Tailoring the Doping Mechanisms at Oxide Interfaces in Nanoscale

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Tailoring the doping mechanisms at oxide interfaces in nanoscale Weitao Dai, Ming Yang, Hyungwoo Lee, Jung-Woo Lee, Chang-Beom Eom, and Cheng Cen Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b02508 • Publication Date (Web): 14 Aug 2017 Downloaded from http://pubs.acs.org on August 15, 2017

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Tailoring the doping mechanisms at oxide interfaces in nanoscale Weitao Dai1, Ming Yang1, Hyungwoo Lee2, Jung-Woo Lee2, Chang-Beom Eom2, Cheng Cen1* 1

Department of Physics and Astronomy, West Virginia University, Morgantown, West Virginia 26506, USA

2

Department of Material Science and Engineering, University of Wisconsin-Madison, Madison, Wisconsin 53706, USA

Here we demonstrate the nanoscale manipulations of two types of charge transfer to the LaAlO3/SrTiO3 interfaces: one from surface adsorbates and another from oxygen vacancies inside LaAlO3 films. This method can be used to produce multiple insulating and metallic interface states with distinct carrier properties that are highly stable in air. By reconfigurably patterning and comparing interface structures formed from different doping sources, effects of extrinsic and intrinsic material characters on the transport properties can be distinguished. In particular, a multi-subband to single-subband transition controlled by the structural phases in SrTiO3 was revealed. In addition, the transient behaviors of nanostructures also provided a unique opportunity to study the nanoscale diffusions of adsorbates and oxygen vacancies in oxide heterostructures. Knowledge of such dynamic processes is important for nano-device implementations.

Keywords: 2DEG, defects, adsorbates, subband engineering, diffusion

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Defects and surface adsorbates have significant influences on two dimensional electron systems1-9. Methods for manipulating local defects and adsorbates are highly desirable for nanoelectronics applications. In addition, once the nanoscale distributions of defects and adsorbates are modified, it is also valuable to understand their dynamic responses over the time. At LaAlO3/SrTiO3 (LAO/STO) interfaces, one significant carrier source originates from water dissociations at polar LAO surfaces, which produce proton surface states that facilitate charge transfer to the interface, and hydroxyl states that trap electrons4-6, 10. Another important carrier donor is oxygen vacancies (OV) in LAO. Unlike OVs in STO substrate which tend to generate three-dimensional carrier distributions11-13, OVs in LAO films1-3, assisted by the built-in electric field, can remotely dope electrons to the interface and produce two dimensional electron gas (2DEG). These two mechanisms often coexist and are very difficult to distinguish in electrical measurements. While the total density of mobile carriers results from the combined effect, different electron sources may have distinct influences on the carrier scattering14, 18

19-23

modulation and spin related phenomena

15

, lattice distortion6,

15-17

, bandwidth

. In order to understand and utilize the rich

functional properties found in oxide 2D systems19,

20, 24-34

, it is important to gain the

capability of tuning individual doping mechanism separately. Here, we show that, starting from the as-grown metallic interface of 8 unit cells (uc) LAO/STO, insulating state and different metallic states formed from either surface adsorbates or OVs can be produced by sequences of three main operation steps: oxygen plasma exposure, biased atomic force microscope (AFM) probe scan and polar solvent treatment. Not only can this method allow a wide range of tunable carrier properties to be reconfigurably tailored, the produced effects are also highly repeatable in different samples and very stable in air for at least months. These performances show clear advantages comparing to previous works on 3 uc LAO/STO where only protons related metastable switching can be observed6, 35. Furthermore, transient electrical measurements and scanning probe imaging were also carried out on 2DEG nanostructures formed from different doping mechanisms, which retrospectively revealed the interesting dynamic migration behaviors of surface protons and OVs. 2 ACS Paragon Plus Environment

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Figure 1 Methodology used in controlling the interface doping mechanisms. (a) Surface regions not treated by oxygen plasma is shown in darker shade, underneath which as-grown 2DEG (marked by green) is present at the interface. Surface region enclosed by black dashed lines were configured to form different interface states as illustrated in b,e,h, and k. (c, f, i) Schematics showing the relative contributions from the


