Nanoscale Control of Oxygen Defects and MetalâInsulator Transition

Subscriber access provided by UNIVERSITY OF TOLEDO LIBRARIES

Article

Nanoscale Control of Oxygen Defects and MetalInsulator Transition in Epitaxial Vanadium Dioxides Yogesh Sharma, Janakiraman Balachandran, Changhee Sohn, Jaron T Krogel, Panchapakesan Ganesh, Liam Collins, Anton V. Ievlev, Qian Li, Xiang Gao, Nina Balke, Olga S. Ovchinnikova, Sergei V. Kalinin, Olle Heinonen, and Ho Nyung Lee ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b03031 • Publication Date (Web): 15 Jun 2018 Downloaded from http://pubs.acs.org on June 15, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

Nanoscale Control of Oxygen Defects and MetalInsulator Transition in Epitaxial Vanadium Dioxides Yogesh Sharma†, Janakiraman Balachandran‡, Changhee Sohn†, Jaron T. Krogel†, Panchapakesan Ganesh‡, Liam Collins‡, Anton V. Ievlev‡, Qian Li‡, Xiang Gao†, Nina Balke‡, Olga S. Ovchinnikova‡, Sergei V. Kalinin‡, Olle Heinonen§, Ho Nyung Lee† *

†

Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA ‡

Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA §

Materials Science Division, Argonne National Laboratory, Lemont, IL 60439, USA

*

[email protected]

1 ACS Paragon Plus Environment

ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ABSTRACT Strongly correlated vanadium dioxide (VO2) is one of the most promising materials to exhibit temperature-driven, metal-insulator transition (MIT) near room temperature. The ability to manipulate the MIT at nanoscale offers both an insight to understand the energetics of phase transition and a promising potential for nanoelectronic devices. In this work, we study nanoscale electrical modifications of the MIT in epitaxial VO2 thin films using a combined approach with scanning probe microscopy (SPM) and theoretical calculations. We find that applying electric voltages of different polarity through an SPM tip locally changes the contact potential difference and conductivity on the surface of VO2 via modulating the oxygen stoichiometry. We observed nearly two orders of magnitude change in resistance between positive and negative biased-tip written areas of the film, demonstrating the electric field modulated MIT behavior at the nanoscale. Density functional theory calculations, benchmarked against more accurate manybody quantum Monte-Carlo calculations, provide information on the formation energetics of oxygen defects that can be further manipulated by strain. This study highlights the crucial role of oxygen vacancies in controlling the MIT in epitaxial VO2 thin films, useful for developing advanced electronic and iontronic devices.

KEYWORDS: vanadium dioxide, metal‒insulator transition, scanning probe microscopy, oxygen vacancy, density functional theory, quantum Monte-Carlo 2 ACS Paragon Plus Environment

Page 2 of 28

Page 3 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

Metal-insulator transition (MIT) in strongly correlated materials such as vanadium dioxide (VO2) has attracted considerable interest for various technological applications.1-4 This thermally induced MIT in VO2 is accompanied by a structural phase transition from a lowtemperature insulating monoclinic phase to a high-temperature metallic rutile phase.5 VO2 can be used in many practical applications based on the drastic physical property changes associated with the phase transition. Recently, a plethora of research has been presented to effectively control the MIT temperature and physical properties of VO2 by means of chemical strain,6-8 epitaxial strain,9-11 electrochemical modulation,12-15 optical excitation,16 hydrogenation,17 and electric field effect,18 to name a few. Even though it is known that oxygen vacancy formation, strain, and electric field couple to one another, we still lack a comprehensive understanding of how this coupling allows us to control MIT. Thin films form perfect model systems for such explorations. In epitaxial VO2 thin films, interfacial strain between the film and the substrate has been reported to reduce the MIT temperature to close to room temperature.8,19 However, the epitaxial strain makes the structural phase transition more complicated because of the additional mechanical deformation resulting from the straining. On the other hand, electric-field tuning of MIT in VO2 films was shown using electrolyte gating experiments.18 Electrolyte gating was used to vary the electron density in VO2 by means of electrostatic effects at the film-electrolyte interface, resulting in a significant reduction of MIT temperatures. Later, it was realized that the electric field at the interface is so large that it leads to electrochemical changes related to the removal of oxygen from certain lattice orientations in the VO2 film.12 These observations have indicated that the MIT behavior of VO2 is very sensitive to the oxygen nonstoichiometry and strain. Therefore, MIT can be accessed by introducing oxygen vacancies and changing the strain


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

state in VO2. There have not been many reports on oxygen vacancy, local strain, and electric field-induced nanoscale or mesoscale changes in the electronic properties of epitaxial VO2 thin films. Scanning probe microscopy (SPM) is a promising technique for the study of local electronic and mechanical properties of various oxide materials on the nanoscale under various external stimuli, such as applied electric bias, mechanical force, and temperature.20-25 In this paper, we present a combined experimental and theoretical approach to elucidating the role of oxygen vacancies in controlling the MIT behavior of epitaxial VO2 thin film. Using Kelvin probe force microscopy (KPFM) and conducting atomic force microscopy (CAFM), we establish that the applied electric voltage of different polarity through the SPM tip can locally control the VO2 film’s surface potential and conductivity and thereby modulate the MIT behavior. We observed two orders of magnitude enhanced current in a negative biased-tip written area compared with a positive biased-tip written area. Tip-bias polarity dependence of the contact potential difference (CPD) and the observed conductance change indicate the dominant role of oxygen-vacancy– assisted electrochemical reactions at the film surface. The changes in oxygen stoichiometry and their strain dependence were quantified using density functional theory (DFT) calculations. We attribute our results to the electric-field–induced local change in the oxygen stoichiometry in VO2 films, which can provide a basis for nanoscale electrochemical control of MIT in VO2 thin films. RESULTS AND DISCUSSION We grew epitaxial VO2 films with a thickness of 16 nm on (001) TiO2 and Nb-doped TiO2 (Nb:TiO2) single-crystal substrates using oxide pulsed laser epitaxy. Figure 1(a) shows an X-ray diffraction (XRD) θ‒2θ pattern of a VO2/Nb:TiO2 thin film, confirming the growth of 4 ACS Paragon Plus Environment

