Controlled Charging of Ferroelastic Domain Walls in Oxide

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Controlled Charging of Ferroelastic Domain Walls in Oxide Ferroelectrics Xian-Kui Wei,*,†,‡,∇ Tomas Sluka,†,∇ Barbara Fraygola,† Ludwig Feigl,†,§ Hongchu Du,‡,∥ Lei Jin,‡ Chun-Lin Jia,‡,⊥ and Nava Setter†,¶ †

Ceramics Laboratory, EPFL-Swiss Federal Institute of Technology, Lausanne 1015, Switzerland Peter Grünberg Institute and Ernst Ruska-Center for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich, 52425 Jülich, Germany § Institute for Photon Science and Synchrotron Radiation, KIT - Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, D-76344 Eggenstein-Leopoldshafen, Germany ∥ Central Facility for Electron Microscopy (GFE), RWTH Aachen University, Aachen 52074, Germany ⊥ The School of Electronic and Information Engineering, Xi’an Jiaotong University, Xi’an 710049, China ¶ Department of Materials Science and Engineering, Tel-Aviv University, Ramat Aviv 69978, Israel ‡

S Supporting Information *

ABSTRACT: Conductive domain walls (DWs) in ferroic oxides as device elements are a highly attractive research topic because of their robust and agile response to electric field. Charged DWs possessing metallic-type conductivity hold the highest promises in this aspect. However, their intricate creation, low stability, and interference with nonconductive DWs hinder their investigation and the progress toward future applications. Here, we find that conversion of the nominally neutral ferroelastic 90° DWs into partially charged DWs in Pb(Zr0.1Ti0.9)O3 thin films enables easy and robust control over the DW conductivity. By employing transmission electron microscopy, conductive atomic force microscopy and phase-field simulation, our study reveals that charging of the ferroelastic DWs is controlled by mutually coupled DW bending, type of doping, polarization orientation and workfunction of the adjacent electrodes. Particularly, the doping outweighs other parameters in controlling the DW conductivity. Understanding the interplay of these key parameters not only allows us to control and optimize conductivity of such ferroelastic DWs in the oxide ferroelectrics but also paves the way for utilization of DW-based nanoelectronic devices in the future. KEYWORDS: tunable charging, ferroelastic domain wall bending, domain wall conduction, doping, PZT

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be easily created, annihilated, and displaced by electric field or mechanical force.13,14 Strongly charged DWs, e.g., with a head-to-head polarization configuration, display superior electronic conductivity. However, their creation is very complicated due to their high formation energy15,16 and their low stability. Moreover, their existence is always mixed with presence of nonconductive DWs, e.g., tail-to-tail charged walls, neutral walls or antiphase boundaries as observed in ordinary ferroelectrics BaTiO3 and LiNbO3,8,17 improper ferroelectrics RMnO3 (R = Er, Ho)9,10 and charge-ordered multiferroics Pr(Sr0.1Ca0.9)2Mn2O7.11 The complex domain configuration hampers the practical usage of the charged DWs. Therefore, seeking for a solution of

INTRODUCTION

Domain walls (DWs) in ferroic oxides are found to possess exceptional and distinct properties against the domain bulk. The typical paradigms include conductive DWs in multiferroic BiFeO3,1,2 ferroelectric translational boundaries in antiferroelectric PbZrO3,3 ferromagnetic antiphase boundaries in Mndoped BaTiO3 multiferroics4 and electric polarization produced by magnetic DWs in multiferroic Lu2MnCoO6.5 These properties enable them to act as functional entities in the design of novel nanoelectronic devices.6 Among the topics, the DW conductivity attracts a widespread attention because of their ubiquitous existence in materials involving the ferroelectric order.7−11 In nonpolar materials, similar phenomena were also observed at heterointerfaces, e.g., at the LaAlO3/ SrTiO3 interface.12 Obviously, the movable DWs are more advantageous than the fixed heterointerfaces because they can © 2017 American Chemical Society

Received: October 28, 2016 Accepted: January 31, 2017 Published: January 31, 2017 6539

DOI: 10.1021/acsami.6b13821 ACS Appl. Mater. Interfaces 2017, 9, 6539−6546

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ACS Applied Materials & Interfaces

