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Article Cite This: ACS Appl. Nano Mater. 2019, 2, 3556−3569

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Electrostriction, Electroresistance, and Electromigration in Epitaxial BaTiO3‑Based Heterostructures: Role of Interfaces and Electric Poling Dana Stanescu,*,† Helene Magnan,† Brice Sarpi,‡ Maxime Rioult,† Thomas Aghavnian,†,‡ Jean-Baptiste Moussy,† Cindy L. Rountree,† and Antoine Barbier† †

Service de Physique de l’Etat Condensé, CEA, CNRS, Université Paris Saclay, CEA Saclay, 91191 Gif-sur-Yvette Cedex, France Synchrotron SOLEIL, Saint-Aubin, BP48, F-91192 Gif-sur-Yvette Cedex, France

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‡

S Supporting Information *

ABSTRACT: Ferroelectric materials hold significant promise for potential applications in a number of fields including spintronics and solar energy harvesting. When integrating them into heterostructures, it becomes of crucial importance to master the ferroelectric properties and to determine the influence of adjacent materials (substrates and/or top layers). We studied the role of interfaces on the ferroelectric properties of BaTiO3-based heterostructures elaborated on 1 atom % Nb:SrTiO3 (001) and Pt (001) substrates. Poled patterns were found to be more stable in time and shape for sandwich structures when the BaTiO3 layer is covered by a nonferroelectric oxide layer due to interface screening. Significant topography deformations occur when poling was performed with voltages above a threshold value and were found undoubtedly depending on both the nature of the substrate and the voltage polarity. Maximum deformations occur for negative poling voltages in BaTiO3 layers grown on Pt (001) and for positive ones for BaTiO3 grown on 1 atom % Nb:SrTiO3 (001). In both cases, the step was associated with an increase of the sample resistance. Clockwise and counterclockwise resistance hysteresis loop cycles obtained for BaTiO3-based heterostructures grown on Pt (001) and on 1 atom % Nb:SrTiO3 (001), respectively, are understood by oxygen ions and vacancies migration through the oxide layers depending on the substrate nature (oxide or antioxidant metal). The present work provides a global view of ferroelectric thin film behaviors and is important for the understanding of physical phenomena occurring upon poling in nanometric ferroelectric layers. KEYWORDS: ferroelectric materials, barium titanate, piezoresponse force microscopy, PFM, Kelvin probe force microscopy, KPFM, electric poling, topography deformation, memristor effect, ferroelectric thin films

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INTRODUCTION Novel nanocomposites structures including ferroelectric layers have attracted considerable research efforts in recent years because of their potential application in a number of fields including spintronics, photovoltaics, and solar water splitting. In spintronics for instance, the study of couplings between ferroic orders in multiferroic artificial systems combining at least two ferroic orders, like ferro(or ferri-)magnetism and ferroelectricity, has attracted the interest of researchers over the past couple of years.1−3 Concerning solar water splitting applications, the use of ferroelectric photoanodes (characterized by a built-in electric field) is a widely explored route to drive the separation of photogenerated carriers, to reduce the recombination rate, and thus to increase the hydrogen production efficiency.4 As a matter of fact, perovskite solar cells benefited in very recent years from huge efficiency improvements linked to the inclusion of ferroelectric layers.5 However, a detailed understanding of the phenomena behind these improvements has not yet been fully achieved, in particular because of the usage of very thin ferroelectric layers, © 2019 American Chemical Society

which is quite recent as compared to the long-standing knowledge of bulk materials and micrometer-sized devices. In heterostructures stacks, interfaces play a major role in the determination of the global properties. Consequently, the characterization of interface effects on the ferroelectric properties (for instance polarization switching with an external electric field, remnant polarization stability after poling, surface potential, etc.), in heterostructures containing a ferroelectric layer, raises nontrivial questions deserving in depth investigations. The efficient manipulation of the ferroelectric polarization by an external electric field in these systems remains very challenging. Indeed, stabilizing a saturated polar state is difficult as it was stressed by Zubko et al. in SrRuO3/ BaTiO3/SrRuO3 heterostructures at room temperature using a combination of piezoresponse force microscopy, dielectric measurements, and structural characterization.6 Received: March 19, 2019 Accepted: May 10, 2019 Published: May 10, 2019 3556