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surface adsorbates to interface charge transfer (nsurf) and oxygen vacancies doping (nov) in stable (solid circles) and metastable (dashed circles) interface states. When the transitions between different states were induced following the brown arrows, interface conductance changes monitored between the interface contacts labeled as “C1” and “C2” are plotted in (d, g, j). Circular insets show the working principles of how each transition was controlled. The effects of the three main operation steps can be understood as the following. Oxygen plasma treatment increases the density of surface hydroxyls and suppresses the charges transferred from the surface adsorbates to the interface ( nsurf )5. While the added hydroxyls bond to the LAO surface very stably, the reduced nsurf can still be changed by manipulating protons that bond and dissociate very frequently. However, the changed proton density always restores to the equilibrium value within hours in air at room temperature6. Therefore, the reduced nsurf produced by oxygen plasma can be considered eq

as a new equilibrium value ( nsurf ) of the surface charge transfer (green dash lines in Figure 1 c,f,i). Local electric field induced by biased AFM probe affects both nsurf and OVs related interface doping ( nOV ) (Fig.1 d,g,j, inset1). First, it can ionize the surface water and produce either protonated (positive bias) or deprotonated surfaces (negative bias)6. Second, when the LAO layer is thick enough, due to the reduced vacancy formation energy at polar LAO surface1, it may also be possible for the probe field to create OVs and drive them toward the interface (positive bias), or attract existing OVs to the surface and annihilate them (negative bias). In the meanwhile, both polarities of probe field were found to favor the hydroxylation of LAO surface6. After the removal of probe field, re-equilibrium of surface proton density in air typically takes several hours. In experiments, this process can be greatly expedited by the quick proton solvation in polar solvents (e. g. isopropanol (IPA), insets 2 in Figure eq 1d,g,j)6. This step quickly restore nsurf back to nsurf but unlikely affect the OVs inside

LAO films. 4 ACS Paragon Plus Environment

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To monitor the transient behaviors of the interface conductance, an isolated conducting strip between two interface contacts was generated by masked oxygen plasma exposure (Fig.1a). Subsequent AFM scans and plasma treatments as shown in Figure 1 were all performed in the same 20 µm × 10 µm region (Fig.1a, enclosed by dashed lines). As shown in Figure 1c,d,e, a metastable insulating interface can be produced by raster-scanning with negative probe bias, which deprotonated the surface and removed the OVs formed during growth. After immersion in IPA, the surface protons instantly reequilibrated but the non-reducing solvent does not facilitate the re-creation of OVs. Consequently, the conductance of the scanned region recovered to a level below the original value, corresponding to a stable metallic state that is dominantly produced by

nsurf (Fig.1e). Then, the dashed region was also exposed to oxygen plasma, which significantly eq reduced nsurf and made this region insulating (Fig.1f). Scanning with positive probe bias

produced a metastable state with conductance much higher than the as-grown value (Fig.1g). After immersion in IPA, the interface conductance dropped significantly due to proton re-equilibrium but didn’t fully restore to the insulating state prior to biased AFM scan (Fig.1g). This residual conductance, also larger than as-grown values, was most likely contributed by OVs generated during positively biased AFM scan (Fig.1h). This OV dominated metallic state can be converted back to an insulating state by negatively biased AFM scan (Fig.1i,j,k). Unlike the metastable behavior observed in aseq grown state (Fig.1d), after the removal of OVs, scanned region with low nsurf remained

insulating even after IPA immersion (Fig.1j). As shown in Figure 1, interface conductance can be widely tuned by different combinations of nsurf and

nOV . Once configured into the four stable states as depicted by

Fig.1b.e.h.k, interface conductance remained constant throughout the weeks to months experiment windows. In addition, interface transport properties tuned by these methods are highly repeatable in different samples (Fig.S2).


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Using the method described above, 2DEG Hall bars in the three stable metallic states (Fig.1b,e,h) were sequentially patterned and characterized. Since all measurements were performed in the same sample area using identical contacts, reliable comparison between different doping sources can be made (Fig.2). At room temperature, after removing OVs from the LAO layer in as-grown sample, carrier density decreased by a small fraction from 3.0 × 1013 cm-2 to 2.4 × 1013 cm-2 (Fig.2c), indicating the major role of surface charge transfer in the spontaneous formation of 2DEG. Though when more OVs were introduced, much larger interface doping ( 6.4 × 1013 cm-2) can be generated even in absence of charge transfer from surface adsorbates.

