Nanoscale Phase Mixture and Multi fi eld-Induced Topotactic Phase Transformation in SrFeO

Nanoscale Phase Mixture and Multi fi eld-Induced Topotactic Phase Transformation in SrFeO

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Junjiang Tian,

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Yang Zhang,

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Zhen Fan,* Haijun Wu,* Lei Zhao, Jingjing Rao, Zuhuang Chen, Haizhong Guo, Xubing Lu, Guofu Zhou, Stephen J. Pennycook, Xingsen Gao, and Jun-Ming Liu

Cite This:ACS Appl. Mater. Interfaces2020, 12, 21883−21893 Read Online

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^sı Supporting Information

ABSTRACT: Nanoscale phase mixtures in transition-metal oxides (TMOs) often render these materials susceptible to external stimuli (electric field, mechanical stress, etc.), which can lead to rich functional properties and device applications. Here, direct observation and multifield manipulation of a nanoscale mixture of brownmillerite SrFeO_2.5 (BM-SFO) and perovskite SrFeO₃ (PV- SFO) phases in SrFeO_x (SFO) epitaxial thin films are reported.

The mixed-phase SFO film in its pristine state exhibits a nanoscaffold structure consisting of PV-SFO nanodomains embedded in the BM-SFO matrix. This nanoscaffold structure produces gridlike patterns in the current and electrochemical strain maps, owing to the strikingly different electrical and electro- chemical properties of BM-SFO and PV-SFO. Moreover, electric

field control of reversible topotactic phase transformation between BM-SFO and PV-SFO is demonstrated by electric-field-induced reversible changes in surface height, conductance, and electrochemical strain response. In addition, it is also shown that the BM-SFO

→PV-SFO phase transformation can be enabled by applying mechanical stress. This study therefore not only identifies a strong nanometric structure−property correlation in the mixed-phase SFO but also offers a new paradigm for the multifield control of topotactic phase transformation.

KEYWORDS: topotactic phase transformation, SrFeO_x, mixed phases, resistive switching, electrochemical strain

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INTRODUCTION

Transition-metal oxides (TMOs) exhibit various competing phases and plenty of fascinating properties. A TMO with a rationally designed phase coexistence often displays exotic properties that are unprecedented in its single-phase counter- parts, such as giant piezoelectric response,^1,2colossal magneto- resistance,^3,4 superconductivity,⁵ and enhanced electrochem- ical activity.^6,7 Therefore, recently, TMOs with mixed phases have been subjected to extensive research in the fields of physics, electronics, chemistry, and so forth.

So far, a variety of approaches have been developed to obtain mixed phases in TMOs, such as chemical alloying,^8,9 strain engineering,¹⁰⁻¹²and oxygen stoichiometry tuning.¹³⁻¹⁷ For example, when PbZrO₃is alloyed with PbTiO₃at a Zr/Ti ratio of 52/48, i.e., the so-called morphotropic phase boundary, both tetragonal and rhombohedral phases coexist and anomalously high piezoelectric and dielectric coefficients (d₃₃ > 500 pC/N and εr > 1000, respectively) arise.⁸ In addition, when BiFeO₃undertakes a critical compressive strain of 4.5%, a mixture of supertetragonal and rhombohedral phases is formed, which yields giant electromechanical response,¹¹ electrically controlled ferromagnetism,¹⁸ and flexoelectric photovoltaic effect.¹⁹ Although chemical alloying and strain

engineering have proven effective in developing mixed phases, they can apply only to a limited number of TMO systems.

Oxygen stoichiometry tuning, however, appears suitable for almost all TMOs, and it can beflexibly applied, either during or after material synthesis.

An ideal platform for implementing the oxygen stoichiom- etry tuning is a family of TMOs exhibiting two interconvertible topotactic phases, namely, perovskite (PV) ABO₃ and brownmillerite (BM) ABO_2.5phases. The PV phase comprises corner-linked BO₆octahedra, while the BM phase consists of alternately stacked layers of BO₆ octahedra and BO₄ tetrahedra.²⁰ By gaining or losing oxygen, the BM and PV phases can be reversibly transformed to each other. Such topotactic phase transformation induces not only structural changes but also dramatic changes in electronic and magnetic properties²¹ because, typically, the BM phase is an

Received: February 26, 2020 Accepted: April 21, 2020 Published: April 21, 2020 Downloaded via NANJING UNIV on June 16, 2020 at 09:06:35 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

antiferromagnetic insulator²² and the PV phase is a ferromagnetic metal^23,24in their ground states. This promises a wide range of applications such as fuel cells,²⁵ catalysts,¹³ resistive switching memories,²⁶⁻²⁸ and magnetoelectric devices.²⁹ Therefore, both BM and PV phases, and the topotactic phase transformation between them via oxygen stoichiometry tuning, have attracted great interest and stimulated ongoing research in the past few years.^15,30−39

Although the emergence of BM + PV mixed phases during the topotactic phase transformation has been observed previously,⁴⁰⁻⁴⁴ the structure−property correlation at the microscopic level in the mixed phases still lacks in-depth research. Moreover, how the mixed phases respond to external stimuli (including electrical and mechanical stimuli) is also important to understand because exotic phenomena such as large resistance change and giant electrochemical strain may arise owing to the electro-chemo-mechanical coupling in the mixed phases.