DOI: 10.1021/acsanm.9b00517 ACS Appl. Nano Mater. 2019, 2, 3556−3569

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ACS Applied Nano Materials

Figure 1. Typical structures of the samples considered in this study: S1 = BTO/Nb:STO, S2 = BTO/Pt, S3 = CFO/BTO/Nb:STO, and S4 = HEM/BTO/Nb:STO.

these measurements can easily be complemented and correlated to variations in film morphology and surface potential via fast-scan atomic force microscopy (FS-AFM) and Kelvin probe force microscopy (KPFM), respectively, on the same AFM platform, after the electric poling. Leakage current measurements were realized using the TF analyzer 1000 from aixACCT Systems GmBH. This work reports recent PFM, FS-AFM, and KPFM investigations revealing specific contrasts as a function of the TOLs and of the nature of the substrates in BTO based heterostructures. Results section focuses on (i) the remnant ferroelectricity and poled domains stability in single BTO and TOL/BTO structures, (ii) voltage threshold as evidenced from PFM poling, (iii) morphological deformations induced by high poling fields and how these are correlated with clockwise and counterclockwise resistance hysteresis cycles obtained on these heterostructures.

Herein, we consider heterostructures based on BaTiO3 (BTO), which is an archetypical ferroelectric material. It is an environmentally friendly lead-free material, which is wellknown for its ferroelectricity at room temperature, high permittivity, wide band gap, and numerous dielectric-based applications.7−9 The para- to ferroelectric phase transition occurs at about 130 °C in bulk BTO. This temperature increases for strained epitaxial BTO layers, suggesting improved ferroelectric properties upon substrate clamping. This transition is accompanied by a cubic to tetragonal structural transition upon decreasing the temperature. The electric polarization orientation lies along the [001] tetragonal axis in the ferroelectric phase configuration.10 Fortunately, this is one of the simplest ferroelectric configurations and is thus very well suited to in depth fundamental research approaches. In this study, we consider four significant samples (S1−S4 as depicted in Figure 1), extracted from large sets of samples (∼30) elaborated by atomic oxygen assisted molecular beam epitaxy (AO-MBE) with various growth conditions and thicknesses (Scheme 1). The highlighted samples are an

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RESULTS The chosen S1, S2, S3, and S4 samples provide a wealth of information concerning how the interfaces influence the ferroelectricity in BTO based heterostructures. More precisely, this study investigates ferroelectric polarization switching, poled regions stability in shape and time, charge screening at the interfaces, poling capability with an external field, and topography modifications induced by the poling. For that purpose, the interdisciplinary multiscale atomic force microscopy platform (IMAFMP) measures, through PFM, AFM and KPFM techniques, the response at the remanence of each ferroelectric system following external electric poling. The TF analyzer 1000 allows leakage measurements during the poling. Remnant Ferroelectricity in BaTiO3-Based Heterostructures. PFM experiments were realized at remanence after PFM poling on all structures. Figure 2 presents poling patterns (top row), PFM amplitude (middle row), and PFM phase images (bottom row) obtained on each sample. The top row indicates the values of the poling potentials for each case. As stated in the Experimental Section, amplitude images contain information about the polarization vector magnitude, while phase images contain information about the polarization vector orientation. Figure 2 indicates a contrast in both amplitude and phase images. The electric polarization orientations are indicated in the images: ⊗ for Pdown and ⊙ for Pup concerning regions polarized with V > 0 or V < 0, respectively. These results indicate that the heterostructures considered herein are ferroelectric and can be poled with an external electric field applied between the metallic PFM tip and the substrate. Although typical images of samples S1, S2, S3, and S4 are shown, the conclusions of global ferroelectric behavior originate from multiple experiments realized on the same

