Homoepitaxial SrTiO3(110) film is grown by molecular beam epitaxy in ultra-high vacuum with oxygen diffusing from substrate as the only oxidant. The resulted oxygen vacancies (VOs) are found to be spatially confined within few subsurface layers only, forming a quasi-two-dimensional doped region with a tunable high concentration. Such a δ-function distribution of VOs is essentially determined by the thermodynamics associated with the surface reconstruction, and facilitated by the relatively high growth temperature. Our results demonstrate that it is feasible to tune VOs distribution at the atomic scale by controlling the lattice structure of oxide surfaces.

Transition metal oxide interfaces have exhibited a variety of novel phenomena, including the high-mobility two-dimensional electron gas,1–5 superconductivity6 and unusual magnetism.7 While these phenomena are inherently related to the artificially designed architecture, oxygen vacancies (VOs) are recognized to be particularly important in determining the exotic behaviors of the oxides systems.8,9 As electron donors in oxides, VOs were observed to be one of the major factors for the high carrier density at the LaAlO3/SrTiO3 interface.10 VOs induce ferromagnetism in the epitaxial LaCoO3 films via the ordering of excess electrons in Co 3d orbit.11 It was also suggested that the existence of VOs plays an important role in the formation of undesirable insulating phase in ultrathin films of many metallic oxide materials,12,13 referred to as the “dead layer” behavior. How VOs are involved in such complex phenomena is still unclear.

One essential issue toward the design of emergent properties of oxide films is how to characterize and control VOs at the atomic scale. Recently, quantitative measurements of VO concentration have been realized with high spatial resolution benefited by the development of the state-of-the-art aberration-corrected scanning transmission microscope (STEM) and related spectroscopy techniques.14–16 However, tuning the VO density precisely by growth control is still extremely challenging. Different oxidants, i.e., molecular O2, O3 and atomic O, have been used with different feeding pressure during the growth.10,17–19 This approach is limited by thermodynamics of the material since the pressure of oxidant has to coordinate with temperature and other parameters in order to ensure the formation of the desired phase.10,19 In many cases the optimal condition for oxygen stoichiometry cannot be reached, and residual VOs are inevitably formed. It has been reported that VOs even form inhomogeneous clusters in SrTiO3 (STO) films.20 Post annealing is another thermodynamic approach and seems to be effective to reduce the VO density. However, it does not allow the precise control of VOs distribution at the atomic scale either.

In the current work, to study the control of VOs at the atomic scale, we grow SrTiO3 homoepitaxial films along the polar [110] direction. Atomically well-defined (4 × 1) reconstruction is maintained on the surface all through the growth [see Fig. 1 (a) and (b)], which compensates the polarity.21 More importantly, the energy configuration associated with the surface reconstruction results in the spatial confinement of VOs within only few subsurface layers of the film, while their concentration remains extremely low on the topmost surface and in the bulk of STO substrate. We show that such a quasi-two-dimensional (2D) VO-doped layer, mimicking the δ-doping case in conventional semiconductors, can be precisely controlled.

FIG. 1.

The structural model of SrTiO3(110)-(4 × 1) surface and the illustration of the homoepitaxial MBE growth in UHV. (a) The top- (upper panel) and cross-sectional (lower panel) views of the (4 × 1) reconstruction. The (4 × 1) unit cell is indicated by the rectangle. (b) STM image (10 nm × 10 nm, +2 V/100 pA) of the surface of 1 ML film that keeps the (4 × 1) reconstruction. (c) The right panel shows the RHEED patterns of the substrate and the 15 ML film surfaces. IR of (01) diffraction spot along [1 1 ¯ 0] is monitored (left panel) as the feedback signal to control the shutter status of Sr source while Ti is deposited continuously. The blue dashed line indicates the trigger timing of shutter “open” and the red dashed line indicates the timing of “close”. Keeping IR evolving within a reversible window,26 which is accompanied by the variation of relative surface Sr concentration (ΔSr), we repeat the “open-close” cycles to obtain the film cation stoichiometry precisely.22 Only a part of the growth is shown.

