS maps tracking the near-edge structure of the Cr L-edge, as shown in Supplementary Fig. 3, is consistent with Cr4+ being detected within the SrO 2 -containing stripes and Cr3+ being found between these stripes, based on a reported 1 eV chemical shift between Cr3+ in LaCrO 3 and Cr4+ in CrO 2 (ref. 11).

Based on this experimental data and ab initio calculations, it is natural to consider the crystal structure of R-SCO in terms of 〈111〉-oriented lamella consisting of , where and have charges of −1 and +2 with respect to the lattice. The packing of these lamella can be understood in terms of the electrostatic interactions between atomic planes. Ab initio calculations reveal that two such lamellae do not interact if separated by five or more Cr planes (Supplementary Fig. 4). The repulsion energy of two lamellae separated by only four Cr planes is ~0.45 eV per (1 × 1) lateral 15R cell, and it rapidly rises as the distance between the [SrO 2 ]2– planes decreases and lamellae begin to overlap. This result is consistent with the experimentally observed structure, and suggests that under reducing conditions the ordered oxygen-deficient layers pack as closely as possible without incurring a repulsive energy penalty, and implies that these 2D structures are mobile and can move along the 〈111〉 direction, either in aggregate or as isolated vacancies. Provided the spacing between lamellae (d 1 ) exceeds the spacing between a central SrO 2 layer and it nearest octahedral Cr3+ layer (d o ), there is no repulsion between the two lamellae.

Evidence for reversible reduction/oxidation processes

The majority of the oxygen-deficient planes are removed by annealing in air for 2 h at only 250 °C as, accordingly, R-SCO is transformed into P-SCO, as clearly seen in the scanning TEM image and FT (inset) in Fig. 1d. This phase transformation involves two processes: (1) insertion of oxygen atoms into SrO 2 layers to form SrO 3 , and (2) electron transfer from octahedrally coordinated Cr3+ ions located between the SrO 2 layers so as to convert the inserted oxygen atoms into O2– ions. As oxidation proceeds, XRD reveals a contraction in the out-of-plane lattice constant by ~2% (Fig. 1a,b). The (103) Bragg reflection for the air-annealed sample indicates that the lattice constant is close to that of the bulk P-SCO (Supplementary Fig.2). Therefore, we refer to the air-annealed sample as P-SrCrO 3-δ . It is difficult to determine the exact value of δ because O vacancies distribute randomly and their concentration is very low; our ab initio calculations indicate that δ should be ≤0.10 in order to form the P-SCO phase.

This oxidation process can be dynamically observed in situ with high spatial resolution using environmental TEM (ETEM). The vacancy planes are eliminated by annealing at 250 °C in 1.5 mbar oxygen gas within ~30 min. To ensure that the apparent oxidation is not a result of beam damage, the sample was imaged at the same temperature and pressure in Ar for ~45 min before introducing oxygen with no apparent changes observed during this time. In oxygen at 250 °C, the oxygen-deficient domains typically vanish more quickly than could be captured in our apparatus (<10 frames per second under imaging conditions). In order to slow the process, experiments were also performed by introducing a partial pressure of 0.05 mbar oxygen at 400 °C. The individual ordered oxygen-deficient SrO 2 planes still oxidize very quickly, but single domains sometimes heal in steps (Supplementary Movie 1). Incorporation of oxygen may go unnoticed until a sudden R-to-P phase transition occurs, as the relevant TEM contrast is more sensitive to the structural discontinuity than to the individual oxygen ions. The grouping of rapidly oxidized SrO 2 planes seems to correlate with the location of {111}-oriented structural defects formed at, for example, grain boundaries between mutually orthogonal {111}-ordered domains (Supplementary Fig. 5). This result suggests that the oxidation rate is limited by defect formation. As oxygen is incorporated into the SrO 2 layers, the local structure changes and the superlattice diffraction spots in the FT disappear. In addition, the out-of-plane lattice parameter decreases by 1.7%, and the in-plane parameter decreases by 0.8%. This result is consistent with XRD data and the results of our Perdew-Burke-Ernzerhoff (PBE) calculations for unconstrained SrCrO 3-δ , which suggest that the R-to-P structural transition is accompanied by an out-of-plane lattice parameter decrease of 3.2–3.6%, depending on the residual V O concentration, while the in-plane lattice parameter decreases by 1.1–1.5%.