Figure 1. Identification of the in-phase and out-of-phase DW bending in the PZT films. (a) The bright-field TEM image viewed along [100] direction showing morphology of ferroelastic 90° domains in 45 nm thick film. (b−e) The dark-field TEM images recorded using the (01̅1̅) reflection spot from the a-domains in films with thickness of (b) 45 nm, (c) 54 nm, (d) 100 nm, and (e) 160 nm. (f) The ABF image of a 90° DW couple in the 45 nm thick film recorded along the [100] direction. The bending angle for wall-1 and wall-2 is α1 = 0.78° and α2 = 0.70°, respectively. (g) The ABF image of a 90° DW couple in the 160 nm thick film recorded along the [100] direction. The bending angle for wall-1 and wall-2 is α1 = 2.16° and α2 = −0.79°, respectively. The pink dashed lines and blue solid lines denote the neutral and real DW positions, respectively. Representative unit cells (Pb-yellow circles, Zr/Ti-blue circles and O-red circles) are magnified to denote the head-to-tail PS configuration across the DWs.

phase-field simulation. Our systematic study reveals that charging of the ferroelastic DWs is controlled by wall bending, type of doping, polarization orientation and work-function of the adjacent electrodes. In spite of intricate coupling of these factors, the type of doping is found to dominate conductivity of the partially charged DWs.

overcoming these obstacles becomes the primary task on the way ahead. Landau-Ginsburg-Devonshire phenomenological theory has predicted that purely ferroelastic 90° DWs, either in the c/a/c/a (c and a are out-of-plane and in-plane oriented domains, respectively) or the a1/a2/a1/a2 (a1 and a2 are in-plane oriented and orthogonal domains) domain structures, can be fabricated by means of strain engineering on the BaTiO3 and PbTiO3 films.18,19 Regarding this type of DW, previous studies pay more attention to their impact on enhancement of the piezoelectric responses.20−22 Recently, the electronic conductivity was discovered at such nominally neutral walls in ferroelectric Pb(Zr1−xTix)O3 (PZT) thin films.23 Unlike the defect-assisted neutral wall conduction in (3dn, n = 5) BiFeO3 multiferroics,1,24 variable-temperature experiments corroborates that the ferroelastic DWs exhibit intrinsic metallic-like conductivity in the 3d0-type ferroelectrics.14 The stable conductivity, uniform DW type and easy manipulation by external forces14,20 imply that such ferroelastic DWs are better candidates for implementation in practical devices. However, a full understanding on the conduction mechanism is still missing with respect to its dependence on the basic material parameters and boundary conditions. In this paper, the mechanism of charging and conductivity of the ferroelastic 90° DWs in the PZT (x = 0.9) thin films is investigated by (scanning) transmission electron microscopy ((S)TEM), conductive atomic force microscopy (c-AFM) and

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EXPERIMENTAL SECTION Materials. The PZT (x = 0.9) thin films were grown by pulsed laser deposition onto (110) plane of DyScO3 (DSO) substrates. According to the lattice parameters and crystallographic orientation relation, the a-axis of c-domain is under a tensile strain of 0.93% and 0.84% along the [001] and [11̅0] direction of the DSO substrate, respectively. While the c-axis of the a-domains is under a compressive strain of −4.53% and −4.62% along these two directions.23 Electron Microscopy Characterization. The STEM and electron energy loss spectroscopy (EELS) experiments were performed on FEI Titan 80−200 Chemi-STEM. The nanobeam electron diffraction (NBED) experiments were carried out on FEI 80−300 Titan microscope. Other characterizations were performed on Tecnai F20 microscope. Quantification of atomic column positions was carried out by fitting the column intensities with two-dimensional Gaussian peaks using a homemade software package DMPFIT25 supplied within the Gatan Digital Micrograph software based on the Levenberg− Marquardt technique from the MPFIT C code26 and a 6540


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0.70°, respectively (Figure 1f). While in the 160 nm thick film, the out-of-phase bending angle for wall-1 and wall-2 is α1 = 2.16° and α2 = −0.79°, respectively (Figure 1g). The domain width for these individual cases is w = 13.5 and 18.6 nm, respectively. Associated with direct imaging of all atoms including oxygen atoms, which are visible from the magnified unit cells, the polarization orientation of domains can be identified by the off-center displacements of the Zr/Ti and O atoms. Across the ferroelastic DWs, it is seen that a head-to-tail polarization configuration is kept in the PZT films. By statistics, we summarize the a-domain width (w) measured along the in-plane direction of the film and the average bending angle of the DW couples as a function of film thickness (Figure 2a). It is found that the bent angles of α1 + α2