Herein, we choose SrFeO_x (SFO) as a model system to study the BM + PV mixed phases because the small Gibbs free energy difference between SrFeO_2.5 (BM-SFO) and SrFeO₃ (PV-SFO) (∼100 meV³¹) may lead to an easy access to the mixed phases. Byfinely controlling the oxygen pressure during the growth of SFO epitaxial thin films, we develop a self- assembled nanoscale mixture of BM-SFO and PV-SFO phases.

The mixed-phase film (pristine state) exhibits a nanoscaffold structure consisting of PV-SFO nanodomains embedded into

the BM-SFO matrix, and this novel structure gives rise to gridlike patterns in both current and electrochemical strain maps. Then, we observe nanoscale phase transformation and associated changes in current and electrochemical strain response upon applications of both electrical and mechanical stimuli, implying great application potential in many electronic devices such as memories, actuators, and sensors.

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RESULTS AND DISCUSSION

The crystal structures of BM-SFO and PV-SFO are sketched in Figure 1a, showing that BM-SFO contains alternately stacked layers of FeO₆octahedra and FeO₄tetrahedra, while PV-SFO consists of only corner-linked FeO₆octahedra. To obtain the coexistence of these two different phases, the SFOfilms were grown on SrRuO₃ (SRO)-buffered (001) SrTiO₃ (STO) substrates under a specific oxygen pressure (pO₂) of 3 Pa. The obtained SFOfilms werefirst characterized by X-ray diffraction (XRD), as shown inFigure 1b. Two sets of diffractions peaks from SFO are clearly observed: (i) the half-order peaks, including (020)_o, (040)_o, (060)_o, (080)_o, and (0100)_o from BM-SFO, and (ii) the (001)_pc and (002)_pc peaks from PV- SFO. Here, the subscripts“o” and“pc” denote orthorhombic and pseudocubic unit cells, respectively. The half-order peaks stem from the alternate stacking of oxygen polyhedral layers, which is a fingerprint of the BM structure. Thus, the observation of these peaks indicates that the BM-SFO phase with the b_o-axis oriented along the out-of-plane direction is Figure 1.Phase analysis of the mixed-phase SFOfilm. (a) Schematics of crystal structures of BM-SFO and PV-SFO viewed along the⟨110⟩pc

direction. (b) XRDθ−2θdiffraction pattern, (c) RSM around (103)_pcreflection of STO, and (d) RSM around (1101)_oreflection of BM-SFO for a mixed-phase SFOfilm. The pink and cyan bars in (b) indicate the positions of (0k0)_opeaks from bulk BM-SFO (PDF#17-0932) and (00l)_pcpeaks from bulk PV-SFO (PDF#40-0905), respectively. AFM topography images of the mixed-phase SFOfilm (e) in the pristine state and (f) after the +7 V electrical writing.

formed. According to the peak positions, the lattice constantb_o of BM-SFO is calculated as∼15.67 Å. On the other hand, the (001)_pcand (002)_pcpeaks are typical reflections from the PV- SFO phase with the c_pc-axis oriented along the out-of-plane direction. The lattice constantc_pcof PV-SFO calculated from the peak positions is∼3.85 Å. The calculated values ofb_oand c_pcagree well with their respective bulk values (for bulk BM- SFO:²⁰b_o=∼15.59 Å,a_o=∼5.67 Å, andc_o=∼5.53 Å; and for bulk PV-SFO:^45,46a_pc=b_pc=c_pc= 3.83−3.85 Å). Therefore,

the XRD results suggest that the SFOfilm grown underpO₂of 3 Pa presents BM + PV mixed phases. Notably, the precise control overpO₂during thefilm growth is critical to obtain the mixed phases because pO₂ largely controls the oxygen stoichiometry of the deposited film. As pO₂ decreases (increases), the fraction of the BM-SFO (PV-SFO) phase tends to increase and eventually a pure BM-SFO (PV-SFO) phase is obtained, as shown inFigure S1.

Figure 2.Hierarchical structures of nanoscale phase mixtures. (a) Low-magnification STEM ABF plan-view image showing two regions, marked as A and B. (b, c) Electron diffraction patterns from the two regions marked as A and B in (a), respectively. The black circles indicate the overlapped reflections from both PV-SFO and BM-SFO phases, while the other superlattice reflections are only from the BM-SFO phase. Particularly, the 1/

4{100}_pc and 1/2{100}_pc diffraction spots from the BM-SFO phase are indicated by the red and black arrows, respectively. (d) Medium- magnification STEM ABF image showing one B region surrounded with four A regions. (e) High-magnification STEM ABF image focusing on one A region surrounded with two B regions. (f1, f2) FFT images obtained from the red and green boxes in (e), respectively. (g1, h1) Atomically resolved STEM ABF and HAADF images from regions A and B, respectively. (g2, h2) FFT patterns corresponding to g1 and h1, respectively. (i1− i3) Atomically resolved STEM HAADF images of three regions showing the coexistence of BM and PV phases.