Page 4 of 28

Page 5 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

highly strained (402) oriented monoclinic VO2 film. The atomic-level scanning transmission electron microscopy (STEM) observations confirmed the very high crystallinity of the films, as shown in the inset of Fig. 1(a). The MIT characteristics of a VO2 film grown on TiO2 substrate were investigated by dc transport measurements as a function of temperature across the phase transition regime. The temperature-dependent resistivity [ρ(T)] curve of the VO2 film is depicted in Fig. 1(b). We obtained the MIT temperature using the derivative of the ρ(T) curve, as shown in the inset of Fig. 1(b). Based on the maxima in the 1/ρ(dρ/dT) curve and considering the rather substantial hysteresis, the MIT temperature was found to be ~314 ± 10 K. The lower value of MIT temperature compared with bulk VO2 is attributed to the tensile strain in the VO2 film.26 The nanoscale electronic phase transition across the MIT temperature of a VO2 film was also studied using CAFM. CAFM images at various temperatures across the MIT are shown in Fig. 1(c). At 288 K, below the MIT, we observe nearly homogeneous low-current values, indicating the insulating phase of the film. As the sample is heated, the nanoscale domains or clusters with a higher current contrast compared with that of the insulating host start to appear, indicating an onset of the electronic phase transition to a metallic state. Further increase in temperature shows that these metallic domains grow gradually. The local current-voltage curves were also measured at low (292 K), intermediate (303 K), and high (330 K) temperatures to compare the conductivity of the nanoscale clusters in the insulator-to-metal transition regime (supporting information in Figure S2). Such observation of the nanoscale metallic clusters and/or metallic puddles that appear at the onset of the insulator-to-metal transition in VO2 have been proposed previously.27,28 We observed that a complete insulating to metallic phase transition can be achieved above 318 K in agreement with the ρ(T) curve (Fig. 1(b)).


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

To study the effect of local electric field on the electrochemical behavior of the film and the role of oxygen vacancies on electronic phase transition, a combination of CAFM and KPFM techniques was employed. Micrometer-sized regions on VO2 film were electrically written at 298 K, below the insulator-to-metal transition temperature, with a conducting dc-biased tip in contact mode atomic force microscopy (AFM). The ±4 V tip bias-induced evolution in CPD in the written areas was measured by KPFM, as shown in Fig. 2(a). In KPFM, the CPD between the tip and the film surface is defined as the offset in the respective vacuum energy levels (in volts).24,29 There was no surface damage or morphology change observed in that area of the film after electric writing, as confirmed by topography image shown in Fig. 2(b). Similarly, the resistance states of the same areas were measured by CAFM and are shown in Fig. 2(c). Both tip-bias polarity dependence of CPD and observed conductance change indicate the dominant role of ionically mediated electrochemical interaction over the charge injection at the film surface during the bias application.22,29 Furthermore, to confirm the effect of the polarity of applied bias voltage on CPD and to check the reversibility of the local electrochemical effect, a KPFM image was recorded after electrical writing using a concentric square method, as shown in Fig. 2(d). The voltage polarity-dependent switching of CPD and CAFM current further confirmed the reversibility of the local electrochemical process at the VO2 film surface (supporting information in Figure S3). Focusing on the electrochemical mechanism of this local electric-field-driven process, we argue that the positive bias-tip repelled the preexisting positively charged oxygen vacancies from the film surface (see Fig. 2(e)). Along with that, negatively charged oxygen ions and/or electrons are attracted to the film surface under this positive bias. On the other hand, a negatively biased-tip repelled oxygen ions and generates more oxygen vacancies with the removal of oxygen, making the VO2 film surface more conducting, as


Page 6 of 28

Page 7 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

confirmed by the CAFM image (Fig. 2(c)). Moreover, based on the temperature-dependent KPFM results (supporting information in Figure S4), we can predict that the creation of oxygen vacancies (negative tip-bias written area) can lower the phase transition temperature, while eliminating the oxygen vacancies (positive tip-bias written area) can enhance the phase transition temperature of VO2. Such a prominent role of oxygen vacancies controlling the phase transition has been recently reported by Zhang et al.30 using systematic postgrowth oxygen annealing of VO2 films. We have estimated the critical electric field required to generate oxygen vacancies in VO2 thin films. A series of negative tip biases were applied to the VO2 film surface for the electrical writing. The seven rectangular regions were patterned by tip bias ranging from –0.4 to –2.8 V with a 0.4 V incremental step at two different temperatures of 300 and 310 K. After electrical writing with incremental tip bias, KPFM images were acquired to visualize the resulting CPD change at each temperature on the VO2 film surface, as shown in Fig. 3(a-b). We observed that the CPD value decreases as the applied negative bias increases, while at a bias voltage of –2.4 V (at 300 K) CPD becomes nearly saturated (Fig. 3(c)). Our results agree with previous reports where CPD was found to be lower in oxygen-deficient VO2 films and became nearly saturated after the full metallization of the film above insulator-to-metal transition temperature.21,31 The saturation of CPD suggests a full charge compensation through oxygen vacancy formation at –2.4 V, which also indicates that a certain bias voltage is required to create oxygen vacancies in VO2. However, during the KPFM measurements, the influence of charge injection from the AFM tip and dissociated species of water can influence the measured CPD, obscuring the true surface potential change that originated because of the creation of oxygen vacancies.32