Figure 2 Transport properties of different interface 2DEG states.

(a, b) Oxygen

octahedral rotations and shifting of three Ti t2g orbitals in (b) cubic and (a) tetragonal phase. Polarized optical microscope image in (a) shows the formation of tetragonal domain patterns in SrTiO3 substrate below T*. Scale bar is 20 µm. (c) Sheet carrier density and mobility measured in 2DEG Hall bars formed from different doping mechanisms.(d) Extracted carrier densities and mobilities in the light ( nL , subbands ( nH ,

µ H ) at the temperature of 2K.


µL ) and heavy

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The carrier densities n of the three states all decreased below 105 K (Fig.2c). We note that the n data shown in Figure 2c are effective values obtained from the slope of the Hall resistance (RH(B)) near zero field. At the interface, the three Ti t2g orbitals split into lighter dxy band and heavier dxz, dyz bands7, 8, 35 due to the 2D confinement. When multiple subbands are occupied, n is a complex function of both the densities and mobilities of carriers in each subband (supporting materials). As shown in Figure S1, the low temperature Hall resistances of the three states exhibited nonlinear dependences on the magnetic field, indicating the occupation of multiple subbands with different mobilities. In comparison, RH(B) above T* was highly linear for field less than 9 T. This means that, at higher temperatures, the mobilities of the occupied subbands were more similar and consequently n can be approximately viewed as the sum of carrier densities in all subbands. One possible mechanism for carrier reduction is the frozen-out of carriers from defects (e.g. OVs). However, in that case, the carrier reduction should start from room temperature following the Arrhenius law36. Instead, n measured here only started to reduce at T* = 105 K. This temperature corresponds to the cubic-to-tetragonal structural phase transition in SrTiO337, evidenced by the tetragonal domain patterns (Fig.2a) found below T* (supporting information). In order to better understand the nature of decreasing n below T*, low temperature carrier densities and mobilities in the light band ( nL, µL ) and heavy bands (

nH , µH ) were extracted from the field dependence of the Hall

resistance (Fig.S1) and shown in Figure 2d. It can be seen that, the carrier density sums ( nL + nH ) in the OV dominated state at 2 K was substantially smaller than the room temperature effective carrier density n(300 K), while the 2 K

nL + nH of the as-grown and

adsorbates dominated states were approximately equal to n(300 K). This observation suggests that the reduced n below T* measured in the three states might be of different origins. In the OV dominated state, when STO transits into the tetragonal phase, in-plane antiferrodistortive rotations of oxygen octahedral near the surface tend to shift the heavy dxz, dyz bands to higher energies38, which can cause a reduction of heavy bands carrier occupations at low temperatures (Fig.2a,b). For as-grown and adsorbates dominated states, the significant charge transfers from the surface protons are accompanied by outof-plane octahedral rotations propagating from the LAO surface into the substrate6, which 7 ACS Paragon Plus Environment

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may counteract the subband energy shifting effect induced by the tetragonal lattice distortion in SrTiO3 near the interface. In the meanwhile, the structure sensitive electronphonon interactions39 may also produce a varied carrier mobility ratio between the light and heavy bands in different structural phases. Once this ratio becomes larger in the low temperature tetragonal phase, the effective carrier density will also appear smaller due to the more significant conduction contribution from the light band with less carrier and larger mobility. Such effect, supported by the temperature dependent changes found in the linearity of RH(B) (Fig.S1) and effective mobilities (Fig.2c), is likely present in all three interface doping states. Figure 2d also shows that the light band mobility in the OV dominated state where surface charge transfer is suppressed by hydroxylation was substantially larger than the other two states at low temperature. Since the carrier densities in heavy and light bands were both similar for adsorbates and OV dominated states, the big difference in light band mobilities cannot be accounted for only by carrier density related variations in electronic screening and electron-electron scattering process40. Instead, the different remote dopants in the three states may have played a more important role in varying the interface mobility41 at low temperatures. Randomly distributed surface protons and ionized OVs not only can generate a fluctuating potential that cause carrier scattering in longer range, their correlation with interface band bending and lattice distortion may also produce disorders directly present at the interface. In particular, the substantial octahedral rotations associated with the charge transfer from the surface protons6 can induce significant impacts on the interface roughness42 and conduction bandwidth18. Once electron-accepting hydroxyls are added to the surface and pair with the protons, surface potential fluctuation and lattice distortion can both be suppressed6, leading to a higher light band carrier mobility observed in the OV dominated state. The heavy band mobility