Then, to further characterize the mixed phases, reciprocal space mapping (RSM) was measured around the (103)_pc reflection of the STO substrate, as shown in Figure 1c. The middle spot with the highest intensity is the STO (103)_pcpeak overlapped with the BM-SFO (1121)_o peak. The upper spot corresponds to the (103)_pcpeak of PV-SFO, while the lower spot corresponds to the SRO (103)_pcpeak. The in-plane lattice constanta_pc(or b_pc equivalently) of PV-SFO can be directly calculated from the (103)_pcpeak to be∼3.91 Å. To accurately calculate the in-plane lattice constanta_o(orc_oequivalently) of BM-SFO without influence from the peak overlap, an alternative peak of BM-SFO, i.e., the (1101)_o peak, was scanned (seeFigure 1d). The value ofa_ois calculated as∼5.52 Å, which is about √2 × 3.91 Å. Therefore, the lattice parameters of both the BM-SFO and PV-SFO phases are fully determined. The in-plane lattice parameters of BM-SFO and PV-SFO (in pseudocubic unit cells) are quite close to that of STO (∼3.905 Å), suggesting that both the BM-SFO and PV- SFO phases are highly constrained by the STO substrate.

To understand how the BM-SFO and PV-SFO phases are mixed together at the nanoscale, atomic force microscopy (AFM) and scanning transmission electron microscopy (STEM) characterizations were performed.Figure 1e presents an AFM topography image of the mixed-phase SFOfilm in its pristine state. A novel nanoscaffold structure is observed, where well-ordered pits (dark regions) distribute ubiquitously in the matrix (bright regions). The pits possess an average diameter of∼150 nm and an average depth of∼0.8 nm. The pits and matrix may contain different phases, and the surface height contrast may come from the different out-of-plane lattice spacings (d_oop) of PV-SFO and BM-SFO. For PV-SFO, d_oop=c_pc=∼3.85 Å. For BM-SFO, two lattice orientations are considered: (i) when theb_o-axis is oriented along the out-of- plane direction,d_oop=b_o/4 =∼3.92 Å, and (ii) when theb_o- axis is oriented along the in-plane direction,d_oop≈ a_o/√2 =

∼3.91 Å. The first orientation has already been revealed by XRD, while the second one will later be revealed by STEM.

Nevertheless, regardless of the orientation, thed_oop values of the BM-SFO phases are always larger than that of the PV-SFO phase. A surface height−phase correlation is therefore speculated: the pit regions might contain mainly PV-SFO whereas the matrix is mainly composed of BM-SFO. This speculation will be verified later by the STEM study, and it is also supported by the observation that the SFO film grown under higher pO₂ exhibits a higher fraction of the PV-SFO phase and also a larger total area of pit regions (seeFigure S1).

Upon electrically writing with an AFM tip bias of +7 V, the morphology of the mixed-phase SFO film changes into a cracked-earth-like pattern (Figure 1f), manifesting as an assembly of rectangular bulges separated by crossed channels.

Again, according to the fact that BM-SFO has a largerd_oopthan that of PV-SFO, the bulge and channel regions may mainly contain BM-SFO and PV-SFO, respectively. ComparingFigure 1e,f, the BM-SFO domains expand in area, while the PV-SFO domains shrink in area, pointing to an electrically induced PV- SFO → BM-SFO phase transformation (to be demonstrated later). Therefore, the mixed-phase SFO film can exhibit two types of morphologies: the nanoscaffold type in the pristine state and the cracked-earth type after the +7 V electrical writing.

While AFM can identify different domains in the mixed- phase SFOfilm, the microstructure in each domain is largely speculated and direct evidence is still lacking. To directly probe

the microstructure, STEM plan views of the mixed-phase SFO film were performed. Figure 2a presents a low-magnification STEM annular bright-field (ABF) plan-view image, which reveals a polydomain pattern containing rectangular domains (marked as B) separated by crossed stripes (marked as A).

This pattern well resembles the cracked-earth-type morphol- ogy observed under AFM in the mixed-phase SFO film after electrical writing (Figure 1f). The phases in regions A and B were investigated by taking electron diffraction patterns from these two regions. As shown in Figure 2b,c, both electron diffraction patterns show the coexistence of BM-SFO and PV- SFO phases. The difference is that Figure 2b from region A shows the PV phase as the majority phase, while Figure 2c from region B shows the BM phase as the majority phase, because the 1/4{100}_pc diffraction spots (characteristic reflections from the BM phase) are clearly observed inFigure 2c, while they are nearly absent inFigure 2b.

Figure 2d presents a medium-magnification STEM ABF image focusing on a B region surrounded with four A regions, further revealing the nanoscale phase mixtures in both regions.