MINPACK-1 Least Squares Fitting Library in C.27 The TEM specimens were prepared by conventional preparation procedures combined with focused ion beam system on FEI Helios Nanolab 400s. Our systematic examination did not find any dislocation either near the interfaces or at the DWs. AFM Characterization. The c-AFM and piezoresponse force microscopy (PFM) experiments were performed using an Asylum Research Cypher AFM (Asylum Research). The cAFM measurements were carried out with an Asylum ORCA current preamplifier. Conductive Asylec Ti/Ir-coated probes with nominal force constant 3 N m−1 were used for the c-AFM scans. Sample bias of +4 V was used. Note that for this bias polarity the out-of-plane polarization component does not switch. The thermal treatments on the 120 nm thick films were performed inside the Cypher chamber as well. After annealing at 300 °C under certain oxygen partial pressure, the samples are cooled down to room temperature at a rate of 10 °C min−1, respectively. Phase Field Simulation. The phase-field simulation results were obtained by numerical solution of a two-dimensional model that incorporates coupling between ferroelectric, electrostatic, elastic and electronic properties of PZT including immobile donors and acceptors. Here, an ideal interface between the film and the DSO substrate - without introduction of any lattice defect - is considered. The temperature used in the simulation is 50 K, but no qualitative difference occurs up to room temperature. The model of the acceptor doping is introduced analogically to the donor doping.14 Donor and acceptor concentrations are both 1012 cm−3, and their energy levels are located 0.6 eV from the conduction and valence bands (inside the bandgap), respectively. The Fermi level energies are precalculated from the charge neutrality condition for each of the two doping cases. The work-function difference between the PZT and electrodes is introduced as a constant potential at the film surfaces.

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RESULTS AND DISCUSSION The PZT thin films, which are grown on 30 nm thick SrRuO3 (SRO) buffered DSO substrates, contain large c-domains with up or down oriented spontaneous polarization (PS) and narrow a-domains with an in-plane polarization orientation. As an example, morphology of the regular c/a/c/a domain arrangement in a 45 nm thick PZT film is shown in a bright-field TEM image (Figure 1a). In our dark-field TEM experiments, it is surprisingly found that most a-domains have a parallelogram shape in the thinner films (thickness t < 100 nm), i.e., the DWs remain almost parallel to each other across the whole film (Figure 1b,c). While in the thicker films (t ≥ 100 nm), the adomains adopt a wedge shape, i.e., the domain width shrinks gradually from the film surface to the interface (Figure 1d,e). More relevant results are provided in Supporting Information, Figure S1. To understand the morphology difference between the adomains, atom-resolved annular-bright-field (ABF) imaging technique is applied to study the 90° DWs in the PZT films. With respect to orientation of the neutral wall planes (pink dashed lines), which strictly run along [011] direction of the tetragonal c-domains, we find that the real DWs (blue solid lines) bend in an in-phase and out-of-phase manner in the thinner and thicker films, respectively. Defining the wall bending angle as αi = αneutral -αreal, the in-phase bending angle for wall-1 (between cI and a-domain) and wall-2 (between aand cII-domain) in the 45 nm thick film is α1 = 0.78° and α2 =

Figure 2. Parameter statistics on the domains and the bent DWs in the PZT films. (a) Domain width and bending angle changes as a function of film thickness. The solid lines are fitting to the experimental points to mimic the trend. (b) The averaged bending angles for wall-1 (α1) and wall-2 (α2) as a function of film thickness. (c) Scaling of domain spacing changes as a function of film thickness. The solid line shows a linear fitting to the double logarithmic scales.

inversely scale with the a-domain width in the PZT thin films. Specifically, the average bending angles of α1 and α2 for films with different thickness are shown (Figure 2b). It is seen that in the thicker films, the α2 reverses its sign and the bending angle in the thicker films becomes larger than that in the thinner films (see details of measurement in Supporting Information, Figure S2). Meanwhile, the a-domain spacing (L) is found to change with the film thickness (Figure 2c). A linear fitting to the double logarithmic scale yields that the slope is n = 0.68 (n ≈ 2/3). Recent studies on the PbTiO3 thin films, grown on SrRuO3-electroded DSO and SrTiO3 substrates,28,29 have reported the similar scaling behavior. According to LandauLifshiftz-Kittel law,30,31 the domain spacing follows a square6541


