Directional ionic transport across the oxide interface enables low-temperature epitaxy of rutile TiO₂

Abstract

Heterogeneous interfaces exhibit the unique phenomena by the redistribution of charged species to equilibrate the chemical potentials. Despite recent studies on the electronic charge accumulation across chemically inert interfaces, the systematic research to investigate massive reconfiguration of charged ions has been limited in heterostructures with chemically reacting interfaces so far. Here, we demonstrate that a chemical potential mismatch controls oxygen ionic transport across TiO₂/VO₂ interfaces, and that this directional transport unprecedentedly stabilizes high-quality rutile TiO₂ epitaxial films at the lowest temperature (≤ 150 °C) ever reported, at which rutile phase is difficult to be crystallized. Comprehensive characterizations reveal that this unconventional low-temperature epitaxy of rutile TiO₂ phase is achieved by lowering the activation barrier by increasing the “effective” oxygen pressure through a facile ionic pathway from VO_2-δ sacrificial templates. This discovery shows a robust control of defect-induced properties at oxide interfaces by the mismatch of thermodynamic driving force, and also suggests a strategy to overcome a kinetic barrier to phase stabilization at exceptionally low temperature.

Introduction

Interfaces formed by two dissimilar materials can break the translational symmetry and thereby provide an opportunity to develop functionality that is unachievable in bulk materials^1,2,3,4,5. When two materials (I, II) that have different work functions (i.e., $\mu _e^I \, < \, \mu _e^{II}$) are brought together at a semiconductor heterojunction, charge carriers near the interface diffuse across the junction (Fig. 1a);^1,2 as a consequence, a conducting channel with high carrier density and high electron mobilities (e.g., two-dimensional electron gas) could be created at the interfaces between normally-insulating materials^2,3. If an external bias is applied to adjust this built-in potential, the amount of transferred charge flow can be controlled by changing the electrochemical potential across chemically-inert interfaces^2,6, which is the basic principle of heterojunction field effect transistors (HFETs)⁷.

Fig. 1: Low-temperature epitaxy of rutile TiO₂ on VO₂ sacrificial templates.

a Schematics of possible directional charge (top) and ionic (bottom) transport due to chemical potential mismatch (Δμ) across the interface with loss of translational symmetry. Symmetrical x-ray scan of b TiO_X/TiO₂ homostructure and d TiO₂/VO₂/TiO₂ heterostructure containing TiO₂ films grown at 150 °C. Contrary to the absence of peak around TiO₂ substrate peak in homostructure (b), (002) Bragg reflections and Kiessig fringes around the peak from TiO₂ substrate in heterostructure in d indicate that the TiO₂ films are epitaxially grown at 150 °C on VO₂ templates with sharp interface. HRTEM images and FFT patterns of c TiO_X /TiO₂ homostructure (scale bar = 2 nm) and e. TiO₂/VO₂/TiO₂ heterostructure (scale bar = 5 nm) projected with [100] zone axis. Unlike the amorphous nature of TiO_X films in homostructures (green square in c), obvious diffraction spots were observed in the FFT pattern of the TiO₂ films on VO₂/TiO₂ (green square in e), and were same as those observed from the TiO₂ substrates (yellow square in e); this similarity indicates an identical epitaxial relationship of TiO₂ epitaxial films with TiO₂ substrates in TiO₂/VO₂/TiO₂ heterostructure.

As an analogy to the reversible control of electric charge transfer at an interface with discontinuity, (electro-)chemical potential mismatch for oxygen (Δμ_O) between two materials ($\mu _O^I \, < \, \mu _O^{II}$) may give rise to charged ionic transfer to bring the equilibrium of the system with heterogeneous junction at the interface (Fig. 1a)^{5,8,9,10,11,12}. In particular, charged ionic defects migrate by ionic diffusion (e.g., diffusion of oxygen ions through vacancies) to mitigate Δμ_O at the interface;^{9,10,11,12,13,14} charged ions, in principle, are transferred to adjacent materials and reconfigured by the redox reaction across the chemically-reacting interfaces in oxide heterostructure until the chemical potentials of the layers match ($\mu _O^I = \mu _O^{II}$ in Fig. 1a)⁵.

For example, the vanadium dioxide (VO₂), the archetypal correlated oxide with metal-insulator (MI) transition near room temperature, is interfaced with ionic liquid (IL), and then (electro-)chemical potential can be adjusted by applying an external electric field across the VO₂/IL interfaces^9,14,15,16. In this case, instead of electric charge transfer, charged oxygen ions out-diffused to the IL to equilibrate the (electro-)chemical potential between VO₂ and IL; the formation of oxygen vacancies (V_O) by oxygen ion migration are responsible for the reversible insulator-to-metal transition and giant lattice expansion in VO₂ films under the positive bias¹⁵. Furthermore, V_O concentrations that develop in LaNiO₃, LaTiO₃, and In₂O₃ can be modulated by the directional oxygen flow to the adjacent layers^10,11,12.

At interfaces where ionic flux ($J_{O^{2 - }}$ in Fig. 1a) is directional, the dynamics of charged ions may be important by assembling other metal oxides that have different μ_O^10,11,14. The redistribution of charged vacancies can screen the electric fields that Δμ_O causes. Therefore, the massive redistribution of charged ions can be accelerated by extremely increasing the thermodynamic Δμ_O across the interfaces; by supplying unidirectional charged ionic flux, this redistribution may offer a spontaneous route that can facilitate synthesis of crystalline materials¹², and may enable robust control of defect-induced properties at oxide interfaces^9,13,14.

