Introduction

Transition-metal oxide epitaxial thin films, which often exhibit behaviors different from those of the bulk oxide, have attracted a great deal of attention as a fascinating platform for exploring novel functionalities1,2,3,4,5. This is in part because strong correlations between charges, spins and lattices determine the functional properties of the films and these correlations are affected by structural distortions from substrate-induced strain effects6,7,8,9. For thin films of perovskite oxides ABO3 consisting of corner-shared oxygen octahedra BO6, such effects result in octahedral distortions including deformations, tilts and rotations, which are essentially characterized by displacements of the oxygen atoms in the film's lattice10. While substrate-induced distortions have been closely correlated with the functional properties of fully oxygenated perovskite oxide thin films, little is known about their effects on thin films of oxygen-deficient perovskites. The oxygen deficiency introduced in the perovskite structure produces various oxygen coordination environments for the transition metals, like tetrahedra and pyramids as well as octahedra11,12,13. Therefore, substrate-induced modifications of such various oxygen coordination units would provide additional routes to controlling the strong correlations and consequently to modifying or even enhancing the functional properties.

We here focus on oxygen-deficient Fe-based perovskite oxides, SrFeOx (SFOx), which exhibit a variety of structural and physical properties depending on their oxygen vacancy concentration11,14,15,16,17,18,19,20,21,22,23. The general formula of the known phases is described as SrFeO3–1/n with n = ∞, 8, 4, 2 and 1. The n = ∞ member, SrFeO3, has the simple cubic perovskite structure with corner-shared FeO6 octahedra and exhibits metallic conduction down to low temperatures14,15,24,25,26,27,28,29. The oxygen vacancies introduced in the SrFeO3 lattice result in the formation of various Fe environments which consist of ordered arrangements of corner-shared oxygen polyhedra including the FeO4 tetrahedra and FeO5 pyramids. An important consequence of these ordered arrangements of the oxygen vacancy is a charge-ordering of Fe with different valence states, which impacts on physical properties. In the n = 8 member SrFeO2.875, with nominal Fe3.75+, the oxygen vacancies order at 523 K and the ordering stabilizes the charge-ordered tetragonal phase with Fe4+O6 octahedra, distorted Fe3.5+O6 octahedra and Fe4+O5 pyramids11,16. The tetragonal phase undergoes another phase transition at ~70 K, which is due to charge disproportionation of Fe3.5+ into Fe4+ and Fe3+ and which is accompanied by an abrupt increase in the electrical resistivity14,27,30,31. Further increase in the oxygen vacancy concentration leads to SrFeO2.75 (n = 4) with nominal Fe3.5+11,16,32 and SrFeO2.5 (n = 2) with nominal Fe3+16,17,33,34,35,36,37,38. Although the oxygen vacancies in SrFeO2.75 are ordered at 598 K and those in SrFeO2.5 are ordered at 1103 K, charge disproportionation does not occur in these phases.

In this study we discovered in an oxygen-deficient SrFeOx (x ~ 2.8) epitaxial thin film a transport behavior markedly different from the corresponding behavior in the bulk material. The present thin film under substrate-induced tensile strain shows a metal-insulator transition, associated with the charge disproportionation of Fe3.66+ into Fe4+ and Fe3+ at 620 K. This transition temperature is much higher than ~70 K reported for the transition in the bulk material and is also much higher than room temperature. We also found the transition to be accompanied by oxygen-vacancy ordering. This significant increase in the metal-insulator transition temperature demonstrates that epitaxial growth of oxygen-deficient oxides under substrate-induced strain is a promising route for exploring novel functionality.

Results

Preparation of SrFeO2.8 epitaxial thin film

A brownmillerite-structure SrFeO2.5 (SFO2.5) epitaxial thin film was first prepared on a (001) SrTiO3 (STO) single-crystal substrate by pulsed laser deposition. An x-ray 2θ-θ profile of that film is shown in Fig. 1a, where we see (101) and (202) reflections of SFO2.5 at slightly lower angles than (001) and (002) reflections of the STO substrate. In the x-ray reciprocal space mapping (RSM) shown in Fig. 1c, we see (206), (484) and (602) reflections whose in-plane positions are the same as that of the STO (204) reflection. These x-ray diffraction results confirm that the deposited SFO2.5 brownmillerite thin film consists of multiple domains with (101) orientation and that the in-plane lattice of the each domain is fixed by the substrate lattice37.