Scheme 1. Schematics of BaTiO3-Based Heterostructures Growth by Atomic Oxygen Assisted Molecular Beam Epitaxy (AO-MBE) in Ultrahigh Vacuum (UHV)

epitaxial ferroelectric BTO layer, deposited on 1 atom % Nb:SrTiO3 (001) (Nb:STO) (S1) or on Pt (001) (Pt) single crystals (S2) or as internal layer in sandwich structures including another top-oxide-layer (TOL): CoFe2O4 (CFO) (S3) and α-Fe2O3 (HEM) (S4). The ferroelectric behavior of these thin layer heterostructures is expected to strongly depend on charge screening mechanisms, layers thicknesses, strains, and chemistry at ferroelectric surfaces.11 To investigate these properties, we use predominately piezoresponse force microscopy (PFM),12 which is a very well-suited experimental approach both to polarize ferroelectric samples and to check the poling efficiency, ferroelectric polarization vector (orientation and magnitude), the stability in time of the ferroelectric remnant state, topography modifications induced by electric poling, etc. Furthermore, 3557


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Figure 2. PFM images as recorded on S1 = BTO (15 nm)/Nb:STO, S2 = BTO (25 nm)/Pt, S3 = CFO (9 nm)/BTO (21 nm)/Nb:STO, and S4 = HEM (5 nm)/BTO (15 nm)/Nb:STO: (a) PFM poling pattern composed of black (Vpoling < 0) and white (Vpoling > 0) regions; (b) PFM amplitude and (c) PFM phase images at ferroelectric remanence after the electric poling.

samples (in different regions with varying poling conditions) and a large number of similar heterostructures grown in the same conditions. All cases, within a total of ∼30 samples, reveal similar contrasts and results. All the PFM images in Figure 2 display with dotted lines the borders of the white/black patterns (from Figure 2 top row) in order to compare them with the effective poled surfaces after the poling. Single BTO layers (S1 and S2 samples) versus TOL/BTO bilayers (S3 and S4 samples) images display significant differences after poling. For the BTO layer case a clear blurring effect occurs at the borders. Moreover, it expands beyond the poling pattern edges. For the case of TOL/BTO bilayers, the read pattern is quasi-identical to the written one; i.e., borders are sharp and well-defined. Moreover, for samples S3 and S4 the polarized regions are more stable in time; PFM contrast remains significant after several days and/or even months. To understand this difference, we have to consider screening effects. In 2002, Kalinin and Bonnell13 found that the surface potential profiles near ferroelectric surfaces for unscreened and completely screened cases are entirely different. Moreover, they found larger potential gradients between two regions poled in opposite directions for unscreened interfaces as compared to screened ones. To explain our results within these assumptions, let us consider that the surface of a single BTO layer is predominantly unscreened. We also assume that in TOL/BTO structures the interface between the ferroelectric layer and TOL is completely (or almost completely) screened with free charges from TOL. In order to compare our results with the ones from Kalinin and Bonnell, we used KPFM to image surface potentials on both systems after electric poling

with a SCM-PIT tip. Figure 3a and Figure 3b present KPFM images obtained on BTO/Nb:STO and HEM/BTO/Nb:STO

Figure 3. KPFM images on (a) BTO (13 nm)/Nb:STO and (b) HEM (10 nm)/BTO (28 nm)/Nb:STO after PFM poling at ±8 V and ±10 V, respectively. (c) Surface potential profiles for the two samples. 3558


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Figure 4. (Top) PFM multivoltages experiment on BTO (25 nm)/Pt: (a) poling pattern for poling voltages between ±4 V and ±12 V; (b) height sensor topography image acquired during a PFM read; (c) height profile along the yellow line in (b). (Bottom) PFM multivoltages experiment on 15 nm CFO/5 nm BTO/Nb:STO. Poling voltages are from left to right ±4 V (d), ± 6 V (e), ±8 V (f), and ±10 V (g).