Here, we demonstrate the formation of high-quality rutile TiO₂ epitaxial films at exceptionally low temperature, which is driven by directional transport of oxygen ions across the TiO₂/VO₂ heterointerfaces. Contrary to the amorphous nature of TiO₂ films directly grown on TiO₂ substrates at 150 °C, single-crystal rutile TiO₂ layer is synthesized by forming the heterointerface with VO₂ template at the lowest growth temperature T_G (<150 °C) ever reported, at which rutile TiO₂ is difficult to be crystallized. By experimental characterization using atomic-resolution electron microscopy and synchrotron x-ray spectroscopy combined with theoretical calculation, we demonstrate that a facile ionic diffusion of oxygen ions from the oxygen reservoir VO₂ along the [001] channel decreases Δμ_O, and thereby enables this unprecedented epitaxy of rutile TiO₂ at low temperature by lowering the activation barrier for formation of stable nuclei. Interestingly, this directional ionic transport improves the registry in the lattice of TiO₂ films at the expense of the structural and electronic modulation in an oxygen-deficient VO_2-δ sacrificial layer. As a result of the mismatch of thermodynamic driving force combined with kinetically-facile migration of oxygen ions across the TiO₂/VO₂ interfaces, the massive redistribution of oxygen ions enables low-temperature epitaxy of high-temperature-stabilized phase by increasing an “internal” oxygen supply across the chemically-reacting oxide interface.

Results

Low temperature epitaxial growth of rutile TiO₂ films on VO₂ template

Prior to TiO₂ growth, the substrates with 12-nm-thick VO₂ template were prepared on (001)-oriented TiO₂ substrates by pulsed laser deposition (Supplementary Fig. 1). X-ray diffraction (XRD) results (Supplementary Fig. 2a) showed sharp VO₂ (002)_R peaks (in rutile notation) at ~ 2θ = 65.9° (c = 0.2839 nm) without other peaks related to vanadium oxides that had valence states other than +4. A steep MI transition (ΔR_S ~ 10^3.3) occurred on the VO₂ films at T_MI ~ 298 K (Supplementary Fig. 2b); this result indicates the formation of coherently tensile-strained VO₂ films with high crystal quality and negligible V_O^9,15,17,18.

Then, the 6 nm-thick TiO₂ films were grown at low T_G ~ 150 °C with the same oxygen pressure ($pO_2$ ~ 12 mTorr) on two substrates: 1) (001) TiO₂ single crystals without the VO₂ template layer (denoted as TiO₂/TiO₂ hereafter) and 2) (001) TiO₂ single crystals with the VO₂ template (denoted as TiO₂/VO₂/TiO₂). Symmetric 2θ-ω scans using synchrotron X-ray scattering on TiO₂/TiO₂ grown at T_G = 150 °C (Fig. 1b) detected no Bragg reflections except for substrate (2θ = 49.54°); this absence indicates no formation of crystalline TiO₂ films¹⁶; the formation of amorphous TiO₂ films on TiO₂ substrates is attributed to the thermodynamic or kinetic instability of the rutile TiO₂ phase, which requires sufficient thermal energy for phase formation^{19,20,21,22,23}. To exclude the possible coincidence of diffraction peak from TiO₂ films and substrates in homostructures in Fig. 1b, TiO₂ films were also grown on (100) Al₂O₃ single crystal substrates (Supplementary Fig. 3). The symmetric 2θ-ω scan in wide range of angle also detected the only peak related to the (100) Al₂O₃ substrate (2θ = 68.22°) due to the formation of amorphous films at T_G = 150 °C.

In contrast, rutile TiO₂ epitaxial films were strikingly stabilized by introducing the VO₂ layers on TiO₂ substrate at the same condition with T_G = 150 °C. Symmetric 2θ-ω scans of TiO₂/VO₂/TiO₂ heterostructures (Fig. 1d) showed two Bragg reflections, one from the rutile TiO₂ (2θ = 49.46°) substrates, and one from VO₂ (2θ = 51.8°) films. More importantly, the TiO₂ substrate peak was resolved to a slightly broad peak from TiO₂ epitaxial films¹⁶, which did not appear in the scans of TiO₂/TiO₂ homostructure. Each film peak exhibited a Kiessig fringe; fitting of the peaks showed that their periodicity differed due to different film thickness (Supplementary Fig. 4); these clear oscillations from peaks represent sharp interface of TiO₂/VO₂/TiO₂ all-epitaxial heterostructures.

The low-temperature epitaxy of rutile TiO₂ films on VO₂-templated substrates was locally visualized by comparing cross-section high-resolution transmission electron microscope (HRTEM) images of both TiO_x (x < 2)/TiO₂ and TiO₂/VO₂/TiO₂ (Fig. 1c, e). Amorphous nature of TiO_x films on TiO₂ was confirmed by the diffused halo feature in fast Fourier transform (FFT) pattern from the films (green square in the right column of Fig. 1c). In contrast, sharp diffraction spots were observed in FFT pattern of the TiO₂ films on VO₂/TiO₂ (green square in the right column of Fig. 1e), and is similar with that from the TiO₂ substrates (yellow square in the right column of Fig. 1e); this observation demonstrates that epitaxial rutile TiO₂ films can be crystallized at 150 °C simply by introducing VO₂ templates on TiO₂ substrate. Considering the dramatic difference of crystallinity in those TiO₂ films grown at the same growth condition, the epitaxial TiO₂ with excellent crystallinity at 150 °C is unusual, because it formed even though thermal energy was insufficient at 150 °C to overcome the activation energy that is required to drive formation of thermodynamically stable crystalline nuclei^20,23.