Figure 1
figure 1

(a, b) X-ray 2θ-θ profiles for (a) as-deposited brownmillerite SrFeO2.5 (SFO2.5) film and (b) oxidized SFO2.5 (SFO2.5+δ) film grown on the STO substrates. (c, d) X-ray reciprocal space mappings around STO (204) reflections for (c) SFO2.5 and (d) SFO2.5+δ thin films. The intensity is plotted on a semi-logarithmic scale. All the data were obtained at room temperature. (e) 57Fe conversion electron Mössbauer spectrum of the SFO2.5+δ thin film at room temperature. The fitting results, which are listed in Table I, are shown with lines in black (total), red (Fe4+ singlet component), green (Fe3.5+ doublet component) and blue (Fe3+ doublet component). Note that the electric field gradients for the Fe3.5+ and Fe3+ doublets are oriented along the in-plane direction. The oxygen content of the SFO2.5+δ thin film was estimated to be 2.83.

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The deposited SFO2.5 thin film was then oxidized into SFO2.5+δ by the air-annealing at 773 K. Figures 1b and 1d show the x-ray 2θ-θ diffraction profile and RSM of the annealed film. One sees that the diffraction peaks shifted to higher 2θ angle positions and the multiple domain structure disappeared, indicating that the out-of-plane lattice spacing of the oxidized film is smaller than that of the as-deposited film. The decreased lattice size seen after the air annealing strongly suggests oxygen incorporation into the brownmillerite structure. Note that the in-plane lattice of the SFO2.5+δ film is still fixed by the substrate lattice.

The valence state of Fe ions in the oxidized thin film was investigated by Mössbauer spectroscopy. Figure 1e shows a room-temperature 57Fe conversion electron Mössbauer spectrum of the oxidized thin film and the result of the peak fitting. (The fitting results are also summarized in Table I.) No signal due to magnetic ordering was observed. The observed spectrum consists of a singlet with an isomer shift (IS) of 0.04 mm/s (the red line in the figure) and two quadrupole doublets with an IS of 0.12 mm/s (the green) and an IS of 0.35 mm/s (the blue). The observed IS values indicate that the singlet originates from Fe4+ and that the doublets with the smaller and larger IS values arise from Fe3.5+ and Fe3+, respectively11,39,40. The relative abundances of the Fe4+, Fe3.5+ and Fe3+ components are respectively 56%, 20% and 24%, giving an average oxidation state for Fe of +3.66 and giving an oxygen content of 2.83. Therefore the thin film obtained by air-annealing the SFO2.5 thin film is identified to be the oxygen-deficient perovskite SrFeO2.8 (SFO2.8).

Table 1 The isomer shift (IS) and the quadrupole splitting (QS) parameters obtained by fitting the 57Fe internal conversion electron Mössbauer spectra of the oxidized SFO2.5 (SFO2.8) thin film at 300 K. The directions of the electric field gradients of the Fe3.5+ and Fe3+ doublet components are set to be parallel to the in-plane direction of the film
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Metal-insulator transition associated with charge disproportionation above room temperature in SrFeO2.8 thin film