µH

was one order of magnitude smaller than the light band and much less affected by

the nature of the remote dopants (Fig.2d). This indicates that

µ H at

low temperatures

might be more critically affected by dopant-independent scattering processes such as the polaronic coupling with phonon39, 43. At room temperature, since phonon scattering is the


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main mobility limiting factor42,

44

, influences exerted by different dopants was not

obvious and all three states exhibited very similar mobilities (Fig.2c). Besides of being created or annihilated, protons and OVs can also migrate on sample surface or in the film. Such process is negligible when large structures are used to study steady state material properties, but will become significant when the structure dimensions become small. We first consider proton motions in two nanostructures with inverted non-equilibrium surface proton distributions: One is an insulating nanogap with depleted proton density created by a single line scan of negatively biased probe across a 20 µm wide as-grown 2DEG channel (Fig.3a), and the other is a conducting nanowire with increased proton density created by a single line scan of positively biased probe across a 10 µm stable insulating gap (Fig.3e). We note that although densities of OVs were also changed, as shown in Figure 1 d and g, their effects on the interface conductivity are much weaker compared to the far-from-equilibrium surface proton densities.


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Figure 3 Effects of proton diffusion on the evolution of 2DEG nanostructures (a-d) show (a) the illustration of experiment setup, (b) diffusion induced changes in the surface proton density, (c) transient conductance measured in air, and (d) calculated transient conductances based on 1D diffusion model for nanogap. (e-h) show similar plots but for nanowire. (i) Transient conductance measured across a nanogap in vacuum. Sample space was vented to air at 2500 s. (j) Illustration of water assisted proton diffusion process When the nanogap was being created, conductance across the gap recovered so quickly that the channel was never completely cut and the conductance fully recovered within minutes (Fig.3c). This is very different from a raster scan created micron size gap which maintained its zero conductance for much longer time (Fig.1d). In contrast, the conductance along the nanowire only decreased slowly after its creation (Fig.3g), very similar to its micron-size counterpart (Fig. 1g). The distinct transient conductances measured across a nanogap and along a nanowire are typical signatures of diffusion process. For a nanogap, since the primary diffusion occurred in the same orientation of the current flow (Fig.3b), a small amount of protons diffusing into the narrow gap can change the conductance from zero to finite values. And in the case of a nanowire, since the current flows perpendicular to the primary diffusion axis, proton diffusion only caused a redistribution of local conductivity transverse to the wire while the net conductance was little affected. Solving 1D diffusion equation numerically and assuming the interface conductivity scales linearly with the surface proton density (supporting information), the calculated transient conductances across nanogaps (Fig.3d) show sensitive dependence on the initial width of the gap, while the constant conductance along nanowires have no size dependence (Fig.3h). Both results are consistent with the experimental observations. Beyond this simplified model, slow proton exchange with air6 and nonlinear conductance dependence on proton density are likely responsible for the hour-level conductance decrease observed in both nanowires and micron size channels. When performing similar nanogap cutting experiment in vacuum, the zero conductance across the gap was very stable (Fig.3i). Though after sample chamber was 10 ACS Paragon Plus Environment

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vented to the atmosphere, the quick conductance recovery highly resembling the prediction of diffusion model took place again, indicating an environment dependent proton diffusion speed. It was reported that the proton diffusion at oxide surfaces can be effectively mediated by the surface water molecules condensed in air

45

(Fig.3j). When

the molecular water desorbs in vacuum, the diffusion of surface protons becomes significantly suppressed.