To determine how this unprecedented rutile TiO₂ phase developed low-temperature epitaxy, annular bright field (ABF) scanning transmission electron microscopy (STEM) data were analyzed with TiO₂/VO₂/TiO₂ heterostructure. The sensitivity of ABF to light-weight atoms permitted visualization of oxygen atomic columns in ABF STEM (Fig. 2a)¹⁷. Magnified ABF-STEM images (red rectangles) show the typical rutile TiO₂ structure in both film and substrate; this result indicates that TiO₂ films had been fully crystallized by coherent epitaxial growth on VO₂ templates. However, the contrasts of oxygen atomic columns in VO₂ are weak and diffuse in the enlarged images of VO₂ (blue rectangle); this result implies that oxygen contents are deficient in the VO₂ template after low-temperature epitaxy of stoichiometric rutile TiO₂ films. Therefore, TiO₂ films with perfect registry of atoms were epitaxially grown on top of defective VO₂ templates; this result is contrary to the general principle that high-quality epitaxial growth is achieved by using low-defect substrates, and suggests that the chemical reaction at the interface is likely to facilitate the low-temperature epitaxy of high-quality TiO₂ films by sacrificing the initially good quality of VO₂ templates.

Fig. 2: Atomic-resolution analysis of TiO₂/VO₂/TiO₂ heterostructure.

a ABF-STEM, b HAADF-STEM and c LAADF-STEM images of rutile TiO₂ epitaxial film grown on VO₂ sacrificial template at T_G = 150 °C with [100] zone axis (scale bar = 5 nm). The sensitivity of ABF technique to light-weight atoms enables observation of oxygen atoms along the [100] zone axis. Note that the regular pattern from TiO₂ epitaxial films (top red square in a) was identical to that from TiO₂ substrates (bottom red square in a), indicating an identical atomic arrangement of films with single crystals by epitaxial growth without oxygen defects. On the contrary, the oxygen-deficient region was observed in the sacrificed VO₂ template near the TiO₂ film (blue square in a). While almost-uniform HAADF contrast was observed across the heterostructures due to similar cation atomic weight across the heterostructures (b), a noticeable strain-field-induced LAADF contrast was observed in VO₂ templates (c) sandwiched between TiO₂ films and substrates. Yellow lines in b and c are the contrast-intensity profiles of the HAADF and LAADF images from the white rectangular areas.

The formation of defective features in VO₂ template could be confirmed by comparing high-angle annular dark field (HAADF) and low-angle annular dark field (LAADF) signals of TiO₂/VO₂/TiO₂ heterostructures from STEM. Contrary to almost identical HAADF contrast across the heterostructures due to similar cation atomic weight across the heterostructures¹⁶ (Fig. 2b), the LAADF contrast was noticeably mottled in VO₂ templates (Fig. 2c) that were sandwiched between TiO₂ films and substrates. The HAADF and LAADF images had distinct intensity profiles along the film growth direction of [001] (insets in Fig. 2b, c) The differences occur because the LAADF signal is more sensitive than the HAADF signal to the frustrated atomic channeling due to oxygen vacancies (V_O)^4,24, so the contrast is blurred and brighter in the LAADF signal of VO₂ templates. Moreover, electron energy loss spectroscopy (EELS) experiments reveal that the t_2g peak of O-K edge was strongly suppressed in the entire VO₂ template (Supplementary Fig. 5b); this result directly visualize the formation of oxygen vacancies²⁵. Therefore, the combined results from LAADF contrast and EELS data in VO₂ templates confirms the significant loss of oxygen atoms from VO₂ during the low-temperature epitaxial growth of TiO₂ films in the heterostructures.

Suppression of metal-insulator transition in VO₂ templates by directional oxygen transport

Interestingly, the degree of oxygen deficiency of VO₂ templates was sensitively modulated by adjusting $pO_2$ (6 mTorr ~ 24 mTorr) during TiO₂ growth at T_G ~ 150 °C on VO₂-templated TiO₂ substrates. As observed in symmetric 2θ-ω synchrotron XRD scans in all heterostructures, Bragg peaks and Kiessig fringes from TiO₂ films were resolved from those from TiO₂ substrates and VO₂ templates with different period of oscillations in fringes (Fig. 3a), which again confirms the importance of VO₂ templates for low-temperature epitaxy of TiO₂ layers. However, unlike the almost identical peak of TiO₂ films and substrates, the (002) reflection of 14-nm-thick VO₂ template films decreased from 2θ = 51.9° (black) to 2θ = 51.5° (green) from symmetric 2θ-ω scans as $pO_2$ was reduced from 24 mTorr to 6 mTorr during the TiO₂ growth (Fig. 3a); this peak shift corresponds to ~ 0.8 % expansion of the out-of-plane lattice parameters in VO₂ templates.

Fig. 3: Structural/electronic modulation by directional ionic transport.

a Symmetrical X-ray scan and b reciprocal space mapping around (112) reflection of TiO₂ grown under $pO_2$ = 6 mTorr (denoted as H6mT) ~ 24 mTorr (denoted as H24mT) on VO₂-templated TiO₂ substrates. These results confirm that entire layers in all heterostructures are coherently strained by TiO₂ substrates, but only VO₂ peaks shifted to lower scattering angles as $pO_2$ decreased during TiO₂ growth. c Temperature-dependent sheet resistance (R_S) in all TiO₂/VO₂/TiO₂ heterostructures with TiO₂ epitaxial films grown at 6 ≤ $pO_2$ ≤ 24 mTorr on VO₂-templated TiO₂ substrates. d Lattice parameter (c) from a and temperature (T_MI) of metal-insulator transition and R_S at 270 K from c as a function of $pO_2$ during TiO₂ growth. e The formation of oxygen vacancies in VO₂ by ionic transfer across the TiO₂/VO₂ interface expanded the lattice to compensate for the larger cation radius of V³⁺ (3d²) than V⁴⁺ (3d¹), and also led to the oxygen-vacancy-induced metallization.