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Moreover, the SPM measurements alone cannot unambiguously detect and identify the precise chemical processes under tip at surface and even in the bulk of the sample. To provide more experimental support of our findings of creation and annihilation of oxygen vacancies using biased tip, we have carried out investigations combining AFM and time-of-flight secondary ion mass spectroscopy (ToF-SIMS).33,34 We employed ToF-SIMS to elucidate surface and bulk modification in chemical composition of VO2 thin films induced by the tip bias. The background pressure in the chamber during measurements was on the order of 10-8 Torr. Initially, we have electrically written two 5 × 5 µm2 regions on the film using ±4 V tip bias, as shown in Fig. 4(a). We have not seen any surface topographic changes after electrical writing. The ToF-SIMS experiments were optimized for the detection of the positive secondary ions around these regions, which allowed detection of the base component V+ and VO+ ions of the film. We observed changes in the concentration of V+ and VO+ ions in the +4 and –4 V written areas, as shown in Figs. 4(a) and (b). The V+ and VO+ ion concentrations were found to be significantly lower in the region written by –4 V, which indicates a reduction of the film in that area, confirming the presence of oxygen vacancies. The variation in the concentration of ions can be further evaluated by line profile (Fig.4 (c)) and depth profile analyses (Fig. 4(d)). It is worth pointing out that we can see a slight decrease in the V+ and VO+ ion concentrations even in the positively written area. However, changes in this case are much insignificant and, thus, could be attributed to surface contamination, for example, by silicone ions from the AFM tip.35 In addition, we have also employed combined ToF-SIMS and AFM studies to investigate the role of surface protonation or the potential hydrogenation effect on enhanced electrical conductivity of the VO2 film. We have not observed any significant changes in the concentration of H+ ions, compared with VO+ ions, in pristine and electrically written areas (supporting information in


Page 8 of 28

Page 9 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

Figure S5), indicating that hydrogenation cannot be a prominent factor for enhanced conductivity observed in our sample. We performed DFT calculations to understand the role of strain on modulating oxygenvacancy–induced electron doping of VO2 thin films. We used the DFT+U scheme36 with the PBE-flavor exchange-correlation functional37 to capture correlation effects, as well as—at least to some extent—the orbital dependence of the full potential.38-40 The value of “U” was determined to be “U = Uopt = 4 eV”, by optimizing the total energy of the solid, coming from a many-body quantum Monte-Carlo (QMC) scheme (supporting information in Figure S6). The DFT+U wavefunction has an occupancy of V-3d of ~1 e-. This is both interesting and important because a localization of 1e- can be thought of as being the driving force behind a V–V dimerization in the low-temperature M1 phase. The calculated bandgap of the M1 phase increases from 0.007 eV at U = 0 to 0.83 eV at U = Uopt, close to the experimentally measured optical gap (0.6 eV).41,42 We incorporated a single oxygen vacancy in a 48-atom supercell of VO2 in the M1 phase, corresponding to a vacancy concentration of 3.125%. For a neutral vacancy, the standard semilocal PBE functional gives rise to a fully delocalized charge density, as shown in Fig. 5(a). In contrast, PBE+Uopt gives rise to a mixed localized and delocalized picture represented in Fig. 5(b). In this case, electrons have a tendency to localize on the two near-neighbor V–d orbitals that formed a dimer before the vacancy was introduced, although there is also some delocalized charge density. This result suggests that in the presence of strong electron localization effects, an oxygen vacancy does not act as a double donor. Such many-body effects were originally shown in SrTiO343,44 and have been recently shown by Spaldin et al.45 in other


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 28

correlated oxides. The exact energy level of these localized states is influenced by vacancy concentration.44,46 To further understand how strain can allow us to control the MIT temperature due to electron-doping (e-doping) from oxygen-vacancies, we constructed the defect phase diagram for oxygen vacancies in the M1 phase using PBE+Uopt, both in the unstrained and epitaxially strained conditions, closely resembling our experiments. The vacancy formation energy of a charge q vacancy is given as = [( ); ] + + − [( )] + !! .

The reference energies are total energies for bulk VO2, [( )], and O2 dimer, , with the formation energy of the O2 dimer corrected to give experimental enthalpy of formation (i.e., 5.15 eV); is the electronic chemical potential (Fermi energy), which can generally vary from the valence band minimum to conduction band maximum (CBM). The corresponding formation energy calculations hence apply to oxygen-rich growth conditions, which are what we expect in experiments since the measurements on the films are made under ambient conditions. The correction term ("#$$ ), predominantly due to image charge interactions, is a constant value47 and does not affect differences in formation energies without and with strain; it is therefore neglected here. The oxygen-vacancy formation energy in unstrained bulk VO2 M1 phase for the two symmetrically inequivalent oxygen sites [O(I) and O(II)] for two different charge states (q = 0, +2) is shown in Fig. 5(c). The defect transition level of the O(I) site is about 0.12 eV below the CBM, while the transition level for the O(II) site is much deeper, about 0.30 eV below the CBM. We incorporated the influence of epitaxial strain in our models considering (402)-oriented epitaxial VO2(M1) films on (001) TiO2 substrates. The corresponding strain values along the b and c crystallographic axes are identified to be 1.21% and 1.31%, respectively, which are similar 10 ACS Paragon Plus Environment