Figure 4 Effects of OV diffusion on the evolution of 2DEG nanostructures (a) EFM images before the OV dominated 2DEG structure marked by white dashed lines was created. Regions marked by red lines are 2DEG regions already fabricated. Only the one connected to the AC driving voltage displayed a large contrast. Scale bars is 1 µm (b,c) EFM images taken (b) three hours and (c) ten days after the center structure was created. Possibly due to the relatively small energy barriers of vacancy hopping in oxides46, thermally activated OV diffusions are also present at room temperature. Unlike protons that re-equilibrate with the environment quickly, OVs created in the LAO film are highly stable, enabling us to study their migration in a much longer time scale. To directly visualize such process, we implemented a tailored electrostatic force microscopy (EFM) imaging mode (supporting information) designed to suppress undesired effects from surface potential and only image conducting features electrically connected to a particular contact (Fig.4a). This way, the migration of electron donating OVs can be monitored through mapping the evolution of OV-dominated 2DEG structures. As in regular EFM mode, spatial resolution of this method is restricted by the long range electrostatic interaction and 50 nm probe lift height. However, features larger than 100 nm can still be faithfully imaged.


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Figure 4 shows EFM images taken before and after an OV-dominated 2DEG pattern was created on an oxygen plasma treated insulating area by biased AFM scan and subsequent IPA immersion. This pattern was assembled from nanowire segments and two loops with diameters of 200 nm and 500 nm (Fig.4a). As can be seen comparing Figure 4b and 4c, the nanoscale details of the pattern became very blur after ten days, indicating the slow diffusion and congregation of oxygen vacancies at room temperature. In conclusion, we have demonstrated the local control of interface doping mechanisms in LAO/STO heterostructures. By controlling the adsorbates at the surface and OVs inside the LAO film separately, multiple levels of interface conductance can be flexibly configured at room temperature. At low temperature, interface carrier density is mainly limited by the structurally driven transition from multi-subband to single subband conduction. The electron mobility, however, is sensitive to the dopant type, and most significantly affected by the ionized surface adsorbates. In addition, the capability of generating non-equilibrium distributions of OVs and surface protons in nanoscale also helps revealing the active diffusions of these electron donors at room temperature. Associated Content. Supporting information is available free of charge via the Internet at http://pubs.acs.org: sample synthesis method, methods of biased AFM scans, polarized microscope, EFM, and 1D diffusion model calculations. Author Information Corresponding Author: *E-mail: [email protected] W.D. and C.C. performed the c-AFM, EFM and magnetotransport measurements. Y. M. performed the optical imaging. H.L., J-W. L., S.R. and C-B.E. performed the PLD thin film growth. C.C. performed the FEM simulations. W.D. and C.C. conceived the study, analyzed the data and wrote the paper. All authors discussed the results and commented on the manuscript. Notes: The authors declare no competing financial interests.