For more detailed structural modulation of TiO₂/VO₂/TiO₂ heterostructures, reciprocal space mapping (RSM) around the (112) reflection of (001) TiO₂ substrate was performed to obtain the information on both in-plane and out-of-plane lattice parameters by adjusting $pO_2$ during TiO₂ growth (Fig. 3b). The RSM data of all heterostructures clearly show sharp and intense (112) Bragg reflections and Kiessig fringes from TiO₂ substrate and film, and from the VO₂ films. The substrate and film peaks showed identical H (i.e., in-plane reciprocal space unit)^26,27, which implicates that entire layers in all heterostructures are coherently strained by TiO₂ substrates along the in-plane direction. Geometric phase analysis (GPA) strain quantification using obtained STEM image confirms coherent interfaces through the heterostructures (Supplementary Fig. 6). However, only the VO₂ peak shifted to lower scattering angle (i.e., characteristic of expansion of out-of-plane lattice parameters) as the $pO_2$ during TiO₂ growth was decreased; these trends are consistent with vanadium valence state switching (V⁴⁺ to V³⁺) by the formation of V_O in VO₂ templates^15,16,28.

Temperature-dependent sheet resistance R_S (T) of TiO₂/VO₂/TiO₂ heterostructures (by van der Paw methods) was measured to quantify how the accelerated oxygen deficiency in VO₂ templates affected electrical transport of the heterostructures (Fig. 3c). The heterostructure that had been formed using $pO_2$~ 24 mTorr during TiO₂ growth (denoted as H24mT hereafter) exhibited slightly suppressed MI transition in terms of R_S(T) (i.e., ~ 3.2 orders of magnitude at T_MI ~ 298 K) compared with as-grown VO₂ films without TiO₂ layers on top. On the contrary, R_S(T) of the heterostructure with $pO_2$ of 6 mTorr (denoted as H6mT) substantially dropped by just less than an order of magnitude; This result indicates that MI transition of VO₂ templates was progressively suppressed and T_MI was monotonically decreased from 298 K (H24mT) to 260 K (H6mT) by the gradual increase of V_O in VO₂ as the TiO₂ films were grown at progressively lower $pO_2$ (Fig. 3d)^4,9. As a result, the formation of V_O by oxygen ionic transfer across the interface not only expanded the lattice to compensate for the larger cation radius of V³⁺ (3d²) than V⁴⁺ (3d¹)²⁸, but also induced the metallic state at TiO₂/VO₂ interfaces even at 270 K (Fig. 3e). Structural and electrical modulation driven by V_O in VO₂ cannot be generated by simple post-annealing without TiO₂ layer growth on top; directional oxygen ionic transport indeed occurs across the TiO₂/VO₂ interfaces by forming V_O in VO₂ layer, as long as TiO₂ layers are grown on VO₂ templates (Supplementary Fig. 14).

To elucidate the origin of metallicity in a VO₂ template interfaced with a TiO₂ layer, we performed polarization-dependent x-ray absorption spectroscopy (XAS) at the V L_2,3-edges for two heterostructures that contained 2.5-nm-thick TiO₂ films grown on VO₂ templates under $pO_2$ = 24 mTorr (H24mT, Fig. 4a) and 6 mTorr (H6mT, Fig. 4b). The XAS signals at the V L_2,3-edges represent a dipole-allowed transition from the V 2p_1/2 and 2p_3/2 core level to the V 3d valence electronic states (i.e., 2p⁶3d¹ → 2p⁵3d²)^29,30,31 only from the VO₂ templates buried under layers of rutile TiO₂ owing to its element-specific character. Linearly-polarized X-rays with the polarization vector parallel (E||c) and perpendicular (E⊥c) to the out-of-plane orientation (c axis), respectively, detect the vacant d_|| and π* electron states³¹, so V_O formation also significantly affects the dichroic signal (Fig. 4a, b) related to selective orbital occupancy of d_|| induced by V-V dimerization^30,31.

Fig. 4: Oxygen-vacancy-induced isotropic orbital occupancy in VO₂ templates.

Polarization-dependent x-ray absorption spectroscopy (XAS) at the V L_2,3-edges at 270 K for two heterostructures composed of 2.5-nm-thick TiO₂ films grown on VO₂ templates under different $pO_2$ (a H24mT, b H6mT). Unlike large difference of XAS signal in H24mT at 270 K due to the orbital polarization with V-V dimerization in the monoclinic VO₂, almost no effect on the XAS signal in H6mT was observed at 270 K; this result indicates isotropic orbital filling in H6mT sample even at 270 K. The XLD (I_|| – I_⊥) are also shown at both 270 K and 320 K for c H24mT and d H6mT. Oxygen vacancies driven by directional ionic transport in H6mT tend to increase the crystal symmetry to close to rutile structure by weakening of V-V dimerization, so selective filling of d_|| (inset of c.) changes to isotropic orbital occupancy of d_|| and π* (inset of d) in VO₂ templates; the oxygen-vacancy-driven isotropic occupancy leads to metallization at 270 K.