FIG. 1.

The structural model of SrTiO3(110)-(4 × 1) surface and the illustration of the homoepitaxial MBE growth in UHV. (a) The top- (upper panel) and cross-sectional (lower panel) views of the (4 × 1) reconstruction. The (4 × 1) unit cell is indicated by the rectangle. (b) STM image (10 nm × 10 nm, +2 V/100 pA) of the surface of 1 ML film that keeps the (4 × 1) reconstruction. (c) The right panel shows the RHEED patterns of the substrate and the 15 ML film surfaces. IR of (01) diffraction spot along [1 1 ¯ 0] is monitored (left panel) as the feedback signal to control the shutter status of Sr source while Ti is deposited continuously. The blue dashed line indicates the trigger timing of shutter “open” and the red dashed line indicates the timing of “close”. Keeping IR evolving within a reversible window,26 which is accompanied by the variation of relative surface Sr concentration (ΔSr), we repeat the “open-close” cycles to obtain the film cation stoichiometry precisely.22 Only a part of the growth is shown.

Close modal

The films were grown on Nb-doped (0.7 wt. %) SrTiO3(110) substrates in a molecular beam epitaxy (MBE) system, following the recipe reported previously22 as described in the supplementary material. The high angle annular dark field (HAADF) and annular bright field (ABF) micrograph STEM images as well as the electron energy loss spectroscopy (EELS) were obtained by the aberration-corrected JEOL JEM-ARM200CF equipped with a cold field-emission electron source. The STEM sample was prepared by mechanical polishing and ion milling at liquid nitrogen temperature. The density functional theory (DFT) calculations were performed with the projector augmented-wave method23 and Perdew-Burke-Ernzerhof functional24 with the slab model and implemented in the “Vienna ab initio simulation package” (VASP) code.25 

To grow the SrTiO3(110) film in UHV, the recently developed shuttered MBE technique22 is used, in which the growth can be precisely controlled by switching the shutter status of the metal (Sr or Ti) evaporation source with the reflection high-energy electron diffraction (RHEED) intensity (IR) as the feedback signal [illustrated in Fig. 1 (c)].26 The film surface maintains the (4 × 1) reconstruction as on the substrate all through the growth. Further characterized by cross-sectional STEM, the SrTiO3 film keeps the perovskite lattice structure without any other phase formed. The interface between film and substrate is not distinguishable, indicating the high quality of the homoepitaxial film (Fig. 2). And the uniform image intensity of Sr/Ti manifests the cation stoichiometry.22 

FIG. 2.

(a) The atomically-resolved cross-sectional HAADF and (b) ABF image of the 15 ML film. The dashed lines indicate the film surface and the interface between the film and substrate. A high concentration of VO is located at the “interface” between surface TiO4 tetrahedra and bulk TiO6 octahedra, i.e. the 2nd layer highlighted in red.

FIG. 2.

(a) The atomically-resolved cross-sectional HAADF and (b) ABF image of the 15 ML film. The dashed lines indicate the film surface and the interface between the film and substrate. A high concentration of VO is located at the “interface” between surface TiO4 tetrahedra and bulk TiO6 octahedra, i.e. the 2nd layer highlighted in red.

Close modal

Shimoyama et al. grew homoepitaxial SrTiO3(001) films by MBE in UHV without introducing any gaseous oxidant.27–29 The continuous oxygen feeding from the oxide substrate to the surface was attained by the fast diffusion of VOs deep into the bulk substrate.28,29 Similarly, in the current work for SrTiO3(110) where the residual oxygen in the growth chamber ( P O 2 < 2 × 1 0 10 mbar) can only supply < 1 % required by the epitaxy with the growth rate of 0.3 ML/min, the substrate has to act as an oxygen source. The key is to keep the growing surface (4 × 1) reconstructed, whose atomistic structure has been resolved,30,31 as shown in Fig. 1 (a) and (b). The surface energy of SrTiO3(110) is effectively reduced by the formation of TiO4 tetrahedra. And the formation of VOs on the reconstructed surface layer will cost energy ( 1.2 e V for each VO relative to the one in bulk (supplementary material)). Therefore, oxygen in the substrate would be driven to diffuse toward the surface to make Ti four-coordinated, as demonstrated by the extremely low VOs density on (4 × 1) surface [see Ref. 31 and Fig. 1 (b)]. The chemical potential of such an oxygen source in solid state is equivalent to that of a gaseous oxidant with the area density of 2 × 1 0 15 /cm2, unachievable in any normal film growth conditions. Therefore the VO density is effectively reduced in the film.