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ACS Applied Materials & Interfaces root law (n = 1/2) as the film thickness is larger than 1−2 μm. However, a linear (n = 1) relationship was reported for a variety of ferroelectric films with thickness less than 1 μm.32,33 With consideration of the azimuthal hoop stress in cylinder nanostructures, which involve the axial and radial stress as well, Scott’s theoretical model well explains that the scaling behavior of nanodomains may fall in the range of 1/2 ≤ n ≤ 1.34 This indicates that the ferroelastic domains in our films are close to their equilibrium state. According to our experimental observation, models for the bent DWs in the PZT films are illustrated in an exaggerated fashion (Figure 3a,c). Measurement on the STEM images

yields that the DW bending is realized by presence of unit-cell scale steps. Although the polarization keeps the head-to-tail configuration across the walls, presence of the wall steps breaks charge neutrality of the ferroelastic DWs. For the in-phase bent DW couple, it can be seen that the wall-1 and wall-2 are positively and negatively charged, respectively (Figure 3b). To lower the electrostatic energy of the system, therefore, free electrons and holes are separately needed to compensate the bound polarization charges. For the out-of-phase bent DW couple, both walls are negatively charged and only the free holes are required to screen the polarization charges (Figure 3d). Clearly, the DW charging is dependent on the wall bending angle (αi). The larger the walls bend, the stronger the walls are charged owing to presence of more unit-cell steps. In order to investigate the charging status of the DWs, the c/ a ratio of the ferroelastic domains was measured to calculate the domain polarization according to the polarization-tetragonality relation P2S = κ(c/a-1).35 Based on positions of Pb atoms in the STEM images, the lattice parameters are measured across the DWs and averaged along the real DW planes (see details in Supporting Information). It is seen that in the thinner films, the a-axis of the c-domains is slightly larger than that of the adomain, while the c-axis of the a-domain is evidently larger than that of the c-domains (Figure 4a). Consequently, the c/a ratio for the a-domain, (c/a)a = 1.059, is larger than those for the cdomains, (c/a)cII = 1.054 and (c/a)cI = 1.048 (Figure 4b). The tetragonality changes are consistent with the results obtained by our NBED analysis (see Supporting Information, Figure S3), which are represented by the filled diamond symbols (Figure 4b). For the 160 nm thick film, the lattice parameter measurements reveal that the a-axis of the c-domains is slightly smaller than that of the a-domain, and the c-axis of the a-domain is evidently smaller than that of the c-domains (Figure 4c). Correspondingly, the domain tetragonality is (c/a)a = 1.050,

Figure 3. Schematic DW bending and charging in the PZT films. (a) The morphology and (b) unit-cell model for the in-phase bent DWs in the thinner film. (c) The morphology and (d) unit-cell model for the out-of-phase bent DWs in the thicker film. The (blue and pink) dashed lines and (black and green) solid lines indicate positions of the neutral and real DWs, respectively. The blue arrows in (b) and (d) denote directions of PS in the ferroelastic domains.

Figure 4. Strain-driven tetragonality modulation among the ferroelastic domains in the PZT films. (a) The lattice parameters and (b) c/a ratio measured from the 45 nm thick film. (c) The lattice parameters and (d) c/a ratio measured from the 160 nm thick film. The solid diamonds in (b) and (d) represent the c/a ratio determined by the NBED experiments. The black arrows in (b) and (d) indicate directions of PS in the ferroelastic domains. 6542


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Figure 5. Phase field simulation of acceptor doped PZT with high work-function Pt electrodes. (a−d) The results for the films with upward PS: the in-phase bent DW couple with with the c/a ratio (a), corresponding distribution of free holes p (b), electrostatic potential φ (c), and schematic illustration of the band diagram (d). (e−h) The results for the film with downward PS: the out-of-phase bent DW couple with the c/a ratio (e), similar concentration of free holes (f), potential profile (g), and the schematic band diagram (h). See also Figure S4 that virtually any bending and charging configuration can be achieved in thin PZT films by polarization orientation, doping and work-function of electrodes.