In H24mT sample, XAS spectra collected at 320 K (T > T_MI) were similar regardless of the polarization direction of the X-ray (Supplementary Fig. 7a); this result was expected because of the isotropic orbital filling in the metallic states of VO₂ due to absence of V-V dimerization. At 270 K (T < T_MI, Fig. 3c), X-ray linear dichroism (XLD, I_|| – I_⊥) was much larger than at T = 320 K (Fig. 4a, c); this increase is a signature of orbital polarization, which is expected to result from the strong V-V dimerization in the insulating states, due to the selective filling of d_|| orbitals in VO₂ films with negligible oxygen loss (inset, Fig. 4c)^30,31. In contrast, XAS spectra of H6mT, which is the sample with the highest driving force for oxygen ion transport from VO₂ sacrificial template, show no polarization-dependence of incident X-ray (Fig. 4d) at either T = 270 K (Fig. 4b) or 320 K (Supplementary Fig. 7b). This result indicates that selective filling of d_|| orbitals did not occur below 270 K^30,31 and explains the V_O-induced metallic behavior at 270 K in H6mT: V_O tends to increase the crystal symmetry toward thermally-induced tetragonal rutile structure by weakening of V-V dimerization^15,32,33 and leads to isotropic orbital occupancy of d_|| and π*³³, (inset of Fig. 4d). The absence of selective filling in d_|| orbitals in H6mT samples provides strong evidence for V_O formation in the entire area of sacrificial VO₂ templates by directional oxygen transport from VO₂ to TiO₂.

Stoichiometric TiO₂ epitaxy induced by directional oxygen transport

To explore the influence of directional ionic transport on the quality of TiO₂ epitaxial layer grown at T_G ~ 150 °C, we evaluated the element-specific Ti L_2,3-edge XAS signal from the TiO₂ layer in TiO₂/VO₂ heterostructures. The Ti L₃ edge peak between 457.9 eV ~ 461.3 eV (i.e., related to e_g orbitals) is split into two peaks due to the distortion of the TiO₆ octahedra in rutile structure; the relative intensities of these e_g doublet (e_g¹ < e_g²) verified a rutile TiO₂ phase in both H24mT and H6mT films^34,35 (Fig. 5a), which is consistent with our results in HRXRD and STEM. The Ti L_2,3-edge signals from H6mT were more intense and sharper than from H24mT^36,37. Furthermore, the contribution of Ti³⁺ L-edge signals slightly increased the dips at 458 eV and 461 eV in the H24mT relative to those in H6mT (inset of Fig. 5a, yellow arrow);³⁶ these results reveal that rutile TiO₂ films toward the stoichiometry with low oxygen deficiency can be formed more easily by the TiO₂ growth with low $pO_2$ than with high $pO_2$. Moreover, EELS data from the top TiO₂ layers in H6mT show that the t_2g peaks of Ti-L edge from TiO₂ layers is exactly same as those from stoichiometric bulk TiO₂ substrates (reference) (Supplementary Fig. 5c); this result represents the formation of stoichiometric TiO₂ layers at the expense of oxygen deficiency in VO₂ layers.

Fig. 5: Facilitated rutile TiO₂ epitaxy by Δμ_O across the TiO₂/VO₂ interface.

a Ti L-edge XAS spectra of TiO₂ grown on VO₂ templates under different $pO_2$ (H24mT, H6mT). The Ti L_2,3-edge XAS signals from the H6mT were more intense and sharper than from H24mT. b XPS spectra of the Ti 2p core level of H24mT, H6mT. negligible Ti³⁺ contribution from the H6mT was observed compared to H24mT in TiO₂/VO₂ heterostructures, which reveals the suppression of V_O formation even at the surface of TiO₂ as a result of increased oxygen transport across TiO₂/VO₂ interface under low $pO_2$. Both XAS and XPS results reveal increased perfection of rutile TiO₂ films after growth at low $pO_2$. c First-principles density functional theory (DFT) calculations to determine values of the lower and upper limit of the chemical potential of μ_O for TiO₂ and VO₂ formation. TiO₂ is the only stable compound at − 9.624 eV ≤ μ_O ≤ − 8.767 eV. Comparison of the formation energies of V_O in rutile VO₂ and TiO₂ as a function of Fermi level in the band gap of TiO₂. d Increased thermodynamic driving force Δμ_O, assisted by high ionic kinetics k, across the interface increased the perfection of registry in the lattice of TiO₂ films by increasing “effective” $pO_2$ and lowering the activation barrier for epitaxy with concurrent emergence of a metallic VO_2-δ sacrificial templates.

In addition to bulk-sensitive XAS and EELS, X-ray photoelectron spectroscopy (XPS) of the Ti 2p core level shows better stoichiometry of rutile TiO₂ film surface grown on VO₂ templates under low $pO_2$ than under high $pO_2$ (Fig. 5b). Deconvolution of The Ti 2p core-level peaks with Ti⁴⁺ (2p_3/2 ~ 458.8 eV, 2p_1/2 ~ 464.6 eV) and Ti³⁺ (2p_3/2 ~ 457.2 eV, 2p_1/2 ~ 463.1 eV) valence states³⁸ showed negligible Ti³⁺ contribution from the H6mT was observed compared to H24mT in TiO₂/VO₂ heterostructures, which reveals the suppression of V_O formation even at the surface of TiO₂ resulting from enhanced oxygen transport across TiO₂/VO₂ interface under low $pO_2$; Both XAS and XPS results contradict the typical observation that stoichiometry could be improved in TiO₂ under high $pO_2$ by removing V_O^21,39.