At higher magnification (Figure 2e), one can see that region B contains a large number of superlattice stripes, while region A contains much fewer. Because the superlattice stripes are the fingerprint of the BM phase, regions A and B thus contain PV- SFO and BM-SFO as the majority phases, respectively. To further confirm it, the fast Fourier transformation (FFT) patterns were taken from these two regions.Figure 2f2 from region B presents strong 1/2{100}_pc diffraction spots belonging uniquely to the BM phase, but these diffraction spots become much weaker inFigure 2f1 from region A. It is therefore concluded that the PV-SFO and BM-SFO phases are the majority phases in regions A and B, respectively.

To further investigate thefine structures in regions A and B, aberration-corrected STEM ABF and high-angle annular dark- field (HAADF) imaging modes, which can produce contrast interpretable by mass thickness (the number of atoms) or the Z contrast,⁴⁷⁻⁴⁹ were employed. The atomically resolved STEM images, as shown inFigure 2g1,h1, clearly reveal that the total area of the PV-SFO domains is larger than that of the BM-SFO domains in region A, but the opposite is true in region B. In addition, region B displays much stronger 1/

4{100}_pc diffraction spots than those of region A (see comparison inFigure 2g2,h2), reflecting a higher fraction of the BM-SFO phase in region B than that in region A. These results confirm that the PV-SFO and BM-SFO phases prevail in regions A and B, respectively.Figure 2g1,h1 also shows the superlattice stripes running along two perpendicular ⟨100⟩ directions, indicating the presence of BM-SFO domains with in-plane-oriented b_o-axes (they were not revealed by our previous XRD study, probably because their amount was very low in the pristine state but increased significantly after the STEM sample preparation; see a detailed discussion in Supplementary Note S1). Interestingly, these BM-SFO domains, together with the PV-SFO domains, form different beautiful architectures at the nanoscale, as shown in Figure 2i1−i3. Detailed analyses on the atomic structure and strain distribution⁵⁰⁻⁵³ near the BM/PV interface are provided in Figures S2−S4. Overall, the above STEM results reveal complex hierarchical phase structures in the mixed-phase SFO film, which consists of rectangular domains (region B) with BM-SFO as the majority phase and crossed stripes (region A) with PV-SFO as the majority phase.

By correlating the above STEM results with the cracked- earth-type morphology observed via AFM (Figure 1f), one can deduce that the rectangular bulge regions mainly contain BM- SFO, while the crossed channel regions mainly contain PV-

SFO. This well validates the above-mentioned surface height− phase correlation: the regions with a higher (lower) surface height mainly contain BM-SFO (PV-SFO). Note that the BM- SFO phases with different orientations are not further Figure 3.Correlations between phase, conductance, and electrochemical strain response in the mix-phase SFOfilm (pristine state). (a) AFM topography image of a 5×5μm²area and (b) C-AFM current map taken over this area (read voltage +3 V). (c) AFM topography image of another 5×5μm²area and ESM (d) amplitude and (e) phase images taken over this area.

Figure 4.Electric-field-induced concurrent changes in phase, conductance, and electrochemical strain response. (a, d) AFM topography image and C-AFM current map, respectively, taken after writing a 3.5×3μm²region with +7 V. (b, e) AFM topography image and C-AFM current map, respectively, taken after rewriting an inner 1.5×1.5μm²region with−7 V. In (d) and (e), the read voltage is +3 V. (c) Surface height profile along the red line in (b). (f) Current profile along the blue line in (e). In (c) and (f), the bold black lines are drawn for eye guidance. (g) AFM topography image and ESM (h) amplitude and (i) phase images taken after writing outer (2×2μm²) and inner (1×1μm²) regions with +7 and

−7 V, respectively.

distinguished because thed_oopvalues of the BM-SFO phase are always larger than those of the PV-SFO phase, as discussed above. This surface height−phase correlation may also reasonably apply to the nanoscaffold-type morphology, although it is not directly observed in STEM. Why the cracked-earth-type morphology rather than the nanoscaffold- type morphology is observed in STEM may stem from phase transformation caused by the STEM sample preparation (see a detailed discussion in Supplementary NoteS1).

As mentioned in theIntroductionsection, the BM-SFO and PV-SFO phases exhibit starkly different conductivities. It is thus of interest to look into the current distribution in the mixed-phase SFOfilm. We started with the mixed-phase SFO film (pristine state) that exhibits a nanoscaffold structure.