Page 11 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

to the experimental strain values. The defect phase diagram corresponding to the strained model is shown in Fig. 5(d). The bandgap increases from 0.83 eV unstrained to 0.87 eV under strain. The defect transition levels for both oxygen sites move closer to each another and to the CBM, with values of 0.07 and 0.09 eV, respectively, below the CBM. Most importantly, our calculations suggest that under epitaxial strain, oxygen vacancies act as shallow donors, compared with unstrained VO2, thereby supporting the hypothesis that the strain-induced tunability of the insulator-to-metal transition originates from an increased ease of e-doping of the material by oxygen-vacancy formation. This in turn can explain the lowering of the MIT temperature observed experimentally under strain. Based on this finding, and the now proven hypothesis that control over MIT via strain is related to an increased ease of e-doping by oxygen vacancies, we try to relate our computed defect phase diagram to the KPFM experiments. For a surface in vacuum, the change in the CPD with increasing negative bias is related to the change in its work function, which in turn can be attributed to the shift in Fermi energy with respect to the band edges.48,49 It can be observed from Fig. 3(c) that at 300 K, on average, an increase in the magnitude of an electrical writing bias of 0.4 V changes the CPD by 0.08 eV, pushing it more towards the metallic value. A corresponding increase in Fermi energy in the defect phase diagram by 0.08 eV towards the CBM (i.e., metallicity) results in an increase in the formation energy of the q = +2 oxygen vacancy by about 0.15 eV. The shift seen from our defect phase diagram should be the lower bound for the applied bias to cause the relative Fermi-level shift and is close to the 0.4 V bias seen in experiments. The difference in the computational and experimental values is likely because the correlation between CPD and work function is only approximate for semiconductors, which can contain surface charges, especially when measurements are made in ambient conditions and not vacuum, as in


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

our experiments. Moreover, oxide surfaces are different from conventional semiconductors or nonoxide insulators, and a single work-function definition might not hold anymore.50,51 A more detailed study is needed to quantitatively predict the coupling between external bias and vacancy-induced doping in a strongly correlated material. CONCLUSIONS In summary, we studied electrochemical modification of the local transport properties of epitaxial VO2 thin films using nanoscale KPFM and CAFM measurements. We identified that the applied electric voltage through an AFM tip could locally modulate the CPD and conductance at the VO2 film surface, mediated by the creation and annihilation of oxygen vacancies. The observed electrochemical behavior in a VO2 film allows for the electrical manipulation of local resistivity and thus the MIT behavior down to the nanometer level. DFTbased calculations on unstrained VO2 show that oxygen vacancies do not act as double donors in VO2, but this electron doping through creation of oxygen vacancies becomes energetically less costly under epitaxial strain. Our results show the strong coupling among oxygen-vacancy creation and local strain and electric fields, thereby providing insight into the mechanism of phase transition and nanoscale control of MIT for the development of next-generation electronic devices.

METHODS Epitaxial film growth: We deposited VO2 epitaxial films on single-crystal TiO2 and Nb-doped (001) TiO2 substrates by pulsed laser epitaxy. We ablated a sintered VO2 target of M1 phase by a KrF excimer laser (248 nm in wavelength) at a laser fluence of 1.2 Jcm−2 and at a laser repetition 12 ACS Paragon Plus Environment

Page 12 of 28

Page 13 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

rate of 5 Hz. The substrate temperature and oxygen partial pressure were set to 300°C and 12 mTorr, respectively. After deposition, the samples were cooled to room temperature at higher oxygen pressure, ~100 Torr to avoid oxygen deficiencies.52 Measurements: The structural characteristics of VO2/Nb:TiO2 (001) thin films were analyzed by a four-circle, high-resolution X-ray diffractometer (X’Pert Pro, Panalytical) using Cu-Kα1 radiation. The electrical transport of VO2 films grown on TiO2 substrates was measured by a standard four-point probe method with temperatures varied from 270 to 390 K with a physical property measurement system (Quantum Design). Ambient and temperature-dependent SPM studies were performed with a commercial AFM system (Asylum Research Cypher) equipped with a variable-bandwidth current amplifier (FEMTO DLPCA-200). CAFM and KPFM measurements were carried out using a conducting PtSi-coated tip (PtSi-FM-20, Nanosensors). The spring constants of each cantilever were obtained from force‒distance curve and thermal tuning methods. The samples were heated to 100°C for 15 min to remove potential surface contaminants and then cooled to room temperature before electrical writing. The atmosphere in our AFM hood was well controlled in terms of temperature and humidity. The relative humidity value was measured at ~36% inside the AFM hood. For electrical writing, the AFM operated in contact mode and was applied with a tip-deflection set point of 0.6 V and a scan frequency of 0.7 Hz. The spring constant of the PtSi tip was 0.9 N/m. KPFM measurements were performed in lift mode, or two pass mode, with the lift height under which KPFM was performed being ~45 nm, ensuring the response was purely electrostatic in nature. CAFM scans were measured by applying a bias of 0.5 V to the tip. The local I-V curves were acquired in the conductive AFM mode, with the conductive tip in contact with the surface. The heating-cooling stage (Integrated temperature control from Asylum research) was used for temperature-dependent CAFM and