Discussion

This study presents two interesting observations. (1) The increase in directional oxygen transport from VO₂ to TiO₂ with decrease in $pO_2$ during TiO₂ growth, and (2) facile formation of rutile TiO₂ epitaxial layer at extremely low temperature (≤150 °C). To identify the driving force for spontaneous oxygen loss from VO₂ templates, firstly, density functional theory (DFT) calculations were performed to determine values of the lower and upper limit of the chemical potential of oxygen (μ_O) for TiO₂ and VO₂ formation in general:

$$\frac{1}{2}\left( {E_{tot}\left[ {TiO_2(or\,VO_2)} \right] - \mu _{Ti}\left[ {Ti(or\,V)} \right]} \right) \le \mu _O\left[ {TiO_2(VO_2)} \right] \le \mu _O[O_2].$$

(1)

Our calculations predict that the lower limit of μ_O is −8.767 eV for VO₂ and −9.624 eV for TiO₂ (top of Fig. 5c), which indicates that TiO₂ is the only stable compound at −9.624 eV ≤ μ_O ≤ −8.767 eV. In the specific μ_O region in which VO₂ is thermodynamically unstable, oxygen atoms can preferentially migrate from it to stable TiO₂.

Indeed, VO₂ and TiO₂ coexist at the interfaces in the heterostructures, so the following thermodynamic reactions occur during TiO₂ growth:

$$TiO_{2 - \delta }\left( s \right) + VO_2\left( s \right) \leftrightarrow TiO_2\left( s \right) + VO_{2 - \delta }(s)$$

(2)

Thermodynamic calculations using this reaction yielded a Gibbs free energy ΔG = −18.69 kJ/mol at δ = 0.125, and −82.77 kJ/mol at δ = 0.5 at T_G = 150 °C^40,41,42. Regardless of the degree of oxygen deficiency in grown TiO_2-δ layer, oxygen ions tend to transfer spontaneously to the TiO₂ layer to equilibrate μ_O between the two layers by forming V_O in VO₂ templates. A thermodynamic driving force between TiO_2-δ and VO₂ still exists even if few monolayer of TiO₂ prevents the direct interface between two dissimilar materials, so “remote” oxygen ionic transport from VO₂ to TiO_2-δ is maintained through the few TiO₂ monolayer as long as oxygen diffusion is kinetically allowed (Supplementary Figs. 9, 10, 15)^16,43. To support our observation on the preferred formation of V_O in VO₂ templates, we also compared the formation energies of V_O (and vanadium interstitials (V_i)) in rutile VO₂ and in TiO₂ as a function of Fermi level in the band gap of rutile TiO₂ (bottom of Fig. 5c). Our calculations predict that the formation energy of V_O is at least 0.69 eV lower in VO₂ than in TiO₂ and is ~ 2.0 eV lower that of V_i in VO₂ (Supplementary Fig. 8).

It should be emphasized that the driving force for ionic flux accelerates with increase in the oxygen deficiency on the formed TiO_2-δ layers (i.e., decrease in $pO_2$ at which the film was grown) by maximizing chemical potential mismatch ($J_{O^{2 - }} \propto \Delta {\mathrm{G}}({\mathrm{or}}\,\Delta {\upmu})$)¹⁰. Due to this strengthened driving force for directional oxygen diffusion from VO₂ to TiO₂, more oxygen vacancies prefer to form in VO₂ templates during TiO₂ growth as $pO_2$ was decreased. Δμ_O across TiO₂/VO₂ interfaces in the heterostructure drives oxygen flux ($J_{O^{2 - }}$) as a directional supply of oxygen ionic radical across the interface from VO₂ sacrificial layers without the dissociation of oxygen gas molecules¹⁰. The transferred oxygen ions can “effectively” increase the oxygen partial pressure ($pO_2$) and μ_O at the TiO₂ side; paradoxically, low $pO_2$ at the ambient is likely to increase “effective” $pO_2$ (= “external” $pO_2$ from O₂ gas + “internal” $pO_2$ across the solid-solid interface) during TiO₂ growth on VO₂ templates. As a result, the increased “effective” $pO_2$ by the enhanced $J_{O^{2 - }}$ across the interface magnifies driving force for the formation of rutile TiO₂ with stoichiometry in heterostructure (Supplementary Fig. 16); lack of oxidation in the deposited TiO_2-δ species during the growth is compensated by transferred oxygen ions from the VO₂ templates below^10,44.

For heterogenous nucleation on the substrate during the film growth, the activation energy (ΔG^*) for the formation of crystalline nuclei could be significantly lowered by increasing the supply of oxygen ions (i.e., increasing the driving force for formation of rutile TiO₂) based on the following expression⁴⁵.

$$\Delta {\mathrm{G}}^ \ast = \frac{{16\pi \gamma _{fv}^3}}{{3\left( {\Delta {\mathrm{G}}_v - \Delta {\mathrm{G}}_S} \right)^2}}S(\theta )$$

(3)

where ΔG_v, γ_fv, ΔG_S, and S(θ) are the chemical free energy change for the formation of solid rutile TiO₂ nuclei, surface free energies, interfacial strain energy and a geometrical factor for heterogeneous nucleation, respectively. At very low T_G ~ 150 °C, the adatoms freeze in metastable form, so the time constant for crystallization $\big(\tau_{cryst}, \, {\mathrm{i.e.}}, \, \frac{1}{{\tau _{cryst}}} = {\mathrm{A}} \cdot {\mathrm{exp}}\big( { - \frac{{\Delta G^ \ast }}{{k_BT_g}}} \big)\big)$ becomes extremely long due to the high ΔG^* and insufficient thermal energy; the amorphized or metastable TiO₂ films (with nonequilibrium structure) are likely to form due to the kinetic hindrance of crystal formation (e.g., insufficient movement of ablated TiO₂ adatoms and/or limited reaction with O₂ gas) as observed in our TiO₂/TiO₂ homostructure.