polarization direction, doping and work-function difference between the film and electrodes. The strongest in-phase bending in the upward polarized films, as presented above (Figure 1f), is achieved in the simulation with acceptor doping and high work-function electrodes (Figure 5a−d), such as platinum used to cover the top surface of the actual TEM specimens. However, the same conditions result in a strong out-of-phase bending when the film is polarized downward (Figures 5e-h). Additionally, the DWs tend to bend in-phase in the thinner films because the elastic energy becomes less significant and comparable with the DW energy.37,38 In this case, the total energy of the system is reduced when the DWs become shorter due to the in-phase bending. The bound charge at the DWs induces the corresponding potential bending, which either enhances the conduction or suppresses it depending also on the three factors: polarization orientation, type of doping and electrode work-function. These factors control the sign of DW charging and the sense of band bending−toward enhanced or suppressed conduction. This classifies the bent ferroelastic DWs in the oxide ferroelectrics as weakly charged DWs,39 and their conductivity depends on the additional factors. The illustrations of band diagrams show that the DWs with p-type conductivity in the acceptor-doped films requires the negative charging of the DWs and a high workfunction electrode (Figure 5d,h and Supporting Information, Figure S4). This can be achieved in the films with upward PS and in-phase bending of the DWs (Figure 5a), or in the films with downward PS and out-of-phase bending of the DWs (Figure 5e). Particularly, the polarization orientation mediated DW conduction in the thinner films has been evidenced by our previous c-AFM experiments.14 When the acceptor doping is replaced by the donor doping, the DW conduction for the above two configurations is completely inhibited, which highlights the significance of type of doping in tuning conductivity of the DWs. Exactly opposite scenario takes place in the donor-doped films (see Supporting Information, Figure S4). Overall, our simulation shows that in t < 45 nm thin PZT films the DW charging and conductivity can be controlled

(c/a)cII = 1.063 and (c/a)cI = 1.056. These values have a good agreement with those determined by the NBED (Figure 4d). It should be noted that the c/a ratio exhibits a jump change across the DWs, regardless of the film thickness. This indicates that the DW width ranges from 4 to 9 unit cells (1.6−3.6 nm) in the PZT films. Based on the domain and DW parameters, our calculation reveals that averagely, the net polarization charge (ΔPS⊥, the component normal to the wall planes) for the outof-phase bent DW couple is ΔPS⊥ = −9.3 μC cm−2 in the 160 nm thick film, which is about 3.6 times of the in-phase bent DW couple in the 45 nm thick film (see the Supporting Information, Table S1). The above results evidence that the DW bending and charging interact intimately with the strain in the PZT system. Since the in-plane lattice constants of the DSO substrate fits much better to that of the c-domains than that of the a-domains (see the Experimental Section), the overall domain and DW behaviors are dominated by the strain relaxation of the cdomains. In the thinner films, preservation of the tensile strain substantially reduces the c/a ratio of the c-domains over the whole film. Associated with the elastic lattice distortion at the interfaces, this in turn allows elongation of the c-axes of the adomains and enhancement of their c/a ratio. With increasing of the film thickness, the mismatch strain inside the c-domains is well relaxed and the stress in the c- and a-domains is relatively equilibrated. As a result, the a-domains become wider with increasing film thickness and they adopt the V-shape. This is evidenced by closeness of the (c/a)cII to the bulk value of PbTiO3 (c/a = 1.064).36 In order to understand the interplay among bending, charging and conduction of the ferroelastic DWs in the PZT films, especially in the thinner films, we designed a phase field model to figure out the structure−property relation. The simulation results well reproduce the structural features observed in the thinner films (Figure 5a). Unlike in the thicker films, where the out-of-phase bending is present always, our simulation further reveals that the DWs in the thinner films may bend in either sense, which depend on the interplay of 6543


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Figure 6. Effect of doping on the DW conduction in the PZT films. The c-AFM map for the out-of-phase bent DWs in the 120 nm thick film treated at 300 °C (2 h) under (a) 1 mbar and (b) 0.1 mbar O2 atmosphere, respectively. The insets are their PFM amplitude images. (c−g) The EELS results obtained on the 54 nm thick film. (c) The HAADF image recorded along the [100] direction illustrates the area for EELS study. The blue solid lines indicate the real DW positions. (d,e) The integrated intensity map for Ti-L2,3 and O-K in the energy range of 455.5−469.5 and 532−549 eV, respectively. (f,g) Comparison of the absorption-edge profiles for Ti-L2,3 and O-K extracted from the domains and DWs binned along the wallplane direction. The cyan shadows denote the energy ranges for integration of the EELS signal intensities.

to a larger extent. In the thicker films, the dependence of the DW conduction on doping and work-function of electrodes validates as well. Considering the identical charging status, the out-of-phase bent DWs in the thicker films allow us to probe and distinguish the effect of doping on the DW conduction. Our c-AFM experiments reveal that the conductivity of the DWs is very sensitive to the thermal treatments. When the film was treated at a high-pressure O2 atmosphere (1 mbar), the conductivity of the DWs was enhanced and can be well identified from the cAFM map (Figure 6a and its inset). While as the film was annealed at a low-pressure O2 atmosphere, e.g., 0.1 mbar, the DWs became nonconductive anymore (Figure 6b and its inset). This confirms that accumulation of the free holes at the negatively charged DWs account for conduction of the out-ofphase bent walls in the PZT films. Otherwise, the oxygen vacancies, which are introduced by annealing at the low O2 pressure and act as donor dopant, annihilate the free holes and therefore inhibits the DW conduction. By using EELS, the doping conditions at the DWs were further examined along the cross-sectional directions of the PZT films with different thickness. As an example, the results obtained on the 54 nm thick film are presented. The high-angle annular-dark-field (HAADF) image illustrates the region for collection of the EELS signal (Figure 6c), in which the parallel DWs bend in an in-phase manner with αi ≈ 0.36° (see the Supporting Information, Figure S5). For intuitive illustration, the signal intensities of Ti-L2,3 and O-K edges are integrated in