It is essential to clarify whether or not the oxygen diffusion from substrate (equivalently, the VO diffusion from surface to subsurface layers) is kinetically feasible in this epitaxy process without gaseous oxygen environment. We explore the possible pathways of VO diffusing from surface to subsurface layers by evaluating the energy barriers with Nudged Elastic Band (NEB) calculations.32 It is found that the two-step diffusion pathway shown in Fig. 3(a) with the barrier of 0.75 eV/0.56 eV is the most favorable one (supplementary material). To diffuse further into the bulk, the barrier that VOs need to overcome is only about 0.6 eV.33,34 In the thermal activation scenario, the diffusion rate of VO, denoted as h, can be estimated by the Arrhenius equation h = ν e x p ( E b k B T ) , where Eb is the energy barrier, ν is the attempt frequency (empirically given by 1012 s−1), and kB is the Boltzmann constant. Given the film growth rate of ∼ 0.3 ML/min as in the current experiments, 3 oxygen atoms are required to diffuse to surface within 200 seconds. As long as the growth temperature (T) is higher than 90 °C, sufficient oxygen can be supplied (supplementary material). Such an oxygen diffusion is undoubtedly kinetically feasible at the current growth temperature (T = 800 °C).

FIG. 3.

Thermodynamics and kinetics of the surface atomistic processes during film growth. (a) The VO diffusion from surface to bulk, (b) cation diffusion on surface and (c) oxygen desorption from the surface.

FIG. 3.

Thermodynamics and kinetics of the surface atomistic processes during film growth. (a) The VO diffusion from surface to bulk, (b) cation diffusion on surface and (c) oxygen desorption from the surface.

Close modal

In MBE growth with the RHEED intensity as feedback control signal, the diffusion of deposited metals needs to be activated to form the long-range order. Considering the cation diffusion barrier in the bulk phase (more than ∼ 3.0 eV35), the required growth temperature is estimated to be ∼800 °C with the current growth rate by the Arrhenius equation. Experimentally the elevated temperature of 800 °C is critical to maintain the surface crystallinity and reconstruction (supplementary material), as shown schematically in Fig. 3 (b).

Oxygen desorption from the surface is another factor that has to be considered during the growth in UHV at such a high temperature. As schematically shown in Fig. 3(c), the desorption energy barrier [variable with the chemical potential of oxygen μ(T, P O 2 )]36 is estimated to be about 3.0 eV at such a reconstructed surface at T = 800 °C and P O 2 = 1 × 1 0 10 mbar (supplementary material). Usually increasing the pressure of gaseous oxidant for growth can increase the desorption barrier and suppress the oxygen desorption. But here, the oxygen desorption can be compensated effectively by the bulk oxygen diffusing to the surface due to the low diffusion barrier. Consequently, VOs are distributed deep in the bulk while their density at the topmost surface remains extremely low.31 Further elevating the grow temperature may increase the desorption rate of oxygen, i.e., the total amount of VOs all through the sample would be increased.