In our case, VO₂ sacrificial layers from the bottom acts as epitaxial templates for rutile TiO₂. They also sacrifice themselves by forming V_O in VO₂ and thereby supply high concentration of oxygen ions to TiO_2-δ by magnifying the driving force for the oxygen transport at the VO₂/TiO₂ interfaces. Since γ_fv and ΔG_S are unlikely to be changed regardless of the existence of VO₂ templates (Supplementary Fig. 6 and Fig. 15), greater ΔG_v induced by $J_{O^{2 - }}$ significantly reduce ΔG^* in heterostructure than in homostructure. As a result, the large reduction of τ_cryst by the decrease in ΔG^* enables unprecedented epitaxy of high-temperature-stabilized TiO₂ rutile phase^19,23 at extremely low T_G ~ 150 °C by changing the initial stoichiometry of two oxides with different μ across the interface.

In addition to thermodynamic viewpoint, the transport of charged oxygen ions is kinetically facilitated along the crystallographic [001] direction, which has open channels in anisotropic rutile VO₂ and TiO₂^15,16. One-dimensional empty channels are aligned along the c axis in our TiO₂/VO₂ heterostructures and provide the advantage of removing significant amounts of oxygen ions due to a high oxygen diffusion coefficient. Thus, the growth direction of [001] TiO₂ films should strongly accelerate the out-diffusion kinetics of oxygen transport (increased k in Fig. 1a) from the VO₂ sacrificial templates as a result of mismatch in chemical potential (Δμ_O in Fig. 1a); significant reduction in oxygen transport across the oxide interface along the [001] direction kinetically facilitates the decrease in ΔG^* for nucleation of rutile TiO₂ phase even at low T_G to support epitaxial growth of rutile TiO₂ films on VO₂ templates.

On the other hand, the kinetics of oxygen transport will eventually limit our epitaxial growth based on “internal” oxygen transport across the TiO₂/VO₂ interface (Supplementary Figs. 11, 12, 13). Since the oxygens should be supplied from the TiO₂/VO₂ interfaces through the intervening TiO₂ layers, the thickness of “epitaxial” TiO₂ will be limited by oxygen diffusion through the intervening TiO₂ layer. In fact, while the “epitaxial” thickness linearly increased with growth time in the TiO₂ films grown at 300 °C, the “epitaxial” thickness appears to be saturated to be ~ 10 nm as a “critical” thickness at t_film > 10 nm at T_G ~ 150 °C due to scarce source of “internal” oxygen transport ($J_{O^{2 - }}$) even if the growth time increases (Supplementary Fig. 11). Therefore, the existence of “critical” thickness provides the convincing evidence of our unprecedented low-temperature epitaxy driven by direction oxygen ionic transport.

In summary, unconventional low-temperature epitaxy of rutile TiO₂ films was achieved by exploiting the directional transport of oxygen ions across TiO₂/VO₂ heterointerfaces. The thermodynamic driving force, assisted by facile ionic pathway along oxygen channel, across the interface enabled more perfect registry in the lattice of TiO₂ films by lowering the activation barrier for stable nuclei, with concurrent emergence of a metallic TiO₂/VO₂ heterostructures. Contrary to typical experimental condition to obtain TiO₂ with better stoichiometry, interestingly, V_O formation was diminished under low external $pO_2$, because the accelerated chemical potential mismatch (Δμ_O) under low external $pO_2$ significantly increased “effective” $pO_2$ by the internal oxygen transport across the TiO₂/VO₂ interface. Therefore, the controlled ionic transport by Δμ_O may offer an opportunity to design a new heterostructure with different degree of freedom at the interfaces as a result of tuning of ionic defects, and also to stabilize thermal-energy-requiring phases simply by interfacing with dissimilar materials with different thermodynamic and kinetic driving force of ionic defects.

Methods

Synthesis of epitaxial TiO₂/VO₂ heterostructures on TiO₂ substrates

Epitaxial VO₂ thin films (10–14 nm thick) were grown on (001) TiO₂ single-crystal substrates, followed by the growth TiO₂ films (2.5–60 nm thick) by pulsed laser deposition (PLD). The stoichiometric targets for the synthesis of heterostructures were prepared by sintering stoichiometric powders of V₂O₅ (99.99%, Sigma Aldrich) at 600 °C for 6 h and TiO₂ (99.95%, Sigma Aldrich) at 1100 °C for 4 h. First, (001) TiO₂ single crystal substrates (Shinkosha CO., LTD) were loaded into the PLD chamber, which was then evacuated to a base pressure of ~ 1 × 10⁻⁶ Torr. Then, the rotating V₂O₅ targets were ablated by focusing KrF excimer laser (Coherent Compex Pro 102 F, λ = 248 nm) with a fluence of 1 J/cm² and repetition rate of 1 Hz. The VO₂ growth was performed at fixed $pO_2$ = 12 mTorr and 300 °C, which was selected to induce a steep metal-insulator transition near room temperature from coherently tensile-strained VO₂ films. After VO₂ growth, the substrate temperature was quenched to 50 ~ 150 °C. Subsequently, TiO₂ films were grown on VO₂ templates under 6 mTorr ≤ $pO_2$ ≤ 24 mTorr to control the thermodynamic driving force for oxygen ionic transport across the interface at low temperature (T_G = 50–150 °C). After the growth of heterostructures, the samples were cooled down to room temperature with rate of 20 °C/min.