In contrast to the bulk SFO2.8, the electrical resistivity of the prepared film is as high as 1 Ωcm at room temperature. The quite high resistivity of the film decreased with increasing temperature and reached ~3 mΩcm at temperatures above 620 K (Fig. 2). We can clearly see that the metal-insulator transition occurs at 620 K and this transition suggests a significant change in the valence states of Fe in the film at that temperature. To clarify the nature of transition, we measured the time spectra of the nuclear resonant (“57Fe Mössbauer”) scattering41,42. The measured spectra of the SFO2.8 film at 300, 573 and 673 K in air are shown in Fig. 3 together with the theoretically simulated spectrum patterns. It is clear that the oscillation patterns in the time spectra are different below and above the transition temperature. The spectrum at 300 K shows time-dependent oscillation and the oscillation pattern is well reproduced with three components with distinct hyperfine parameters as listed in Table II. The result is in good agreement with that independently obtained in the 57Fe conversion electron Mössbauer measurement shown in Fig. 1e. The time spectrum at 573 K is essentially the same as that at 300 K. The parameters obtained from its analysis (also listed in Table II) confirm that the valence states of Fe in the film and the relative abundance do not change with changes in temperature below the metal-insulator transition temperature and that the thin film below the metal-insulator transition temperature contains Fe4+ (55%), Fe3.5+ (15%) and Fe3+ (30%). Importantly, the relative abundance of each Fe components in the film below the transition temperature is essentially the same as that in the charge-disproportionated insulating phase of bulk SFO2.827. The results indicate that below the transition temperature the most of Fe in the film have the integer valence state of either Fe4+ or Fe3+ and thus the charge-disproportionated state is formed. This is also consistent with the observed high resistivity of the SFO2.8 thin film.

Table 2 The isomer shift (IS) and the quadrupole splitting (QS) parameters obtained by fitting the time spectra of the nuclear resonant scattering of the SFO2.8 thin film at 300, 573 and 673 K. The directions of the electric field gradients of the Fe3.5+ and Fe3+ doublet components are parallel to the in-plane direction of the film
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Figure 2
figure 2

Temperature dependence of the electrical resistivity of SFO2.8 film from 720 to 300 K in air (red line) and SrFeOx (x ≈ 2.81) bulk from 300 to 5 K (black line).

The arrows denote the metal-insulator transition temperatures TMI. The resistivity data of the SFO2.81 bulk were taken from a report by P. Adler et al. (Ref. 27).

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Figure 3
figure 3

Time spectra of the nuclear resonant spectra of SFO2.8 thin film at 300, 573 and 673 K in air.

The inset in each panel shows the time spectrum of reflectivity taken with the reference K2MgFe(CN)6 sample. The ISs of the Fe ion components in the film were determined by fitting the inset spectra. Black dots and red lines correspond to the experimental data and fitting results, respectively. The spectra at 300 and 573 K can be reproduced with three components of Fe4+, Fe3.5+ and Fe3+. The spectrum at 673 K can be fitted with a single component of Fe3.66+. The parameters obtained from the fittings are listed in Table II.

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The spectrum at 673 K (above the transition temperature), on the other hand, shows no oscillation pattern and can well be reproduced by a single component with an IS of –0.04 mm/s. This IS value is quite small, suggesting that the Fe ion in the SFO2.8 thin film at this temperature has an unusually high oxidation state. Given that the oxygen content of the film does not change up to this temperature16, the result indicates that the film above the metal-insulator transition temperature has a single Fe site with a mixed-valence state of Fe3.66+. This single Fe site also suggests that the oxygen vacancies are disordered above the transition temperature43. Our analysis of the time spectra of the nuclear resonant scattering therefore leads to a conclusion that the metal-insulator transition seen in the SFO2.8 thin film at 620 K is due to the charge disproportionation of Fe3.66+ above the transition temperature into Fe4+ and Fe3+ below the temperature.

Structural transition due to oxygen-vacancy ordering

We also investigated whether a structural change at the metal-insulator transition temperature occurs in our SFO2.8 thin film. Figure 4 shows temperature dependence of the out-of-plane lattice spacing of the film together with that of the cubic STO substrate, both of which were obtained in the x-ray 2θ-θ measurements at temperatures between 720 and 300 K in air. We found that with decreasing temperature the out-of-plane lattice spacing of the film decreased abruptly at 620 K, where the metal-insulator transition occurred. It is clear that the structural change in the film was not caused by the substrate lattice, which exhibits the normal thermal change. We also confirmed from the RSM measurements that the in-plane lattice of the film (the open red square in Fig. 4) is fixed by the substrate lattice both above and below the transition temperature. The quite large change in the out-of-plane lattice spacing seen at 620 K indicates that the structural transition in our oxygen-deficient SFO2.8 thin film is due to a change in oxygen-vacancy ordering. Thus the charge disproportionation transition in Fe in the SFO2.8 thin film is accompanied by both a metal-insulator transition and an oxygen-ordering structural transition.