Figure 3a,b presents the topography of a 5 ×5μm²area and the current map taken over this area using conductive AFM (C-AFM), respectively. Interestingly, the current map displays a gridlike pattern having almost one-to-one correspondence with the topography, namely, the pit regions exhibit higher currents (30−80 pA), while the matrix exhibits lower currents (<10 pA). This in turn confirms that the pit and matrix regions mainly consist of PV-SFO and BM-SFO, respectively, because the PV-SFO phase is more conductive than the BM-SFO phase.²⁶⁻²⁸ Therefore, the mixed-phase SFO film (pristine state) looks like an ordered assembly of nanocables with PV- SFO conducting cores surrounded by BM-SFO insulating shells. Note that the current enhancement in the pits compared to that in the matrix is not simply due to the thickness reduction because the surface height difference between the pits and the matrix is only∼0.8 nm (∼1.3% of the totalfilm thickness). The thickness reduction alone is thus unable to cause about 1 order of magnitude enhancement in current. In addition, the topographic crosstalk during tip scanning is excluded as a factor leading to the current enhancement because the enhanced currents are observed in the whole pit regions (rather than only at the boundaries between the pits and the matrix⁵⁴) and these currents are quite stable (Figure S5). It is therefore believed that the phase difference, as discussed above, is the major origin for the different currents measured in the pit and matrix regions.

Besides the different conductivities, the BM-SFO and PV- SFO phases should also exhibit different electrochemical strain responses because of the different oxygen stoichiometries. The electrochemical strain microscopy (ESM) technique was thus employed to detect the electrochemical strain arising from ionic redistribution and surface electrochemical reaction under an applied tip bias⁵⁵⁻⁵⁸ for the mixed-phase SFO film. To minimize influences from the topographic crosstalk on the ESM signals, the dual AC resonance tracking (DART) mode, which can track the resonant frequency, was used.⁵⁸Figure 3c− e presents the topography of a 5×5μm²area and the ESM amplitude and phase images taken over this area, respectively.

It is not unexpected that both the ESM amplitude and phase images show almost one-to-one correspondence with the topography. The ESM amplitude in the pits is higher than that in the matrix, implying that the PV-SFO phase exhibits a larger electrochemical strain response than the BM-SFO phase. This may be because the PV-SFO phase has either higher oxygen diffusivity or higher activity of oxygen incorporation/extraction near the surface.^59,60

While both the conductance and electrochemical strain response are strongly correlated to the phases in the mixed- phase SFOfilm, an interesting question remains how the phase

along with conductance and electrochemical strain response evolves upon applications of electrical and mechanical stimuli.

We therefore performed further AFM-based characterizations.

As can be seen fromFigure 4a (also seeFigure 1f), electrical writing with +7 V leads to the cracked-earth-type morphology and a noticeable increase in surface height. Accompanied with the topographic change, a decrease in current is also observed upon +7 V writing (see Figure 4d). The changes in both surface height and conductance can be reversed by an electrical writing with opposite polarity. As shown in Figure 4b,c, the surface height in the inner−7 V written area is decreased by

∼0.8 nm, compared with that in the outer +7 V written area (note that only the surface height decrease is observes, while the nanoscaffold-type morphology is not recovered; also see Figure 4g). Meanwhile, the current is increased from ∼3 to

∼35 pA, corresponding to an on/offconductance ratio of∼10 (Figure 4e,f). To exclude the transient current contributions from the charging/discharging effect, the time-dependent current measurements were conducted. As shown in Figure S5, the different currents read in the +7 and−7 V written areas are rather stable against time, confirming that the conductiv- ities of these two areas are truly different. Because the BM-SFO phase has a largerd_oopand is less conductive compared with the PV-SFO phase, the above results therefore suggest that a reversible phase transformation from PV-SFO to BM-SFO (BM-SFO to PV-SFO) may occur upon +7 V (−7 V) writing.

To further characterize the resistive switching phenomenon, local current−voltage (I−V) measurements were performed.

Figure 5a illustrates typical I−V characteristics measured by locating the tip in the pit and matrix regions and sweeping the voltage as 0 → −6 V → +6 V → 0. Much smaller I−V hysteresis is observed in the pit region than in the matrix region as the voltage goes along 0→−6 V→0, indicating that the BM-SFO→ PV-SFO phase transformation occurs, while the PV-SFO phase remains unchanged during the application of a negative voltage. After the voltage sweeps from−6 V back to 0, the phases in both pit and matrix regions may become the same PV-SFO phase; therefore, both regions exhibit similar hysteresis as the voltage sweeps along 0→+6 V→0. These localI−Vhysteresis loops are well consistent with the above C- AFM mapping results, confirming that the topotactic phase transformation between PV-SFO and BM-SFO can be driven by the electricfield and it can cause resistive switching.

The topotactic phase transformation also triggers a significant change in electrochemical strain response. As shown in Figure 4g−i, +7 V (−7 V) writing leads to a decrease (increase) in ESM amplitude and noticeable changes in the phase contrast, concurrently with the topographic changes. This observation is consistent with the scenario of topotactic phase transformation between BM-SFO and PV- SFO because the PV-SFO phase is shown to have larger electrochemical strain response than the BM-SFO phase (see Figure 3d).

Besides the ESM mappings, local ESM amplitude loops were also recorded. Figure 5b shows the comparison of ESM amplitude hysteresis loops measured in the pit and matrix regions by applying a voltage cycle of 0→+6 V→−6 V→0.