Structural and electrical characterization of heterostructures

To characterize crystal-structure modulation in TiO₂/VO₂/TiO₂ heterostructures with different degrees of chemical potential mismatch using $pO_2$ during TiO₂ growth, high-resolution X-ray scattering measurements were performed using synchrotron radiation at 3A MP-XRS (λ ~ 0.11145 nm, energy ~ 11,125 keV at Si (111)) and at 3D XRD (λ ~ 0.12398 nm, energy ~ 10 keV at Si (111) beamline of Pohang Light Source-II (PLS-II, Pohang, Republic of Korea), and using an in-house HRXRD (Bruker Discover 8 X-ray diffractometer) with Cu K_α1 radiation (λ = 0.15406 nm). The detailed information on in-plane and out-of-plane lattice parameters and strain states of each film in the heterostructures were obtained by using both symmetric 2θ-ω scan and asymmetric reciprocal space mapping (RSM) around the (112) reflection. The simulation of symmetric 2θ-ω scans was performed using LEPTOS software program. The surface morphology of the films was observed using an atomic force microscope (AFM, VEECO Dimension 3100).

For atomic resolution analysis of crystal structure, the samples were prepared using a dual-beam focused ion beam (FIB) system (Helios G3, FEI). HRTEM and STEM analyses (JEM-ARM200F, JEOL) were performed at 200 kV equipped with a 5^th order aberration corrector (ASCOR, CEOS GmbH) for forming 0.7 Å probe. The collection semi-angles were 68 to 280 mrad for HAADF, 27 to 110 mrad for LAADF and 10 to 20 mrad for ABF. The obtained raw images were band-pass filtered to reduce background noise (HREM Research Inc.).

The sheet resistance R_S was measured as a function of temperature during the heating and cooling from 250K to 340K using Hall measurement system. Measurements were carried out in van der Pauw geometry with square samples (5 mm × 5 mm) and indium Ohmic contacts (<1 mm × 1 mm) in the sample corners. The four-terminal resistances were measured using a 10-μA current.

To investigate electronic structure of TiO₂/VO₂ heterostructures, X-ray absorption spectroscopy (XAS) and linear dichroism (XLD) were performed using high sensitivity at the 2A MS beamline at PLS-II. The total electron yield mode with an energy resolution of ~0.1 eV was used for both measurements at a base pressure of 5 × 10⁻¹⁰ Torr in the analysis chamber by measuring the sample current (I₁) divided by the beam current (I_o) to remove the variation of beam intensity. The linear dichroism of V L_2,3-edge was carried out by using horizontally-polarized or vertically-polarized X-ray beams with photon incidence angle of 22.5° at the measurement temperatures below (270 K) and above the T_MI (320 K) of as-grown VO₂ films. And then, the Ti L_2,3-edges XAS measurements were performed on TiO₂/VO₂/TiO₂ heterostructures; photon incidence angle was 45°, and measurement temperatures were 270 K and 320 K. Due to its element-resolved characterization with several nanometer probing depth, Ti L_2,3-edges spectra were obtained only from the TiO₂ epitaxial films on top.

To evaluate the surface stoichiometry of the TiO₂ rutile films as a result of directional oxygen transport from VO₂ films in our TiO₂/VO₂/TiO₂, XPS spectra of Ti 2p core level were acquired on the 4A2 SARPES and 4D PES beam line (PLS-II) in an ultra-high vacuum chamber (2 × 10^-10 Torr). Before measurement, we carefully removed possible contaminants by using gentle Ar surface treatment. Ion sputtering was performed for 4 min and 20 min in the preparation chamber under the Ar pressure of 8 × 10⁻⁶ Torr at anode voltage of 500V and 2.5 kV. For collect Ti 3d spectra, we measured from 475 eV to 445 eV with 50-meV steps at 300 K. The measured spectra were deconvoluted using XPSPEAK41 software.

First-principles calculation

First-principles density functional theory (DFT) calculations were performed using the Projector Augmented Wave (PAW) method and the generalized gradient approximation of Perdew, Burke, and Ernzerhof (PBE) for the exchange-correlation potential as implemented in Vienna Ab-initio Simulation Package (VASP) code⁴⁶. Periodic boundary condition and Monkhorst-Pack k-point sampling with a Г-centered k-point grid of up to 8 × 8 × 8 was used for the Brillouin zone integration. An energy cutoff of 450 eV was used for the plane-wave representation of the wavefunctions and the 3s electrons of V and Ti ions were considered as valence electrons. A Hubbard U correction term was applied to the V (U = 3.25 eV) and Ti (U = 3.00 eV) to properly reproduce the strong on-site Columbic repulsion of 3d-electrons⁴⁷. Atomic structures were relaxed until all Hellman-Feynman forces were below 0.01 eV/Å. The optimized lattice parameters are a = 4.67 Å and c = 2.52 Å for Ti metal and a = 3.31 Å for V metal. The optimized lattice parameters are a = 4.559 Å and c = 2.889 Å for rutile VO₂, and a = 4.608 Å and c = 2.989 Å for rutile TiO₂.