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

KPFM measurements. ToF-SIMS measurements were performed with TOF.SIMS.5-NSC (IONTOF GmbH, Germany), combining AFM and ToF-SIMS in the same vacuum chamber without breaking the vacuum. Measurements were done in two stages. First, the surface of the VO2 was modified by application of ±4 V electric bias through a conductive AFM tip while scanning a region of 5 × 5 µm2 in contact mode. Second, modified regions were studied using ToF-SIMS to identify possible changes in the sample’s chemical composition. The estimated time between AFM electrical writing and SIMS measurements was considerably short (< 1–2 min). We used a bismuth ion gun (energy 30 keV, current 0.5 nA, spot size ~120 nm) as a primary ion source. Measurements were performed in positive ion detection mode with tracking of V+ and VO+ mass peaks. Three-dimensional maps of the peaks’ spatial distribution were used to identify local changes in the chemistry of the studied sample. After the ToF-SIMS measurements, the sputtered crater was studied with AFM to perform depth calibration of the collected data. Ab initio Modeling: Ab initio calculations were made using the GGA+U approach as implemented in the Vienna ab initio Simulation Package (VASP).53,54 This approach consists of only one effective parameter (Ueff = U–J), where J is the screened exchange parameter and U is the Hubbard parameter. The calculations were performed on a 1 × 2 × 2 supercell of VO2 M1 phase using experimental lattice constants (a = 5.752, b = 4.538, c = 5.382 Å), a 4 × 4 × 4 kpoint grid, and a plane wave energy cutoff of 400 eV; and the internal coordinates were relaxed until the forces were lower than 0.01 eV/Å. The QMC calculations were performed employing the open-source QMCPACK software with workflows driven by the Nexus automation system.55 The simulations were performed in a 24-atom supercell of bulk VO2 with an imaginary time step of 0.01 Ha-1. The energies were twist averaged over a 2 × 2 × 2 supercell k-point grid. Based on 14 ACS Paragon Plus Environment

Page 14 of 28

Page 15 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

experience from prior studies, the obtained optimal U is expected to be insensitive to these choices. The hard norm-conserving pseudopotentials used in the QMC calculations were developed and tested previously in the context of atoms and dimers, showing comparable performance to all electron quantum chemistry methods.56 The nonlocal nature of the pseudopotentials was accounted for by the variational T-moves scheme.57 In case of oxygenvacancy formation energies, the chemical potential of oxygen is obtained from experimental values of oxygen bond formation enthalpy (5.15 eV) since the DFT predictions of O2 dimer formation is poor. The charge density difference plots shown in Fig. 5 correspond to an M1 model system with an O(I) vacancy for two different charge states (q = 0, +2). The atomic structure to plot the charge density difference is obtained by relaxing the internal coordinates of the VO2 M1 phase with O(I) vacancy for q = 0 with the corresponding functionals (PBE, PBE+Uopt, respectively). These plots are plotted with an isosurface value of 0.005 e/Å3. ASSOCIATED CONTENT Supporting Information. Results of the local I-V curves in CAFM mode, concentric-square KPFM and CAFM, temperature-dependent KPFM measurements, ToF-SIMS detection of H+ ions, and details of DFT+U calculations. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding author Ho Nyung Lee ([email protected]) Notes: The authors declare no competing financial interest.


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACKNOWLEDGMENTS This work was supported by the US Department of Energy (DOE), Office of Science, Basic Energy Sciences (BES), Materials Sciences and Engineering Division (synthesis) and as part of the Computational Materials Sciences Program (characterization and theory). Scanning probe microscopy and scanning transmission electron microscopy studies were performed as user projects at the Center for Nanophase Materials Sciences, which is sponsored at Oak Ridge National Laboratory (ORNL) by the Scientific User Facilities Division, BES, DOE. This research used resources of the Oak Ridge Leadership Computing Facility at ORNL, which is supported by the Office of Science of the DOE under Contract DE-AC05-00OR22725. This research used resources of the National Energy Research Scientific Computing Center, a DOE Office of Science User Facility supported by the Office of Science of DOE under Contract DEAC02-05CH11231. REFERENCES 1. Driscoll, T.; Kim, H. T.; Chae, B. G.; Kim, B. J.; Lee, Y. W.; Jokerst, N. M.; Palit, S.; Smith, D. R.; Di Ventra, M.; Basov, D. N. Memory Metamaterials. Science 2009, 325,1518‒21. 2. Zhou, J.; Gao, Y.; Zhang, Z.; Luo, H.; Cao, C.; Chen, Z.; Dai, L.; Liu, X. VO2 Thermochromic Smart Window for Energy Savings and Generation. Sci. Rep. 2013, 3, 3029‒34. 3. Liu, M.; Hwang, H. Y.; Tao, H; Strikwerda, A. C.; Fan, K; Keiser, G. R.; Sternbach, A. J.; West, K. G.; Kittiwatanakul, S; Lu, J; Wolf, S. A.; Omenetto, F. G.; Zhang, X; Nelson, K. A.; Averitt, R. D. Terahertz-Field-Induced Insulator-to-Metal Transition in Vanadium Dioxide Metamaterial. Nature 2012, 487, 345−348. 16 ACS Paragon Plus Environment

Page 16 of 28

Page 17 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

4. Lee, D.; Lee, J.; Song, K.; Xue, F.; Choi, Si-Y.; Ma, Y.; Podkaminer, J.; Liu, D.; Liu, SH.; Chung, B.; Fan, W.; Cho, S. J.; Zhou, W.; Lee, J.; Chen, L-Q.; Oh, S. H.; Ma, Z.; Eom, C-B. Sharpened VO2 Phase Transition via Controlled Release of Epitaxial Strain. Nano Lett. 2017, 17, 5614−5619. 5. Morin, F. J. Oxides Which Show a Metal-to-Insulator Transition at the Neel Temperature. Phys. Rev. Lett. 1959, 3, 34‒35. 6. Lee, S.; Cheng, C.; Guo, H.; Hippalgaonkar, K.; Wang, K.; Suh, J.; Liu, K.; Wu, J. Axially Engineered Metal−Insulator Phase Transition by Graded Doping VO2 Nanowires. J. Am. Chem. Soc. 2013, 135, 4850‒4855. 7. Strelcov, E.; Tselev, A.; Ivanov, I.; Budai, J. D.; Zhang, J.; Tischler, J. Z.; Kravchenko, I.; Kalinin, S. V.; Kolmakov, A. Doping-Based Stabilization of the M2 Phase in FreeStanding VO2 Nanostructures at Room Temperature. Nano Lett. 2012, 12, 6198−6205. 8. Pouget, J. P.; Launois, H. Metal-Insulator Phase Transition in VO2. J. Phys. Colloques 1976, 37, 49-57. (https://doi.org/10.1051/jphyscol:1976408). 9. Liu, M. K.; Wagner, M.; Abreu, E.; Kittiwatanakul, S; McLeod, A.; Fei, Z.; Goldflam, M.; Dai, S.; Fogler, M. M.; Lu, J.; Wolf, S. A.; Averitt, R. D.; Basov, D. N. Anisotropic Electronic State via Spontaneous Phase Separation in Strained Vanadium Dioxide Films. Phys. Rev. Lett. 2013, 111, 096602. 10. Cao, J.; Gu, J.; Fan, W.; Chen, L. Q.; Ogletree, D. F.; Chen, K.; Tamura, N.; Kunz, M.; Barrett, C.; Seidel, J.; Wu, J. Extended Mapping and Exploration of the Vanadium Dioxide Stress-Temperature Phase Diagram. Nano Lett. 2010, 10, 2667‒2673.