samples after ±8 V and ±10 V poling, respectively. A significant contrast occurs between the two regions poled with positive or negative voltages. Figure 3c presents potential profiles extracted from Figure 3a and Figure 3b along the highlighted yellow dashed lines. A single BTO layer displays a larger potential gradient (560 mV) as compared to HEM/ BTO bilayers (80 mV). This result confirms larger potential gradient in unscreened surface. Larger potential gradients between regions poled in opposite directions imply larger inplane electric fields on the poled surface as schematically drawn in Figure 3a. Therefore, we show that the blurring aspect of the poled regions borders observed in all single BTO layers (see S1 and S2 as examples) can be intuitively explained as a consequence of ferroelectric domains nucleation processes induced by non-negligible in-plane electric fields on the unscreened surfaces. For TOL/BTO structures (see S3 and S4), we show that the in-plane electric fields are less important, because they are determined by smaller potential gradients, and in consequence the polarized pattern is more stable and no propagation was observed. Voltage Thresholds Evidenced by PFM Poling. Height sensor images recorded in contact mode during PFM experiments or in tapping mode during KPFM experiments inform us about physical surface features (height, roughness, ...) of the sample surfaces. Figure S1 (from Supporting Information) presents PFM height sensor images for the same samples considered in Figure 2, which were recorded at the same time as the PFM images presented in Figure 2. With the exception of S2, the other three samples display a contrast in topography similar to the poling pattern. Although the contrast is very small for S1, it is still visible. An analysis over an extended number (>30) of experiments that we realized during the past several years on equivalent samples leads to the following observation: a contrast in topography appears when

the poling voltage (V) is larger than a threshold value (V > VL). VL depends on multiple parameters, including the quality of contact between the back side of the sample and the AFM sample plate (in our case the sample is fixed by vacuum aspiration), the leakage current through the ferroelectric layer (depending on layer thickness, crystallinity, chemical composition), air humidity, tip quality and lifetime, spatial resolution during writing process, etc. Therefore, in order to determine the VL value for one sample, it is necessary to take some precautions. The measurements have to be realized (1) on the same sample without unmounting it, (2) with the same tip, (3) during the same day in order to minimize tip aging/oxidation, (4) such that minimal changes in the ambient conditions (temperature or humidity) occur, and (5) without changing the PFM parameters during scans. For example, in the last case, by variation of the deflection set point parameter, the pressure exerted by the tip on the ferroelectric surface changes, which may locally modify the ferroelectric state. On top of that, the quality of the PFM poling and of the PFM reads after poling depends on the contact between the tip and the surface during the entire scan. Initial roughness of the sample, as for instance atomic steps coming from the substrate or defects, may perturb PFM writing or reading procedures and should be considered when analyzing contrasts. To illustrate the determination of the VL value, let us consider the following studies. Figure 4 presents a series of poling experiments realized on a BTO/Pt sample (S2) and on a CFO/BTO sample (S3). During this experiment, voltages were varied between ±4 V and ±12 V for BTO/Pt (Figure 4a) and between ±4 V and ±10 V for CFO/BTO/Nb:STO. The BTO/Pt sample has a significant surface roughness, which is due to the initial roughness of the single crystal Pt substrate (i.e., a significant miscut). Figure 4b presents the topography 3559


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Figure 5. (a) PFM height sensor image and (b) section profile for S3 (CFO (3 nm)/BTO (5 nm)/Nb:STO) after ±8 V PFM poling. (c) FastScan height sensor image and (d) section profile obtained on the same polarized pattern with FS-AFM.

image acquired during the PFM read. For voltages less than 10 V, no deformation is observed in topography. For high voltages (between ±10 V and ±12 V), a clear contrast exists; for this sample the deformation induced by high poling voltages is low with respect to substrate step heights (poling step ≪ atomic step). Hence, it is difficult to quantify the deformation in the surface. Nevertheless, it provides the threshold potential (VL) for the PFM poling which ranges between 9 and 10 V for this experiment. For the CFO/BTO/Nb-STO sample (Figure 4d− g, bottom) the patterns were written on the same sample at nearby places with increasing voltages. The pattern written at ±4 V does not lead to any measurable deformation, while patterns written above ±6 V reveal clearly a deformation linked to the pattern. One may note that the ±6 V pattern is only partially observable when read. In this case the threshold potential is thus around and below 6 V. Let us note that the vertical lines observed in Figure 4d−g do not correspond to real topography deformations but to a sinusoidal noise originating from the AFM setup; however, poled patterns remain visible. Other poling experiments providing better characterization of the deformation induced by a high poling voltage will be presented in the following. Morphological Surface Deformations Induced by PFM Poling at High Voltages. As mentioned in the Experimental Section, artifacts on PFM or KPFM topography images may originate from hardly quantifiable electrostatic forces between the tip and the ferroelectric surface. Therefore, when a contrast in topography occurs, its validity is frequently questioned. In order to qualify and quantify precisely the actual topographical changes, we undertook a complex experiment using two scanning heads, ICON and fast-scan, on the same AFM platform, on the same sample, and on the same object