an energy window of 455.5−469.5 and 532∼549 eV, respectively. With respect to the domain areas, it is seen that the integrated intensity of Ti-L2,3 is evidently reduced at the DWs (Figure 6d). In contrast, the intensity reduction at the DWs for the O-K edge is inconspicuous and seems random over the detected area (Figure 6e). The profiles clearly display the intensity changes as identified from the maps (Figure 6f,g). This indicates that the Ti is deficient at the DWs, which acts as acceptor dopant and leads to the p-type conductivity in the thinner films.14,23 As a comparison, similar EELS experiments were carried out on the thicker films, but the DWs are nonconductive. It is found that the elemental deficiency was not detected at the nonconductive DWs (see the Supporting Information, Figure S6). This implies that the small density of bound charges at the bent DWs allows them to exist even without compensation of any free carrier. Clearly, formation of the regular ferroelastic domain pattern in the PZT films23,28 benefits from minimization of the electrostatic energy at the PZT/SrRuO3 interface. This excludes presence of the intricately coupled charged and uncharged 90°/ 180° domains as observed in the PbTiO3 thin films grown on the insulating substrates.40,41 Regulated by the interface strain, the wall bending leads to charging of the DW array. In bilayer tetragonal/rhombohedral PZT thin films, it has been shown that with thickness reduction of the top tetragonal PZT layer, the c/a/c/a domain pattern is altered to the a1/a2/a1/a2 pattern.42,43 Similarly, modification on the elastic mismatch stress at the interface drives rearrangement of the ferroelastic 6544



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domains. According to first-principles calculations,15 the formation energy of the 90° DW in the PbTiO3 is merely 35 mJ m−2 and the barrier for its motion is negligible. This provides favorable conditions to realize bending of the ferroelastic DWs and their conductivity via doping in both kinds of domain patterns.

CONCLUSIONS Our systematic electron microscopy investigation combined with c-AFM and phase field simulation demonstrates that charging of the ferroelastic DWs in the PZT films can be effectively tuned by wall bending, type of doping, polarization orientation and property of the electrodes. Among these coupled parameters, the doping plays a key role in controlling the DW conductivity. Understanding the interplay of these key parameters not only enables us to control and optimize the DW conductivity in the PZT films, but also provides a widespread adaptability of this methodology over other oxide ferroelectrics. It is believed that revealing of the coupled degrees of freedom may facilitate the development of DW-based nanoelectronic devices in the future. ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b13821. Figure S1: Overview of the a-domain morphology changes as a function of film thickness and their electron diffraction patterns; Figure S2: Method of measuring the domain wall bending; Figure S3: Nanobeam electron diffraction data from films with different thickness; Figure S4: Effects of polarization orientation, doping and work-function difference on domain wall charging simulated by phase field models; Figure S5: Domain wall bending measurement on the 54 nm thick film; Figure S6: EELS results for the nonconductive domain walls in the 160 nm thick film; Table S1: Estimation of the domain polarization and charging of the DWs based on the NBED results (PDF). (PDF)

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REFERENCES

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Research Article

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Xian-Kui Wei: 0000-0003-4320-1120 Tomas Sluka: 0000-0002-8092-2026 Hongchu Du: 0000-0002-4661-4644 Lei Jin: 0000-0001-6924-2364 Author Contributions ∇

These authors (X.K.W. and T.S.) contribute equally to this work. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The research leading to these results has received funding from the European Research Council under the EU seventh Framework Programmme (FP7/2007-2013)/ERC grant agreement no (268058) Mobile-W and Grant Agreement 312483ESTEEM2 (Integrated Infrastructure InitiativeI3). We thank C.S. Sandu, D. Meertens, and M. Kruth for sample preparation. 6545


Research Article


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Controlled Charging of Ferroelastic Domain Walls in Oxide

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