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 28

11. Nagashima, K.; Yanagida, T.; Tanaka, H.; Kawai, T. Stress Relaxation Effect on Transport Properties of Strained Vanadium Dioxide Epitaxial Thin Films. Phys. Rev. B 2006, 74, 172106. 12. Jeong, J.; Aetukuri, N.; Graf, T.; Schladt, T. D.; Samant, M. G.; Parkin, S. S. P. Suppression of Metal-Insulator Transition in VO2 by Electric Field-Induced Oxygen Vacancy Formation. Science 2013, 339, 1402−1405. 13. Jeong, J.; Aetukuri, N. B.; Passarello, D.; Conradson, S. D.; Samant, M. G.; Parkin, S. S. P. Giant Reversible, Facet-Dependent, Structural Changes in a Correlated-Electron Insulator Induced by Ionic Liquid Gating. Proc. Natl. Acad. Sci. U. S. A. 2015, 112, 1013−1018. 14. Passarello, D.; Altendorf, S. G.; Jeong, J.; Rettner, C.; Arellano, N.; Topuria, T.; Samant, M. G.; Parkin, S. S. P. Evidence for Ionic Liquid Gate-Induced Metallization of Vanadium Dioxide Bars over Micron Length Scales. Nano Lett. 2017, 17, 2796–2801. 15. Lee, S.; Sun, X-G.; Lubimtsev, A. A.; Gao, X.; Ganesh, P.; Ward, T. Z.; Eres, G.; Chisholm, M. F.; Dai, S.; Lee, H. N. Persistent Electrochemical Performance in Epitaxial VO2(B). Nano Lett. 2017, 17, 2229−2233. 16. Cavalleri, A.; Tóth, C. S.; Siders, C. W.; Squier, J. A.; Ráksi, F.; Forget, P.; Kieffer, J. C. Femtosecond Structural Dynamics in VO2 during an Ultrafast Solid-Solid Phase Transition. Phys. Rev. Lett. 2001, 87, 237401. 17. Hong, W. K.; Park. J. B.; Yoon, J.; Kim, B. J.; Sohn, J. I.; Lee, Y. B.; Bae, T. S.; Chang, S. J.; Huh, Y. S.; Son, B.; Stach, E. A.; Lee, T.; Welland, M. E. Hydrogen-Induced Morphotropic

Phase

Transformation

of

Single-Crystalline

Nanobeams. Nano Lett. 2013, 13, 1822‒1828.


Vanadium

Dioxide

Page 19 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

18. Nakano, M.; Shibuya, K.; Okuyama, D.; Hatano, T.; Ono, S.; Kawasaki, M.; Iwasa, Y.; Tokura, Y. Collective Bulk Carrier Delocalization Driven by Electrostatic Surface Charge Accumulation. Nature 2012, 487, 459–462. 19. Muraoka, Y.; Hiroi, Z. Metal-Insulator Transition of VO2 Grown on TiO2 (001) and (110) Substrate. Appl. Phys. Lett. 2002, 80, 583–585. 20. Kim, H.; Slusar, T. V.; Wulferding, D.; Yang, I.; Cho, J. C.; Lee, M.; Choi, H. C.; Jeong, Y. H.; Kim, H. T.; Kim, J. Direct Observation of the M2 Phase with Its Mott Transition in a VO2 Film. Appl. Phys. Lett. 2016, 109, 233104. 21. Sohn, A.; Kanki, T.; Sakai, K.; Tanaka, H.; Kim, D. W. Fractal Nature of Metallic and Insulating Domains Configurations in a VO2 Thin Film Revealed by Kelvin Probe Force Microscopy. Sci. Rep. 2015, 5, 10417. 22. Kim, Y.; Kelly, S. J.; Morozovska, A.; Rahani, E. K.; Strelcov, E.; Eliseev, E.; Jesse, S.; Biegalski, M. D.; Balke, N.; Benedek, N.; Strukov, D.; Aarts, J.; Hwang, I.; Oh, S.; Choi, J. S.; Choi, T.; Park, B. H.; Shenoy, V. B.; Maksymovych, P.; Kalinin, S. V. Mechanical Control of Electroresistive Switching. Nano Lett. 2013, 13, 4068‒4074. 23. Kim, H.; Hong, S.; Kim, D.-W. Ambient Effects on Electric Field-Induced Local Charge Modification of TiO2. Appl. Phys. Lett. 2012, 100, 022901. 24. Kim, Y; Bae, C.; Ryu, K.; Ko, H.; Kim, Y. K.; Hong, S.; Shin, H. Origin of Surface Potential Change during Ferroelectric Switching in Epitaxial PbTiO3 Thin Films Studied by Scanning Probe Microscopy. Appl. Phys. Lett. 2009, 94, 032907. 25. Chang, S. H.; Kim. J.; Phatak, C.; D’Aquila, K.; Kim, S. K.; Kim, J.; Song, S. J.; Hwang, C. S.; Eastman, J. A.; Freeland, J. W.; Hong, S. X-ray Irradiation Induced Reversible Resistance Change in Pt/TiO2/Pt Cells. ACS Nano 2014, 8, 1584–1589.