(i.e., poled pattern). Scanning modes, tip and head details for PFM, KPFM, and FS-AFM microscopy techniques are summarized in Table S1 in Supporting Information. First, a PFM tip was used to pole a ±8 V pattern on a CFO/ BTO/Nb:STO (S3) sample. Figure 5a presents the height sensor image after poling. The step height (ΔH) induced by electric poling at high voltages is defined here as the difference between the heights (h) induced by negative (V < 0; Pup region) and positive (V > 0; Pdown region) poling voltages: ΔH = h(V < 0) − h(V > 0)

(1)

Figure 5b depicts the cross-section profile averaged over the blue zone highlighted in Figure 5a. ΔHPFM (ΔH from PFM measurements) varies from −2 nm on the left side to −1 nm on the right side of the image. From this profile, one observes that a positive voltage induces a significant elevation of the layer, while negative voltages give a small reduction of the layer thickness with respect to the as-grown state. Afterward, the AFM head was changed from the ICON to the FastScan head. As stated above, the tip was changed from a conductive tip to an insulating tip (FastScan A). A gold lithography grid with 30 × 60 μm2 windows enables a coordination system ensuring the same scan zone (3 × 2 μm2 region) when swapping the AFM head. The PFM poling was realized inside a well-identified window. Figure 5c presents the height sensor images from the FastScan of the very same pattern as in Figure 5a. Some variations between the patterns are noticeable. Figure 5d depicts the cross-section profile average over the red zone in Figure 5c. ΔHFS (ΔH from FastScan measurements) is now equal to −1.5 nm as shown in Figure 5d. Now, the same value occurs on the left and the right sides of the image corresponding to regions poled with the 3560


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Figure 6. PFM height sensor images on BTO (15 nm)/Nb:STO (a) and BTO (19 nm)/Pt (c); corresponding scan height profiles (b and d). Poling voltages are indicated for each sample.

same potential. The FastScan profile reveals that both the positive and negative voltages induce a dilatation of the layer with respect to the initial state; a larger deformation occurs for the positive poling voltages. From these observations, it is clear that a real morphology step is induced when writing to the sample with NanoMan. Nonetheless, the measured step value from the PFM read measurements is distorted with respect to FastScan AFM measurement. Conductive versus nonconductive tips make a significant difference in these measurements. Nevertheless, PFM topography measurements provide the right sign of ΔH, which is thus a pertinent parameter. Figure 5 reveals a negative ΔH step from both PFM and FastScan images. This provides the potential sign leading to the dominant deformation, V > 0 for S3. Thus, a visible contrast in the PFM height sensor image indicates a real deformation after the PFM poling. The reported values, however, have to be considered as qualitative since, as we have shown here, the real deformation is not accessible with a PFM conducting tip (both in contact and in tapping mode, see Figure S7 in Supporting Information) but needs to be determined with a nonconductive AFM detection setup in each case. Following this finding, we determined the ΔH signs acquired from several single BTO films deposited on two different substrates (Nb:STO and Pt) written with voltages above the threshold VL. Figure 6 presents typical height sensor images and corresponding profiles on BTO/Nb:STO (Figure 6a,b) and BTO/Pt (Figure 6c,d), while the raw deformation values are reported in Figure 7. On Nb:STO substrates, the region poled with positive voltages is significantly higher than the negative poled ones. Therefore, the ΔHPFM sign on BTO/ Nb:STO sample is negative (Figures 6a,b and 7a,c). On Pt substrates, the behavior reverses; a higher step occurs for

Figure 7. Observed deformations ΔHPFM (a); h (V < 0) and h (V > 0) for the BTO/Pt (b) and for the BTO/Nb:STO (c).

