Interfaces formed by two dissimilar materials can break the translational symmetry and thereby provide an opportunity to develop functionality that is unachievable in bulk materials1,2,3,4,5. 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 materials2,3. 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 interfaces2,6, which is the basic principle of heterojunction field effect transistors (HFETs)7.

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

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 TiOX/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 TiOX /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 TiOX 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 (ΔμO) 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 ΔμO 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 interfaces9,14,15,16. 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 (VO) 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, VO concentrations that develop in LaNiO3, LaTiO3, and In2O3 can be modulated by the directional oxygen flow to the adjacent layers10,11,12.

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 μO10,11,14. The redistribution of charged vacancies can screen the electric fields that ΔμO causes. Therefore, the massive redistribution of charged ions can be accelerated by extremely increasing the thermodynamic ΔμO 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 interfaces9,13,14.

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 TG (<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 ΔμO, 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 VO2-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.


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 (ΔRS ~ 103.3) occurred on the VO2 films at TMI ~ 298 K (Supplementary Fig. 2b); this result indicates the formation of coherently tensile-strained VO2 films with high crystal quality and negligible VO9,15,17,18.

Then, the 6 nm-thick TiO2 films were grown at low TG ~ 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 TG = 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 formation19,20,21,22,23. 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 TG = 150 °C.

In contrast, rutile TiO2 epitaxial films were strikingly stabilized by introducing the VO2 layers on TiO2 substrate at the same condition with TG = 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 TiOx (x < 2)/TiO2 and TiO2/VO2/TiO2 (Fig. 1c, e). Amorphous nature of TiOx 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 nuclei20,23.

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.
figure 2

a ABF-STEM, b HAADF-STEM and c LAADF-STEM images of rutile TiO2 epitaxial film grown on VO2 sacrificial template at TG = 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 (VO)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 t2g 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 TG ~ 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.
figure 3

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 (RS) 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 (TMI) of metal-insulator transition and RS 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 VO in VO2 templates15,16,28.

Temperature-dependent sheet resistance RS (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 RS(T) (i.e., ~ 3.2 orders of magnitude at TMI ~ 298 K) compared with as-grown VO2 films without TiO2 layers on top. On the contrary, RS(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 TMI was monotonically decreased from 298 K (H24mT) to 260 K (H6mT) by the gradual increase of VO in VO2 as the TiO2 films were grown at progressively lower \(pO_2\) (Fig. 3d)4,9. As a result, the formation of VO 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 VO 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 VO 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 L2,3-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 L2,3-edges represent a dipole-allowed transition from the V 2p1/2 and 2p3/2 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 (Ec) to the out-of-plane orientation (c axis), respectively, detect the vacant d|| and π* electron states31, so VO formation also significantly affects the dichroic signal (Fig. 4a, b) related to selective orbital occupancy of d|| induced by V-V dimerization30,31.

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

Polarization-dependent x-ray absorption spectroscopy (XAS) at the V L2,3-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 > TMI) 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 < TMI, 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 K30,31 and explains the VO-induced metallic behavior at 270 K in H6mT: VO tends to increase the crystal symmetry toward thermally-induced tetragonal rutile structure by weakening of V-V dimerization15,32,33 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 VO 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 TG ~ 150 °C, we evaluated the element-specific Ti L2,3-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 eg orbitals) is split into two peaks due to the distortion of the TiO6 octahedra in rutile structure; the relative intensities of these eg doublet (eg1 < eg2) verified a rutile TiO2 phase in both H24mT and H6mT films34,35 (Fig. 5a), which is consistent with our results in HRXRD and STEM. The Ti L2,3-edge signals from H6mT were more intense and sharper than from H24mT36,37. 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 t2g 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 ΔμO across the TiO2/VO2 interface.
figure 5

a Ti L-edge XAS spectra of TiO2 grown on VO2 templates under different \(pO_2\) (H24mT, H6mT). The Ti L2,3-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 VO 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 μO for TiO2 and VO2 formation. TiO2 is the only stable compound at − 9.624 eV ≤ μO ≤ − 8.767 eV. Comparison of the formation energies of VO in rutile VO2 and TiO2 as a function of Fermi level in the band gap of TiO2. d Increased thermodynamic driving force ΔμO, 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 VO2-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+ (2p3/2 ~ 458.8 eV, 2p1/2 ~ 464.6 eV) and Ti3+ (2p3/2 ~ 457.2 eV, 2p1/2 ~ 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 VO 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 VO21,39.


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 (μO) for TiO2 and VO2 formation in general:

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

Our calculations predict that the lower limit of μO 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 ≤ μO ≤ −8.767 eV. In the specific μO 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:

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

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 TG = 150 °C40,41,42. Regardless of the degree of oxygen deficiency in grown TiO2-δ layer, oxygen ions tend to transfer spontaneously to the TiO2 layer to equilibrate μO between the two layers by forming VO in VO2 templates. A thermodynamic driving force between TiO2-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 TiO2-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 VO in VO2 templates, we also compared the formation energies of VO (and vanadium interstitials (Vi)) 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 VO is at least 0.69 eV lower in VO2 than in TiO2 and is ~ 2.0 eV lower that of Vi 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 TiO2-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. ΔμO 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 μO 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 TiO2-δ species during the growth is compensated by transferred oxygen ions from the VO2 templates below10,44.

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.

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

where ΔGv, γfv, ΔGS, 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 TG ~ 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 VO in VO2 and thereby supply high concentration of oxygen ions to TiO2-δ by magnifying the driving force for the oxygen transport at the VO2/TiO2 interfaces. Since γfv and ΔGS are unlikely to be changed regardless of the existence of VO2 templates (Supplementary Fig. 6 and Fig. 15), greater ΔGv induced by \(J_{O^{2 - }}\) significantly reduce ΔG* in heterostructure than in homostructure. As a result, the large reduction of τcryst by the decrease in ΔG* enables unprecedented epitaxy of high-temperature-stabilized TiO2 rutile phase19,23 at extremely low TG ~ 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 TiO215,16. 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 (ΔμO 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 TG 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 tfilm > 10 nm at TG ~ 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, VO formation was diminished under low external \(pO_2\), because the accelerated chemical potential mismatch (ΔμO) 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 ΔμO 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.


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 (TG = 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 Kα1 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 5th 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 RS 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 (Io) to remove the variation of beam intensity. The linear dichroism of V L2,3-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 TMI (320 K) of as-grown VO2 films. And then, the Ti L2,3-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 L2,3-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.