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

26. Liu, M.; Sternbach, A. J.; Wagner, M.; Slusar, T. V.; Kong, T.; Budko, S. L.; Kittiwatanakul, S.; Qazilbash, M. M.; McLeod, A.; Fei, Z.; Abreu, E.; Zhang, J.; Goldflam, M.; Dai, S.; Ni, G. X.; Lu, J.; Bechtel, H. A.; Martin, M. C.; Raschke, M. B.; Averitt, R. D. et al. Phase Transition in Bulk Single Crystals and Thin Flms of VO2 by Nanoscale Infrared Spectroscopy and Imaging. Phys. Rev. B 2015, 91, 245155. 27. Qazilbash, M. M.; Brehm, M.; Chae, B. G.; Ho, P. C.; Andreev, G. O.; Kim, B. J.; Yun, S. J.; Balatsky, A. V.; Maple, M. B.; Keilmann, F.; Kim, H. T.; Basov, D. N. Mott Transition in VO2 Revealed by Infrared Spectroscopy and Nano-Imaging. Science 2007, 318, 1750‒3. 28. Qazilbash, M. M.; Tripathi, A.; Schafgans, A. A.; Kim, B. J.; Kim, H. T.; Cai, Z.; Holt, M. V.; Maser, J. M.; Keilmann, F.; Shpyrko, O. G.; Basov, D. N. Nanoscale Imaging of the Electronic and Structural Transition in Vanadium Dioxide. Phys. Rev. B 2011, 83, 165108. 29. Kim, Y.; Morozovska, A. N.; Kumar, A.; Jesse, S.; Eliseev, E. A.; Alibart, F.; Strukov, D.; Kalinin, S. V. Ionically-Mediated Electromechanical Hysteresis in Transition Metal Oxides. ACS Nano 2012, 6, 7026–7033. 30. Zhang, Z.; Zuo, F.; Wan, C.; Dutta, A.; Kim, J.; Rensberg, J.; Nawrodt, R.; Park, H. H.; Larrabee, T. J.; Guan, X.; Zhou, Y.; Prokes, S. M.; Ronning, C.; Shalaev, V. M.; Boltasseva, A.; Kats, M. A.; Ramanathan, S. Evolution of Metallicity in Vanadium Dioxide by Creation of Oxygen Vacancies. Phys. Rev. Applied 2017, 7, 034008. 31. Ko, C.; Yang, Z.; Ramanathan, S. Work Function of Vanadium Dioxide Thin Films Across the Metal-Insulator Transition and the Role of Surface Nonstoichiometry. ACS Appl. Mater. Interfaces 2011, 3, 3396‒3401.


Page 20 of 28

Page 21 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

32. Bi, F.; Bogorin, D. F.; Cen, C.; Bark, C. W.; Park, J. W.; Eom, C. B.; Levy, J. “Water Cycle” Mechanism for Writing and Erasing Nanostructures at the LaAlO3/SrTiO3 Interface. Appl. Phys. Lett. 2010, 97, 173110. 33. McDonnell, L. A.; Heeren, R. M. A. Imaging Mass Spectrometry. Mass Spectrom. Rev. 2007, 26, 606‒643. 34. Ievlev A. V.; Maksymovych P.; Trassin M.; Seidel J.; Ramesh R.; Kalinin S. V.; Ovchinnikova O. S. Chemical State Evolution in Ferroelectric Films during Tip-Induced Polarization and Electroresistive Switching. ACS Appl. Mater. Interfaces 2016, 8, 29588‒ 29593. 35. Ievlev, A. V.; Brown, C.; Burch, M. J.; Agar, J. C.; Velarde, G. A.; Martin, L. W.; Maksymovych, P.; Kalinin, S. V.; Ovchinnikova, O. S. Chemical Phenomena of Atomic Force Microscopy Scanning. Anal. Chem. 2018, 90, 3475–3481. 36. Dudarev, S. L.; Botton, G. A.; Savrasov, S. Y.; Humphreys, C. J.; Sutton, A. P. ElectronEnergy-Loss Spectra and the Structural Stability of Nickel Oxide: An LSDA+ U Study. Phys. Rev. B 1998, 57, 1505. 37. Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865. 38. Xiao, B.; Sun, J.; Ruzsinszky, A.; Perdew, J. P. Testing the Jacob’s Ladder of Density Functionals for Electronic Structure and Magnetism of Rutile VO2. Phys. Rev. B 2014, 90, 085134. 39. Eyert, V. VO2: A Novel View from Band Theory. Phys. Rev. Lett. 2011, 107, 016401. 40. Zhu, Z.; Schwingenschlogl, U. Comprehensive Picture of VO2 From Band Theory. Phys. Rev. B 2012, 86, 075149.