Figure 4
figure 4

Temperature dependence of the out-of-plane (circles) and in-plane (squares) lattice spacings for the SFO2.8 film (red) and the STO substrate (black) from 720 K to room temperature in air.

The in-plane lattice of the film is fixed by the substrate in the entire range of temperatures. The dashed line denotes the metal-insulator transition temperature TMI ( = 620 K) of the SFO2.8 film.

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Discussion

The experimental results described above indicate that the present SFO2.8 thin film undergoes not only a structural transition due to oxygen-vacancy ordering at 620 K but also the charge disproportionation of Fe3.66+ into Fe4+ and Fe3+ at that temperature. Similar oxygen-vacancy ordering and charge disproportionation transitions were reported in bulk SrFeOx (x ≈ 2.8) samples, but in those samples the oxygen-vacancy ordering occurred at 523–598 K and the charge disproportionation took place around 70 K16,27. Interestingly, while the oxygen-vacancy ordering in bulk at 523–598 K is not accompanied by changes in transport properties, the charge disproportionation at 70 K induces a metal-insulator-like resistivity jump. Thus the metal-insulator (and also charge disproportionation) transition temperature of 620 K of the film is significantly higher than the transition temperature reported in the bulk material27,30. We note that the oxygen ordering temperature also increases. The behaviors different from those of the bulk samples are likely to be related to the substrate-induced strain effect in our epitaxial thin film.

Assuming that the low-temperature charge-disproportionated phase (< 620 K) of the present thin film is similar to that of the bulk SrFeOx (x ≈ 2.8), the oxygen-deficient octahedra Fe3.66+O5.66 in the high-temperature phase changes into Fe4+O5 pyramids, Fe4+O6 octahedra, Fe3.5+O6 octahedra and Fe3+O6 octahedra at the transition temperature of 620 K. Our experimental results indicate that because the SFO2.8 film is subjected to significant tensile strain by the substrate lattice, the stretched film's lattice stabilizes long Fe-O bonds like the Fe3+-O bond. In fact, the average Fe3+-O bond distance observed in the charge-disproportionated bulk sample was much shorter than expected for a typical ionic bond14. Thus, the lattice stretched by the tensile strain is preferable for stabilizing the low-temperature Fe3+O6 octahedra and consequently the charge-disproportionated phase in the film is more stable than the phase in the bulk, leading to the increase in transition temperature.

It is also interesting to point out that the transition temperatures of the charge disproportionation for fully-oxygenated perovskite oxides such as CaFeO3 and La0.33Sr0.67FeO3, unlike the oxygen-deficient perovskite SrFeO2.8, are little influenced by substrate-induced strain44,45,46. No significant changes in the charge disproportionation transition temperature were observed when the strained thin films were compared with the bulk materials. This implies that the fully oxygenated octahedral network is pretty rigid, whereas the oxygen-deficient compounds, which contain various oxygen coordination environments, have structural flexibility that lets them accommodate the strain-induced deformation of the polyhedra. Given the diverse oxygen coordination environments seen in oxygen-deficient transition-metal oxides like cobaltites, strain-induced modification of oxygen coordination environments can therefore be useful to improve the functional properties and even to explore novel functionality of oxygen-deficient oxide thin films.