Spectroscopy and the semiconductor-to-metal transition

Cr 2p core-level and valence band (VB) X-ray photoemission spectra, along with conductivity and optical absorption, reveal the expected changes in Cr valence and the electronic structure of the SCO lattice as a whole as the phase transition takes place. Figure 2a shows that for R-SCO, there is, in addition to a main peak at ~576.5 eV associated with Cr4+, a strong shoulder at ~575.2 eV attributed to Cr3+ (ref. 7). After air annealing, the intensity of this shoulder is significantly reduced as the P-SCO phase forms, corresponding to oxidation of Cr3+ to Cr4+ as the phase transition occurs. This change in Cr valence is accompanied by distinct changes in the electronic and optical properties of the film. Temperature-dependent resistivity ρ(T) data, shown in Fig. 2c and Supplementary Fig. 1c, indicate that R-SCO is semiconducting (dρ/dT<0), whereas P-SCO is metallic (dρ/dT>0). In addition, the resistivity of P-SCO (2 × 10−3 Ω cm) is almost three orders of magnitude smaller than that of R-SCO (6 × 10−1 Ω cm). This marked change in resistivity is accompanied by a colour change from transparent grey to opaque dark (Fig. 2d, inset). The optical absorption spectra for R-SCO films show no discernible absorption below hν=2.0 eV, consistent with its semiconductor behaviour. In contrast, P-SCO exhibits a strong absorption onset for hν less than ~1.0 eV, consistent with light scattering by carriers in the metallic phase.

Figure 2: Significant differences in electronic and optical properties for the R and P phases. (a) Cr 2p core-level spectra for as-grown (R-SCO) and air-annealed (P-SCO) films on LAO(001), showing that the Cr3+ contributions to both spin-orbit components are significantly reduced by air annealing. (b) X-ray-excited VB photoemission spectra for the as-grown and air-annealed films. Inset shows the spectra near the Fermi level (E F ), along with the spectrum from an Au foil in direct contact with the films for energy calibration purposes. There is little intensity at E F for as-grown films, whereas substantial intensity appears at E F for air-annealed films, indicating a semiconductor-to-metal transition. (c) Temperature-dependent resistivity ρ(T) curves for as-grown and air-annealed films, revealing semiconducting behaviour for R-SCO and metallic behaviour for P-SCO. (d) Optical absorption spectra for the as-grown and air-annealed films, as well as an LAO substrate. Full size image

The VB spectrum for R-SCO (Fig. 2b) is primarily of O 2p character in the region of 2.2–8.5 eV, and exhibits a Cr 3d-derived feature at ~1.0 eV. There is only very slight spectral intensity detected at the Fermi level (E F ) (Fig. 2b, inset), suggesting semiconducting behaviour. In contrast, pronounced intensity at E F is seen for P-SCO, revealing its metallic character. It is noteworthy that the physical properties of SrCrO 3 are controversial. A critically important issue is whether SrCrO 3 is metallic or insulating12,13,14. This controversy most likely exists because of poorly defined structure and/or composition in powder samples. Our experimental results provide definitive evidence that single-crystal stoichiometric SrCrO 3 is metallic. The metallic state of P-SCO is attributed to strong hybridization of the Cr 3d2 t 2g configuration with O 2p states, which results in delocalized holes at the top of the VB. Cr 3d-O 2p hybridization and hole delocalization are further supported by the presence of an intense pre-edge peak in the O K-edge EELS (Supplementary Fig. 6). This assignment is consistent with those for SrCoO 3 (ref. 15) and SrFeO 3 (ref. 16), as well as CrO 2 (ref. 17).

Remarkably, the P-SCO to R-SCO phase transition is fully reversible at moderately low temperature. P-SCO can be transformed back into R-SCO by annealing in vacuum at 500 °C for 1 h, as confirmed by XRD and ρ(T) data (Supplementary Fig. 7). Likewise, R-SCO can be re-oxidized to P-SCO by annealing in air at 250 °C. Using photoemission intensity at E F to monitor metallicity, the redox process was followed during in vacuo and air annealing. The results are shown in Fig. 3a. The count rate at E F starts to decrease as a result of vacuum annealing at 300 °C, and goes to zero at 500 °C after 1 h, revealing a metal-to-insulator transition. The Cr 2p core-level and VB spectra for the 500 °C vacuum-annealed sample are similar to those for the as-grown R-SCO (Supplementary Fig. 7b,c). Upon annealing in air at 200 °C or below, there is no significant change in the intensity at E F (Fig. 3a, blue). However, the R-to-P phase transition starts at 250 °C. These results illustrate that the V O concentration can be used to tune the conductivity of complex oxides in a way analogous to the role of dopants in conventional semiconductors with two important distinctions. First, in the case of V O in oxides, the doping process is reversible, revealing a way to dynamically control functionality not possible in conventional semiconductors18. Second, isolated anion vacancies are typically n-type dopants that result in higher electron conductivity. However, in the case of SCO, V O result in a lower conductivity because the electrons associated with V O populate a partially occupied t 2g sub-band of Cr, and thus decrease the number of charge carriers and increase optical transparency.