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

41. Eyert, V. The Metal-Insulator Transitions of VO2: A Band Theoretical Approach. Ann. Phys. (Leipzig) 2002, 11, 650–702. 42. Verleur, H. W.; Barker Jr, A.; Berglund, C. Optical Properties of VO2 between 0.25 and 5 eV. Phys. Rev. 1968, 172, 788. 43. Lin, C.; Demkov, A. A. Electron Correlation in Oxygen Vacancy in SrTiO3. Phys. Rev. Lett. 2013, 111, 217601. 44. Lopez-Bezanilla, A.; Ganesh, P.; Littlewood, P. B. Magnetism and Metal-Insulator Transition in Oxygen-Deficient SrTiO3. Phys. Rev. B 2015, 92, 115112. 45. Marthinsen, A.; Faber, C.; Aschauer, U.; Spaldin, N. A.; Selbach, S. M. Coupling and Competition between Ferroelectricity, Magnetism, Strain, and Oxygen Vacancies in AMnO3 Perovskites. MRS Commun. 2016, 6, 182–191. 46. Zhang, L.; Liu, B.; Zhuang, H.; Kent, P. R.; Cooper, V. R.; Ganesh, P.; Xu, H. Oxygen Vacancy Diffusion in Bulk SrTiO3 from Density Functional Theory Calculations. Comput. Mater. Sci. 2016, 118, 309–315. 47. Lany, S.; Zunger, A. Accurate Prediction of Defect Properties in Density Functional Supercell Calculations. Modelling Simul. Mater. Sci. Eng. 2009, 17, 084002. 48. Nonnenmacher, M.; O’Boyle, M. P.; Wickramasinghe, H. K. Kelvin Probe Force Microscopy. Appl. Phys. Lett. 1991, 58, 2921–2923. 49. Mohn, F.; Gross, L.; Moll, N.; Meyer, G. Imaging the Charge Distribution within a Single Molecule. Nature Nanotech. 2012, 7, 227–231. 50. Zhong, Z.; Hansmann, P. Tuning the Work Function in Transition Metal Oxides and their Heterostructures. Phys. Rev. B 2016, 93, 235116.


Page 22 of 28

Page 23 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

51. Jacobs, R.; Booske, J.; Morgan, D. Understanding and Controlling the Work Function of Perovskite Oxides Using Density Functional Theory. Adv. Funct. Mater. 2016, 26, 5471– 5482. 52. Lee, S.; Meyer, T. L.; Park, S.; Egami, T.; Lee, H. N. Growth Control of the Oxidation State in Vanadium Oxide Thin Films. Appl. Phys. Lett. 2014, 105, 223515. 53. Kresse, G.; Furthmüller, J. Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B 1996, 54, 11169. 54. Kresse, G.; Furthmüller, J. Efficiency of Ab Initio Total Energy Calculations for Metals and Semiconductors Using a Plane-Wave Basis Set. Comput. Mater. Sci. 1996, 6, 15–50. 55. Krogel, J. T. Nexus: A Modular Workflow Management System for Quantum Simulation Codes. Computer Physics Communications 2016, 198, 154. 56. Krogel, J. T.; Santana, J. A.; Reboredo, F. A. Pseudopotentials for Quantum Monte Carlo Studies of Transition Metal Oxides. Phys. Rev. B 2016, 93, 075143. 57. Casula, M. Beyond the Locality Approximation in the Standard Diffusion Monte Carlo Method. Phys. Rev. B 2006, 74, 161102(R).

FIGURES


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 1. (a) Room-temperature XRD θ‒2θ scan of a VO2 thin film grown on a (001) Nb:TiO2 substrate. The inset shows the cross-sectional HAADF-STEM image of the VO2/Nb:TiO2 interface. (b) Resistivity vs. temperature curve of the VO2 thin film. The inset shows the corresponding derivative curve. (c) CAFM images at selected temperatures in the vicinity of the MIT of VO2.


Page 24 of 28

Page 25 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

Figure 2. (a) KPFM image showing the change in the CPD after the electrical writing processes. (a) Patterns were written in the contact mode by scanning the film surface under the tip-bias voltages of –4 V (top) and +4 V (bottom). (b) Topographic image of the same area after electrical writing showing no surface damage or morphology change. (c) CAFM image in (a) was taken under the reading voltage of 0.5 V. (d) KPFM image of a VO2 film after boxpatterning using the tip-bias voltages of ±4 V. (e) Schematic illustration of the possible electrochemical mechanism and KPFM imaging on the VO2 film surface.


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 3. (a and b) KPFM images of a VO2 film at 300 and 310 K, which were electrically written by incrementally increasing the negative tip-bias voltage. (c) The CPD value was found to decrease with increasing the applied tip voltage, where saturation in CPD was observed at the tip-bias voltage of ~ –2.4 V (at 300 K), indicating full metallization through oxygen-vacancy formation.


Page 26 of 28

Page 27 of 28 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano

Figure 4. ToF-SIMS imaging of the sample surface including the regions electrically written by tip-bias voltages of –4 V (left) and +4 V (right). Spatial distribution of (a) VO+ and (b) V+. (c) Line profiles of the variation of VO+ and V+ ion concentrations along the scan areas in (a) and (b). (d) Depth profile of the VO+ and V+ ion concentrations within the electrically written regions. Note that the pristine part of the sample in the ToF-SIMS images does not show any notable changes in the ion concentrations as compared to the electrically written areas, indicating that the oxygen vacancies are created during the writing process.

Figure 5. Charge density difference plots obtained for (a) PBE and (b) PBE+Uopt for VO2 M1 phase model with an O(I) vacancy for two different charge states (q = 0, +2). (c) Defect phase diagram of O vacancy formation in unstrained M1 phase. The black lines correspond to the O(I) phase, and the red lines correspond to the O(II) phase. (d) Defect phase diagram of O


ACS Nano 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

vacancy formation in M1 phase under strain.

Table of content graphic:

Three dimensional KPFM image showing the change in surface chemical potential after the electrical writing processes.


Page 28 of 28

Nanoscale Control of Oxygen Defects and MetalâInsulator Transition

Recommend Documents

Nanoscale Control of Oxygen Defects and MetalâInsulator Transition