The matrix region shows negligible hysteresis, while the pit region exhibits a butterfly winglike hysteresis as the voltage sweeps from 0 to +6 V and back to 0, suggesting that a positive voltage induces the phase transformation from PV-SFO to BM-SFO while keeping the BM-SFO phase unchanged. After the voltage sweeps from +6 V back to 0, the phases in both pit

and matrix regions may become the same BM-SFO phase, giving rise to similar hysteresis during the voltage sweep of 0

→−6 V→0. The above ESM results therefore demonstrate that the topotactic phase transformation can lead to a change in the electrochemical strain response.

To further study the effect of electricfield on the topotactic phase transformation, the topographic changes as a function of both magnitude and duration time of the write voltage were monitored.Figure 5c,d shows that the surface height becomes gradually larger as the write voltage increases from +4 to +9 V.

Meanwhile, as the scan rate decreases from 3.5 to 1 Hz, the surface height also gradually increases (seeFigure 5e,f). Note that a lower scan rate means a longer tip dwelling time on each point in the scanned area. These results indicate that a larger magnitude and longer duration time of a positive voltage can drive a higher degree of PV-SFO → BM-SFO phase transformation. Similar dependences of phase transformation on write voltage magnitude and duration time can be found in the writing process with negative voltages (seeFigure S6).

Notably, the surface height change (Δh) can be retained for a prolonged period of time, e.g., 40 h, after removing the electricfield (see Figure S7), attesting to the nonvolatility of the phase transformation and associated electrochemical strain.

If the strain is defined asΔh/d, where dis thefilm thickness, theΔhof 800 pm observed at +9 V with a scan rate of 1 Hz corresponds to a strain of ∼1.3%. This huge electrochemical strain is on the same order of magnitude as the highest strains achieved in state-of-the-art relaxor ferroelectrics,¹ which may have potential applications in actuators.

We now focus on the possible mechanism underlying the electric-field-induced topotactic phase transformation. Accord- ing to previous studies, a chemical equilibrium is generally reached near the SFOfilm surface¹⁶

F + e′ +

O₂ 4 2V_O^.. 2O_O^x (1)

By applying a negative bias on the film surface, the forward reaction ofeq 1 occurs. The oxygen atoms in the ambient are thus incorporated into thefilm,filling the vacancy sites in BM- SFO. Consequently, the BM-SFO phase transforms into the PV-SFO phase. The region with the PV-SFO phase then gradually grows as the oxygen ions migrate toward SRO under the negative bias. On the contrary, applying a positive bias on the film surface triggers the reverse reaction of eq 1. The oxygen atoms in the lattice are pulled out, creating oxygen vacancies. This leads to the conversion of the PV-SFO phase into the BM-SFO phase.

The above mechanism is different from that proposed by Ferreiro-Vila et al.⁶¹They proposed that the oxygen vacancies migrate and accumulate underneath the tip when applying a negative tip bias, leading to the PV-SFO → BM-SFO phase transformation. However, since the reverse phase trans- formation (BM-SFO → PV-SFO) induced by the negative bias was observed in our sample, we think that the oxygen incorporation from the ambient into the SFOfilm [seeeq 1]

may play the dominant role. Indeed, both oxygen migration and oxygen incorporation/extraction may co-work and which one dominates depends on the sample and the electricfield distribution. Their and our samples (both in the pristine state) mainly contain the PV-SFO phase and the mixed phases, respectively. In addition, no bottom electrode was used in their sample, while we used SRO as the bottom electrode, which should lead to different electricfield distributions. Considering these differences, we think that both their and our observations are explainable, and the oxygen incorporation/extraction is a more suitable mechanism for our sample. To further verify it, checking whether the O₂ in the ambient is an indispensable medium for the BM-SFO→PV-SFO phase transformation is of necessity. Hence, the topographic changes in the SFOfilm after applying −7 V in different atmospheres were examined.

Figure S8 shows that the topographic change becomes less observable as the volume fraction of O₂decreases, confirming the critical role played by O₂ in the topotactic phase transformation and also supporting our proposed mechanism.

Mechanical stimuli were also applied to trigger the topotactic phase transformation and property changes in the mixed-phase SFOfilm. A region of 10×10 μm²was initially written with +7 V. Consequently, this region contains mainly BM-SFO and has a relatively large surface height. This pretreatment was used because it can result in a larger contrast between the regions with and without mechanical writing in the topography image. Then, three small regions of 3×3μm² inside the pretreated region were scanned with loading forces of∼9,∼22, and∼40μN, respectively, while applying zero tip bias. Figure 6a shows the topography image taken after the mechanical writing. As the loading force increases from∼9 to

∼40 μN, the sink of the film surface emerges and becomes Figure 5.Hysteretic responses of current and electrochemical strain

and dependences of topographic change on write voltage magnitude and duration time. (a) HystereticI−Vcharacteristics measured in pit and matrix regions. (b) ESM amplitude hysteresis loops measured in pit and matrix regions. The red arrows indicate the starting step of a voltage cycle, while the black arrows indicate the following steps of this voltage cycle. Insets show the locations where hysteresis loops were measured. AFM topography images obtained after the positive writing with (c) different voltages (scan rate 1 Hz) and (e) different scan rates (voltage +7 V). Surface height changes as a function of (d) voltage and (f) scan rate. The data in (d) and (f) are the statistical results obtained from (c) and (e), respectively.