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Acknowledgement Experimental work at West Virginia University was supported by the Department of Energy Grant No. DE-SC-0010399 and National Science Foundation Grant No. NSF-1454950. The work at University of Wisconsin-Madison was supported by AFOSR under grant number FA9550-15-1-0334 and National Science Foundation DMREF grant under grant number 1629270. References 1. Yu, L.; Zunger, A. Nat. Commun. 2014, 5. 2. Li, Y.; Phattalung, S. N.; Limpijumnong, S.; Kim, J.; Yu, J. Phys. Rev. B 2011, 84, (24), 245307. 3. Lei, Y.; Li, Y.; Chen, Y. Z.; Xie, Y. W.; Chen, Y. S.; Wang, S. H.; Wang, J.; Shen, B. G.; Pryds, N.; Hwang, H. Y.; Sun, J. R. Nat. Commun. 2014, 5. 4. Fongkaew, I.; Limpijumnnong, S.; Lambrecht, W. R. L. Phys. Rev. B 2015, 92, (15), 155416. 5. Dai, W.; Adhikari, S.; Garcia-Castro, A. C.; Romero, A. H.; Lee, H.; Lee, J.-W.; Ryu, S.; Eom, C.-B.; Cen, C. Nano Lett. 2016, 16, (4), 2739-2743. 6. Adhikari, S.; Garcia-Castro, A. C.; Romero, A. H.; Lee, H.; Lee, J.-W.; Ryu, S.; Eom, C.-B.; Cen, C. Adv. Funct. Mater. 2016, 26, (30), 5453-5459. 7. Meevasana, W.; King, P. D. C.; He, R. H.; Mo, S. K.; Hashimoto, M.; Tamai, A.; Songsiriritthigul, P.; Baumberger, F.; Shen, Z. X. Nat. Mater. 2011, 10, (2), 114-118. 8. Santander-Syro, A. F.; Copie, O.; Kondo, T.; Fortuna, F.; Pailhes, S.; Weht, R.; Qiu, X. G.; Bertran, F.; Nicolaou, A.; Taleb-Ibrahimi, A.; Le Fevre, P.; Herranz, G.; Bibes, M.; Reyren, N.; Apertet, Y.; Lecoeur, P.; Barthelemy, A.; Rozenberg, M. J. Nature 2011, 469, (7329), 189-193. 9. Walker, S. M.; Bruno, F. Y.; Wang, Z.; de la Torre, A.; Riccó, S.; Tamai, A.; Kim, T. K.; Hoesch, M.; Shi, M.; Bahramy, M. S.; King, P. D. C.; Baumberger, F. Adv. Mater. 2015, 27, (26), 3894-3899. 10. Bi, F.; Bogorin, D. F.; Cen, C.; Bark, C. W.; Park, J.-W.; Eom, C.-B.; Levy, J. Appl. Phys. Lett. 2010, 97, (17), 173110. 11. Basletic, M.; Maurice, J. L.; Carretero, C.; Herranz, G.; Copie, O.; Bibes, M.; Jacquet, E.; Bouzehouane, K.; Fusil, S.; Barthelemy, A. Nat. Mater. 2008, 7, (8), 621-625. 12. Herranz, G.; Basletić, M.; Bibes, M.; Carrétéro, C.; Tafra, E.; Jacquet, E.; Bouzehouane, K.; Deranlot, C.; Hamzić, A.; Broto, J. M.; Barthélémy, A.; Fert, A. Phys. Rev. Lett. 2007, 98, (21), 216803. 13. Kalabukhov, A.; Gunnarsson, R.; Börjesson, J.; Olsson, E.; Claeson, T.; Winkler, D. Phys. Rev. B 2007, 75, (12), 121404. 14. Lin, C.; Demkov, A. A. Phys. Rev. Lett. 2013, 111, (21), 217601. 15. Lin, C.; Mitra, C.; Demkov, A. A. Phys. Rev. B 2012, 86, (16), 161102. 16. Shen, J.; Lee, H.; Valentí, R.; Jeschke, H. O. Phys. Rev. B 2012, 86, (19), 195119. 17. Luo, W.; Duan, W.; Louie, S. G.; Cohen, M. L. Phys. Rev. B 2004, 70, (21), 214109. 18. Imada, M.; Fujimori, A.; Tokura, Y. Rev. Mod. Phys. 1998, 70, (4), 1039-1263. 13 ACS Paragon Plus Environment