We have further studied the VO distribution in the epitaxial films with the EELS of STEM. In the simple ionic scenario, each VO normally donates two electrons that reduce two adjacent Ti4+ ions to Ti3+ in SrTiO3. Therefore, the VO density and distribution at the atomic scale can be quantitatively determined by analyzing the layer-dependent Ti valence shown in EELS. Figure 4 (a) presents the layer-resolved EELS spectra of Ti L2,3 edge from the surface all the way across the interface to substrate. Beneath the 6th layer from the topmost surface, the Ti core level spectra of different layers are uniformly identical to Ti4+ (supplementary material).37 This implies that VOs are homogeneously distributed with the density as an undetectable background, and the STEM sample preparation has no significant effect on EELS. The spectra taken at the Ti columns within 6 layers from the topmost surface show clear chemical shift to the lower energy relative to those for the bulk, indicating the existence of Ti3+ associated to VOs. Quantitatively, we fit the spectra with two components of Ti3+ and Ti4+ by the multiple linear least-squares (MLLS) method to determine the VO density (supplementary material). Figure 4 (b) shows the results of the 15 ML film. Note that the VOs concentration at topmost layer (extremely low) was determined by the in situ STM and DFT analyses,31 rather than from EELS results to avoid possible artifact introduced by the ex situ sample transfer. The VOs show the highest density (∼ 14%) on the 2nd layer and then quickly diminishes to the background reference level (undetectable by EELS) beneath the 6th layer. The layer dependence of VO distribution of the sample with 10 ML homoepitaxial film is also resolved and plotted in Fig. 4(b) for comparison. In both cases, the most important finding is that VOs are mainly confined between the 2nd to 6th layer from the surface with the density maximum at the 2nd layer. The only difference is quantitative-lower VO density in the thinner film at each respective layers (2nd-5th). This is due to the shorter growth time for the thinner film, which means less oxygen desorption from the surface.

FIG. 4.

The layer-dependent distribution of VO near the surface. (a) The atomically-resolved EELS of Ti L2,3 edge of the 15 ML as-grown film acquired at each Ti site from film surface to the bulk substrate. close The spectra to surface with a fraction of Ti3+ is plotted with different colors. (b) The layer-resolved VO concentration of 15 ML, 10 ML as-grown and 15 ML post-annealed samples (15 ML-PA). The values are determined from the Ti3+ to Ti4+ ratio in EELS. The VO distribution calculated from DFT for 15 ML film (assuming the same total amount of VO as experimentally determined) is plotted for comparison. The layer index in (b) is aligned to the EELS spectra in (a) and the background reference level of VOs in the substrate is valued to zero.

FIG. 4.

The layer-dependent distribution of VO near the surface. (a) The atomically-resolved EELS of Ti L2,3 edge of the 15 ML as-grown film acquired at each Ti site from film surface to the bulk substrate. close The spectra to surface with a fraction of Ti3+ is plotted with different colors. (b) The layer-resolved VO concentration of 15 ML, 10 ML as-grown and 15 ML post-annealed samples (15 ML-PA). The values are determined from the Ti3+ to Ti4+ ratio in EELS. The VO distribution calculated from DFT for 15 ML film (assuming the same total amount of VO as experimentally determined) is plotted for comparison. The layer index in (b) is aligned to the EELS spectra in (a) and the background reference level of VOs in the substrate is valued to zero.

Close modal

Due to the high diffusion rate of VOs at the high growth temperature, the film surface almost reaches the thermodynamic equilibrium immediately after Sr/Ti deposition. The spatial confinement of VOs can be understood by resolving the energetics associated with their distribution that is determined by the (4 × 1) reconstruction. On the 2nd layer from the topmost surface, each TiO6 octahedron shares two oxygen atoms with the surface TiO4 tetrahedra [see Fig. 1 (a)]. Such a bonding configuration is quite different from that in the bulk. Given a VO in the sample, our DFT calculations verify a significant energy gain of ∼ 0.6 eV when it is in the 2nd layer over any other possible locations (supplementary material). Note a VO in the topmost surface is the most unfavorable and its energy loss is even larger relative to the 2nd layer (∼ 1.8 eV). We systematically calculate the formation energy of a VO on the i-th layer ( 𝜖 i ) and estimate the according VO concentration determined by the thermodynamics as e 𝜖 i k B T . The calculated results in Fig. 4(b) show that almost all VO would have been confined on the 2nd layer. Experimentally, the observed VO distribution is broadened due to the repulsive Coulomb interaction between ionized VOs which is not considered in the calculations. In addition, electrons ionized from VO may diffuse within a certain range, resulting in further broadening of the reduced Ti3+ ions distribution.