increasingly noticeable. Along with the topographic change, the conductance also becomes higher with an increase in loading force (seeFigure 6c). The simultaneous changes in topography and conductance suggest that the BM-SFO→PV-SFO phase transformation may occur during mechanical writing. After rewriting the whole 10×10 μm²region with +7 V, both the surface height and conductance of the mechanically written regions are largely recovered to be similar to those of the surrounding regions (seeFigure 6b,d). The recovery of both the surface height and conductance further confirms the occurrence of topotactic phase transformation, rather than plastic deformation or wear, induced by mechanical writing.

In terms of the underlying mechanism, one can also refer to eq 1. The mechanical pressure applied by a tip can lower the formation energy of lattice oxygen,^62,63promoting the forward reaction of eq 1. In addition, the pressure may push oxygen vacancies away from the region underneath the tip.⁶⁴ Both effects may contribute to the occurrence of BM-SFO →PV- SFO phase transformation under mechanical pressure.

However, since the tip cannot apply tensile forces to the sample, the PV-SFO→BM-SFO phase transformation is not available.

Up to this point, it has been demonstrated that the conductance and electrochemical strain response are strongly coupled with the phase in the mixed-phase SFO film, and reversible phase transformation between PV-SFO and BM- SFO can be driven by the electricfield. The mixed-phase SFO film with well-ordered nanosized PV-SFO and BM-SFO domains may therefore work as a high-density dual-functional memory device that can store the information of both conductance and electrochemical strain response. Besides, the giant electrochemical strain of 1.3% makes the mixed-phase SFO film potentially useful for actuator applications. In addition, the topotactic phase transformation in the mixed- phase SFOfilm can also be induced by the mechanicalfield, making it attractive for pressure sensing. One last thing worth mentioning is that the AFM technique is a facile tool to study

the topotactic phase transformation. As demonstrated here and also in Ferreiro-Vila’s work,⁶¹ an AFM tip with mild applied electrical and mechanical stimuli (several volts and several micronewtons, respectively) can readily trigger the topotactic phase transformation, unlike thermal annealing,^13,31 ionic liquid gating,^14,15,32 and plasma-assisted ion implantation,⁶⁵ which require high-temperature heating, long-term electric poling, and introducing foreign ions, respectively.

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CONCLUSIONS

In summary, SFO epitaxial thinfilms with a nanoscale mixture of BM-SFO and PV-SFO phases have been fabricated on the SRO-buffered STO substrates byfinely controlling the oxygen pressure during the film growth. The mixed-phase SFO film (pristine state) exhibits a nanoscaffold structure, where PV- SFO nanodomains are embedded in the BM-SFO matrix.

Coupled with this novel structure, gridlike patterns are formed in both current and electrochemical strain maps. Applying a positive (negative) tip bias can drive a phase transformation from PV-SFO to BM-SFO (BM-SFO to PV-SFO), resulting in an increase (decrease) in surface height as well as a decrease (increase) in conductance and ESM amplitude. Particularly, an on/offconductance ratio of∼10 is obtained upon applying±7 V, and an electrochemical strain as large as∼1.3% is achieved at +9 V with a scan rate of 1 Hz. In addition, a mechanical pressure exerted by the tip is also able to trigger the BM-SFO

→ PV-SFO phase transformation, which manifests as a decrease in surface height and an increase in conductance.

The ability to exhibit significant changes in both conductance and electrochemical strain in response to electrical and mechanical stimuli promises the mixed-phase SFO films a wide range of applications such as memories, actuators, and sensors.

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Film Deposition.^METHODS Mixed-phase SFO thin films were epitaxially grown on SRO-buffered (001)-oriented STO substrates using pulsed laser deposition (PLD) with a KrF excimer laser (λ= 248 nm). SRO bottom electrode layers (∼30 nm) werefirst deposited at a substrate temperature of 680 °C in an oxygen atmosphere of 15 Pa.

Subsequently, ∼60 nm mixed-phase SFO films were deposited at 660°C under apO₂of 3 Pa. For comparison, the SFOfilms were also grown underpO₂of 1, 2, 5, and 10 Pa. During deposition, the energy fluence of the laser wasfixed at 1.2 J/cm²with a repetition rate of 4 Hz. After deposition, the samples were cooled to room temperature at 10°C/min in an oxygen atmosphere the same as that used for SFO deposition.

Structural Characterizations. The crystalline phases were examined using XRD and RSM (X’Pert PRO, PANalytical). STEM was conducted on a Jeol ARM200F microscope equipped with a cold field-emission gun and ASCOR probe corrector. The STEM plan- view specimens were prepared by grinding/polishing and ion milling from the substrate side.

AFM-Based Characterizations. Nanoscale observation and manipulation of phases and physical properties were conducted using an integrated AFM system (Asylum Cypher, Oxford Instru- ments) with conductive Pt-coated silicon tips (EFM, Nanoworld).