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19. Brinkman, A.; Huijben, M.; van Zalk, M.; Huijben, J.; Zeitler, U.; Maan, J. C.; van der Wiel, W. G.; Rijnders, G.; Blank, D. H. A.; Hilgenkamp, H. Nat. Mater. 2007, 6, (7), 493-496. 20. Kalisky, B.; Bert, J. A.; Klopfer, B. B.; Bell, C.; Sato, H. K.; Hosoda, M.; Hikita, Y.; Hwang, H. Y.; Moler, K. A. Nat. Commun. 2012, 3, 922. 21. Pavlenko, N.; Kopp, T.; Tsymbal, E. Y.; Sawatzky, G. A.; Mannhart, J. Phys. Rev. B 2012, 85, (2), 020407. 22. Lin, C.; Demkov, A. A. Phys. Rev. Lett. 2014, 113, (15), 157602. 23. Lee, M.; Williams, J. R.; Zhang, S.; Frisbie, C. D.; Goldhaber-Gordon, D. Phys. Rev. Lett. 2011, 107, (25), 256601. 24. Reyren, N.; Thiel, S.; Caviglia, A. D.; Kourkoutis, L. F.; Hammerl, G.; Richter, C.; Schneider, C. W.; Kopp, T.; Rüetschi, A.-S.; Jaccard, D.; Gabay, M.; Muller, D. A.; Triscone, J.-M.; Mannhart, J. Science 2007, 317, (5842), 1196-1199. 25. Caviglia, A. D.; Gariglio, S.; Reyren, N.; Jaccard, D.; Schneider, T.; Gabay, M.; Thiel, S.; Hammerl, G.; Mannhart, J.; Triscone, J. M. Nature 2008, 456, (7222), 624-627. 26. Gariglio, S.; Reyren, N.; Caviglia, A. D.; Triscone, J. M. J. Phys.: Condens. Matter 2009, 21, (16), 164213. 27. Cheng, G.; Tomczyk, M.; Lu, S.; Veazey, J. P.; Huang, M.; Irvin, P.; Ryu, S.; Lee, H.; Eom, C.-B.; Hellberg, C. S.; Levy, J. Nature 2015, 521, (7551), 196-199. 28. Bert, J. A.; Kalisky, B.; Bell, C.; Kim, M.; Hikita, Y.; Hwang, H. Y.; Moler, K. A. Nat. Phys. 2011, 7, (10), 767-771. 29. Ariando; Wang, X.; Baskaran, G.; Liu, Z. Q.; Huijben, J.; Yi, J. B.; Annadi, A.; Barman, A. R.; Rusydi, A.; Dhar, S.; Feng, Y. P.; Ding, J.; Hilgenkamp, H.; Venkatesan, T. Nat. Commun. 2011, 2, 188. 30. Bi, F.; Huang, M.; Ryu, S.; Lee, H.; Bark, C.-W.; Eom, C.-B.; Irvin, P.; Levy, J. Nat. Commun. 2014, 5. 31. Irvin, P.; Ma, Y.; Bogorin, D. F.; Cen, C.; Bark, C. W.; Folkman, C. M.; Eom, C.B.; Levy, J. Nat. Photon. 2010, 4, (12), 849-852. 32. Tebano, A.; Fabbri, E.; Pergolesi, D.; Balestrino, G.; Traversa, E. ACS Nano 2012, 6, (2), 1278-1283. 33. Rastogi, A.; Pulikkotil, J. J.; Auluck, S.; Hossain, Z.; Budhani, R. C. Phys. Rev. B 2012, 86, (7), 075127. 34. Ma, Y.; Huang, M.; Ryu, S.; Bark, C. W.; Eom, C.-B.; Irvin, P.; Levy, J. Nano Lett. 2013, 13, (6), 2884-2888. 35. Joshua, A.; Pecker, S.; Ruhman, J.; Altman, E.; Ilani, S. Nat. Commun. 2012, 3, 1129. 36. Liu, Z. Q.; Leusink, D. P.; Wang, X.; Lü, W. M.; Gopinadhan, K.; Annadi, A.; Zhao, Y. L.; Huang, X. H.; Zeng, S. W.; Huang, Z.; Srivastava, A.; Dhar, S.; Venkatesan, T.; Ariando. Phys. Rev. Lett. 2011, 107, (14), 146802. 37. Shirane, G.; Yamada, Y. Phys. Rev. 1969, 177, (2), 858-863. 38. Chang, Y. J.; Bostwick, A.; Kim, Y. S.; Horn, K.; Rotenberg, E. Phys. Rev. B 2010, 81, (23), 235109. 39. Cancellieri, C.; Mishchenko, A. S.; Aschauer, U.; Filippetti, A.; Faber, C.; Barisic, O. S.; Rogalev, V. A.; Schmitt, T.; Nagaosa, N.; Strocov, V. N. Nat. Commun. 2016, 7. 40. Gold, A. Phys. Rev. B 1987, 35, (2), 723-733. 41. Xie, Y.; Hikita, Y.; Bell, C.; Hwang, H. Y. Nat. Commun. 2011, 2, 494. 14 ACS Paragon Plus Environment