Post annealing slightly affects the VO distribution. At 300 °C, annealing in vacuum for 14 hours broadens the VO distribution to the 7th layer [Fig. 4(b)], while the maximum still locates at the 2nd layer. The broadening might be due to the repulsive interaction between VOs and the annealing enhances the diffusion of VOs. More importantly, the density of electrons (ionized from the confined VOs) in the 2D system beneath the surface can be modified. For the as-grown 15 ML film, the VO density on the 2nd layer is 14%, corresponding to the electron density of 4 × 1014 /cm2. Annealing at an intermediate temperature [300 °C, see Fig. 4 (b)] allows the VO diffusing towards the surface region but suppresses the oxygen desorption. As the result, VO density on the 2nd layer decreases to 7% and the electron density decreases to 2 × 1014/cm2. And the total amount of VOs through the 7 layers from the surface is significantly reduced by ∼ 17%. Beneath the 7th layer, VOs are uniformly distributed VOs with a low density beyond EELS sensitivity.

During the growth, the oxygen chemical potential provided by the solid-state substrate is related to the deposition of metal atoms on the surface. Although the oxygen desorption is difficult to control at high temperature, it is relatively easy to control the growth of cations. This enables us to achieve the robust spatial confinement of VOs – we obtain the same VO distribution in STO film with different thickness and actually even in the single crystalline substrate without epitaxy, as long as the (4 × 1) reconstruction is formed on the surface. To conform this, We have used a undoped SrTiO3(110) with the the (4 × 1) reconstructed surface and perform ex situ electric transport measurements (supplementary material). We found that the bulk is insulating while the subsurface layers is metallic due to the existence of confined VOs. The resistance of the VO layer is much lower than that of the bulk by about 3 orders of magnitude in the measured temperature range. Combining with the evidence that the topmost surface is insulating provided by STM spectroscopy and DFT calculations, we concluded that a 2D metallic layer is flanked (and protected against atmosphere exposure) by the insulating bulk and reconstructed surface, mimicking the δ-doping in conventional semiconductors. Such a buried metallic phase beneath the reconstructed SrTiO3(110) surface has also observed by photoemission spectroscopy.4 The above discussions can also be applied to other (n × 1) surfaces since they have the same TiO4 building blocks.30 In general, using the reconstructed surface as a confinement boundary, one can tune VO density in a range comparable or even wider than traditional methods such as the gate voltage in field effect transistors.

In conclusion, we have grown homoepitaxial SrTiO3(110) films in UHV and studied the controllability of VOs by precisely tuning the surface cation concentration. By preserving the (4 × 1) surface reconstruction, oxygen in the substrate bulk is driven to diffuse toward the surface to supply oxidant continuously. The film growth is associated with the formation of VOs, which are confined within a few subsurface layers between the fully oxidized topmost surface and the bulk substrate with a low VOs concentration as undetectable background, mimicking the δ-doping case in conventional semiconductors. The thermodynamics associated with the surface reconstruction determines such VOs distribution. It is demonstrated that the density and distribution of VOs in oxides can be controlled at the atomic scale by tailoring the lattice structure of the surfaces, thus providing an opportunity to construct and manipulate novel quantum phenomena in low-dimensional oxide systems.

See supplementary material for details about the sample preparation, multiple atomistic process during film growth, the determination of VO concentration from EELS, as well as the transport measurement of the VO layer flanked by insulating bulk and the reconstructed surface.

This work is supported by NSFC (11634016 & 11474334), MOSTC (2016YFA0202300) and the Strategic Priority Research Program B of CAS (XDB07010100). EWP and JZ are partially supported by US Department of Energy (DOE) under Grant No. DOE DE-SC0002136.

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