More specifically, AFM, C-AFM, and ESM were used to characterize the surface morphology, electrical conduction, and electrochemical strain response, respectively. When performing C-AFM, a read voltage of +3 V was used for current mappings. In the ESM measurements, the dual AC resonance tracking (DART) mode with an AC drive amplitude of 1.5 V was employed. Unless otherwise specified, all of the AFM-based characterizations were performed at ambient conditions and the voltage is defined to be positive when the Pt tip is positively biased.

Figure 6. Mechanical-stress-induced phase transformation. (a, c) AFM topography image and C-AFM current map, respectively, taken after writing three 3×3μm²regions with different mechanical forces.

Prior to the mechanical writing, the whole 10×10μm²region was electrically written with +7 V. (b, d) AFM topography image and C- AFM current map, respectively, taken after rewriting the whole 10× 10μm²region with +7 V.

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ASSOCIATED CONTENT

*^sı Supporting Information

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsami.0c03684.

AFM and XRD results of SFOfilms grown on different pO₂; detailed analyses on the atomic structure and strain distribution near the BM/PV interface; time-dependent current measurements; surface height change induced by negative writing; retention of phase transformation;

electrical writing in different atmospheres; and origin for the observation of the cracked-earth-type morphology in STEM (PDF)

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AUTHOR INFORMATION Corresponding Authors

Zhen Fan−Institute for Advanced Materials, South China Academy of Advanced Optoelectronics and Guangdong Provincial Key Laboratory of Optical Information Materials and Technology, South China Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China; orcid.org/0000-0002-1756-641X;

Email:[email protected]

Haijun Wu−Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 Singapore;

State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China; orcid.org/

0000-0002-7303-379X; Email:[email protected] Authors

Junjiang Tian− Institute for Advanced Materials, South China Academy of Advanced Optoelectronics and Guangdong Provincial Key Laboratory of Optical Information Materials and Technology, South China Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China

Yang Zhang−Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 Singapore;

Instrumental Analysis Center of Xi’an Jiaotong University, Xi’an Jiaotong University, Xi’an 710049, China

Lei Zhao−Institute for Advanced Materials, South China Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China

Jingjing Rao− Institute for Advanced Materials, South China Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China

Zuhuang Chen−School of Materials Science and Engineering, Harbin Institute of Technology, Shenzhen 518055, China;

orcid.org/0000-0003-1912-6490

Haizhong Guo−School of Physical Engineering, Zhengzhou University, Zhengzhou 450001, China; orcid.org/0000- 0002-6128-4225

Xubing Lu−Institute for Advanced Materials, South China Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China; orcid.org/0000-0002-2552- Guofu Zhou9571 −Guangdong Provincial Key Laboratory of Optical Information Materials and Technology, South China Academy of Advanced Optoelectronics and National Center for International Research on Green Optoelectronics, South China Normal University, Guangzhou 510006, China; orcid.org/

0000-0003-1101-1947

Stephen J. Pennycook−Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 Singapore; orcid.org/0000-0002-3210-6323 Xingsen Gao−Institute for Advanced Materials, South China

Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China; orcid.org/0000-0002-2725- Jun-Ming Liu0785 − Institute for Advanced Materials, South China

Academy of Advanced Optoelectronics, South China Normal University, Guangzhou, China; Laboratory of Solid State Microstructures and Innovation Center of Advanced

Microstructures, Nanjing University, Nanjing 210093, China;

orcid.org/0000-0001-8988-8429 Complete contact information is available at:

https://pubs.acs.org/10.1021/acsami.0c03684 Author Contributions

¶J.T. and Y.Z. contributed equally to this work.

Notes

The authors declare no competingfinancial interest.

■

ACKNOWLEDGMENTS

The authors would like to thank the National Key Research P r o g r a m of C h i na ( N o s. 2 0 16 Y F A 02 0 10 0 2 a n d 2016YFA300101), the State Key Program for Basic Researches of China (No. 2015CB921202), the National Natural Science Foundation of China (Nos. U1932125, 11674108, 51431006, and 11574365), the Natural Science Foundation of Guang- d o n g P r o v i n c e ( N o s . 2 0 1 6 A 0 3 0 3 0 8 0 1 9 a n d 2020A1515010996), the Guangdong Provincial Key Labo- ratory of Optical Information Materials and Technology (No.

2017B030301007), MOE International Laboratory for Optical Information Technologies, and the 111 Project. X.L. and Z.F.

acknowledge the Project for Guangdong Province Universities and Colleges Pearl River Scholar Funded Scheme 2016 and 2018, respectively. H.W. would like to acknowledge the financial support from Singapore Ministry of Education Tier 1 grant, R-284-000-212-114, for the Lee Kuan Yew Postdoctoral Fellowship. S.J.P and H.W. would like to acknowledge the financial support from the Ministry of Education, Singapore under its Tier 2 grant (No. MOE2017-T2-1-129). Z.F. would like to thank Prof. P. Yu for the valuable discussion.

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