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Nano Letters

42. Su, S.; Ho You, J.; Lee, C. J. Appl. Phys. 2013, 113, (9), 093709. 43. Wang, Z.; McKeown Walker, S.; Tamai, A.; Wang, Y.; Ristic, Z.; Bruno, F. Y.; de la Torre, A.; Ricco, S.; Plumb, N. C.; Shi, M.; Hlawenka, P.; Sanchez-Barriga, J.; Varykhalov, A.; Kim, T. K.; Hoesch, M.; King, P. D. C.; Meevasana, W.; Diebold, U.; Mesot, J.; Moritz, B.; Devereaux, T. P.; Radovic, M.; Baumberger, F. Nat. Mater. 2016, 15, (8), 835-839. 44. Verma, A.; Kajdos, A. P.; Cain, T. A.; Stemmer, S.; Jena, D. Phys. Rev. Lett. 2014, 112, (21), 216601. 45. Merte, L. R.; Peng, G.; Bechstein, R.; Rieboldt, F.; Farberow, C. A.; Grabow, L. C.; Kudernatsch, W.; Wendt, S.; Lægsgaard, E.; Mavrikakis, M.; Besenbacher, F. Science 2012, 336, (6083), 889-893. 46. Inoue, S.; Kawai, M.; Ichikawa, N.; Kageyama, H.; Paulus, W.; Shimakawa, Y. Nat. Chem. 2010, 2, (3), 213-217.


Nano Letters

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

As-grown

Adsorbate dominant

Page 16 of 20

OV dominant

Insulating

LaAlO3 SrTiO3 H

ACS Paragon Plus Environment

O

OV

Al

a

Page 17 of 20

C1

O OV

C2

Al

c

d

n

Adsorbate dominant As-grown 2

nOV

n

Insulating

1 OV dominant

n

g

Insulating 2

H O OV

1

Time (10 s) 3

2

3 2

After several hours

IPA

h

OV dominant

2 0 0

1 1

Time (10 s) 3

0

1 2

EProbe

After several hours

0.5

k

IPA

0

nOV

Adsorbate dominant

4

j

OV dominant

1 1 E Probe

nOV

nsuf

eq suf

0 0

2

i

IPA

0.1

G (µS)

Oxygen Plasma

2

EProbe

e

0.2

1

f

n

1

G (µS)

n

As-grown

H

G (µS)

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

b

Nano Letters

0

1


Time (103s)

0

1

Insulating

E EF dxy

E

dyz

dxz

EF

dyz dxy

kx

c

dxz

2

As-grown

2

nsurf nOV

1

Adsorbate dominant

0

nsurf nOV

1

nsurf nOV

10

T (K)

100

2

1


T=2K

Page 18 of 20

nH

nL

0

OV dominant

6 4

3

kx

n (1013 cm-2)

T* = 105 K

n (1013 cm-2)

z

d

T > T*

µ (103 cm2/Vs)

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

b

T < T*

3

µ (103 cm2/Vs)

a

Nano Letters

µL

2

1

µH 0

Adsorbate OV As-grown dominant dominant

G

Vtip=10V

G (µS)

0

f

G

0.8 0.4

0

-w0/2 w0/2

x

g

I

j

4 t (103s)

6

8

1 atm

0

2 t (103s)

Proton diffusion

0.0 2

2 t (10 s)


0

3

200 nm

Dt (nm ) 2

300

h1

1

1 atm

w0=50 nm

0

2

0

-w0/2 w0/2

0.8

0

1 atm

3

x

0.4 10-5 mbar

d1

0.0

1

0

c

G/G0

I

n/n0

e

i

Nano Letters

n/n0

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

1

G/G0

b

G (µS)

Vtip=-10V

Page 19 of 20

G (µS)

a

0 0

Dt (nm2)

300

a

1

b

Nano Letters Va.c. After 3 hr


Va.c.

Page 20 of 20

1

IEFM (a.u.)

0

1

c After 10 days

IEFM (a.u.)

IEFM (a.u.)

OV dominant

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17

Va.c.

0

0

Tailoring the Doping Mechanisms at Oxide Interfaces in Nanoscale

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