Methods

Preparation of SrFeO2.8 thin films

SrFeO2.8 (SFO2.8) thin films were prepared by oxidizing brownmillerite-structured SrFeO2.5 (SFO2.5) thin films. The SFO2.5 films with thicknesses of 20–100 nm were epitaxially grown on (001) SrTiO3 (STO) single-crystal substrates by pulsed laser deposition. During the thin film deposition, the stoichiometric target was ablated at 2 Hz with a KrF excimer laser (λ = 248 nm, COHERENT COMPex-Pro 205 F) with a laser spot density of 1 J/cm2 on the target surface. The oxygen partial pressure and the substrate temperature were kept at 1 × 10−5 Torr and 973 K. The deposited SFO2.5 films were oxidized to SFO2.8 by annealing them in air at 773 K. The oxygen content in the films was evaluated from the relative abundances of Fe4+, Fe3.5+ and Fe3+ obtained from the fitting of 57Fe internal conversion electron Mössbauer spectrum measured at room temperature.

Structural characterizations

The crystal structures and growth orientations of the as-deposited SFO2.5 thin films and the SFO2.8 films were characterized with a conventional four-circle X-ray diffractometer (PANalytical X'Pert MRD) equipped with a high-temperature sample stage (DHS1100) and operated with Cu Kα1 radiation. The indices of the diffraction peaks of SFO2.5 are based on a bulk orthorhombic cell (a = 5.67 Å, b = 15.59 Å and c = 5.53 Å) and those of SFO2.8 are based on a pseudo-cubic perovskite cell. The out-of-plane lattice spacing of the SFO2.8 film was determined from the (002) Bragg reflection position in a 2θ-θ diffraction pattern and the in-plane lattice spacing was estimated from the (204) reflection position in the reciprocal space mapping. All diffraction measurements were carried out in air.

Electrical characterizations

Electrical resistivity of the SFO2.8 thin films was measured, in a voltage-source mode, by a two-terminal method. Pt metal electrodes were sputtered at room temperature. The temperature dependence of resistivity was obtained while cooling the samples from 773 K to room temperature in air.

57Fe internal conversion electron Mössbauer measurements

The internal conversion electron Mössbauer spectroscopy measurements with the 57Fe excitation energy, 14.4 keV, were performed to see the valence states of Fe in the SFO2.8 thin films. The source velocity was calibrated by using a pure α-Fe film. A 10 mm × 10 mm SFO2.8 thin film with thickness of 100 nm was used for the measurement. The Mössbauer spectrum was fitted with Lorentzian functions by using the “Fit;o)” program.

Measurements of time spectra of the nuclear resonant scattering

The measurements were conducted in a total reflection geometry (a grazing incidence angle of ~0.3°) at the beam line BL09XU of SPring-8. Incident photons were tuned to the 57Fe Mössbauer resonance at 14.4 keV by using a monochromator with 2.2 meV resolution. An eight-cell avalanche photodiode detector made from a thin silicon wafer with a depletion region of around 60 mm was used. The temperature of the thin-film samples was controlled by a resistive heater and the spectra were measured at 300, 573 and 673 K in air.

To determine isomer shift (IS) for the Fe ion components in the film from the time spectra, two types of spectra were collected at each measuring temperature. One was measured with only the film. The other was measured with the film and the reference. For the measurement of the latter spectra, a plate of K2MgFe(CN)6 (RITVERC GmbH) was used as a reference sample and was placed in front of the films during the measurements. The forward scattering time spectrum of the reference sample was also measured separately and its scattering amplitude spectrum was determined by fitting this time spectrum. The obtained nuclear resonance characteristics of the reference were in compliance with initially given parameters: IS = –0.101 mm/s relative to the IS of α-Fe, the line width is 0.165 mm/s and the 57Fe density 0.25 mg/cm2.

The theoretical simulations were performed by using the “REFTIM” program package47, which was modified in order to take into account the existence of the reference sample in the beam in some experiments and was adjusted for calculations of the forward scattering time spectrum as well (for the fit of the spectra for the reference sample). At the beginning the prompt reflectivity curve from the SFO2.8 film on the STO substrate was fitted and the obtained parameters of the electronic density depth distribution were used. These parameters were fixed in the model intended for the fit of the time spectra of the nuclear resonant reflectivity. The exact thickness of the SFO2.8 film was 107.4 nm and the mean square roughness of the surface was 0.8 nm.