negative poling voltages. Consequently, ΔHPFM on this BTO/ Pt sample is positive (Figures 6c,d and 7a,b). We realized numerous experiments on BTO layers grown on Nb:STO or Pt substrates with varying BTO thicknesses. Figure 7a presents ΔHPFM as a function of the BTO thickness grown on Nb:STO (red squares) or Pt (black squares). Consistently, without any exception, the BTO grown on a Nb:STO substrate always provides a negative ΔHPFM, while the BTO grown on a Pt substrate provides a positive ΔHPFM. Same sign (ΔHPFM < 0) was obtained for CFO/BTO grown on Nb:STO as shown in Figure S2 (Supporting Information). This means that the substrate plays a nontrivial role on the stacks response to poling: a BTO layer on Nb:STO (Pt) substrate deforms more for positive (negative) voltages than for negative (positive) 3561


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Table 1. Minimum Potential Values for Which x2 (Piezoelectric Term) and x3 (Electrostriction Term) Are Larger than x1 (Ferroelectric Term) and for Which x3 Is Larger than x2 for 1 mm BaTiO3 Crystal, 0.2 μm BaTiO3 Film, 15 nm BaTiO3 Film, and 0.2 μm PZT Filma Vmin x2 > x1 sample

A

1 mm BaTiO3 crystal 0.2 μm BaTiO3 film 15 nm BaTiO3 film 0.2 μm PZT film

5 kV 1V 0.075 V 2V

x3 > x1 B

x3 > x2

A

1.5 V 0.1 V 1.5 V

2V 0.2 V 4V

B

A

B

3V 0.2 V 3V

25 kV 4V 0.3 V 8V

6V 0.4 V 6V

Set A: εr = 1000, PS = 0.18 C/m2, Q = 0.018 for PZT film; εr = 3000, PS = 0.26 C/m2, Q = 0.03 for BTO films. Set B: εr = 90, PS = 1 C/m2 for PZT film; PS = 0.35 C/m2 for 15 nm BTO films; PS = 0.4 C/m2 for 200 nm BTO films.

a

where x1 = xS is the strain due to the ferroelectric spontaneous polarization (PS) that may reach in the case of a single domain up to xS = QPS2, Q is the electrostriction coefficient, and ε = ε0εr the electrical permittivity. For single crystals, the d coefficient decomposes into a longitudinal and into a transverse contribution depending on PS. However, for thin films it is more appropriate to consider an effective parameter given by18

voltages. Moreover, intriguingly, the absolute deformation values (h) are fairly larger in the Pt case for negative poling potential, with respect to the Nb:STO case (Figure 7b,c). The dependence of the deformation as a function of the nature of the substrate will be discussed in the following. Topography step values measured on these samples after PFM poling are higher (nanometer scale step heights) than the typical deformations expected in ferroelectric samples during electric poling. As a matter of fact, the piezoelectric constant d33 (pm/V) of BaTiO3 gives us the order of magnitude of the deformation value per voltage. This constant equals 85.6 pm/V for a BaTiO3 single crystal14 and varies as a function of the domain structure (416 pm/V was obtained for BaTiO3 ceramics with nanodomain structure15) or of the crystallographic orientation of the substrate (BaTiO3 thin layers grown on SrTiO3 (001) or SrTiO3 (111) substrates have different d33 constants, 62.2 pm/V and 129.4 pm/V, respectively16), etc. Additional effects thus have to be taken into account. To understand the origin of the microscopic deformation observed here, let us recall some important principles that apply when exposing a material to an external electric field17 and in particular in the case of ferroelectric materials. Any material exposed to an electrical field will experience electrostriction resulting in a homogeneous distortion while ions will move eventually from a centrosymmetric crystal positions to noncentrosymmetric positions. A contraction in the direction of the field is usually expected and compensated by dilatation in the perpendicular directions; the reverse effect has also been observed. Upon field reversal the reverse sign of electrical field has to produce the same electrostriction so that the strain depends on even powers of the electric field E. Complementarily, an incremental field applied to an electrostrictive material will lead to incremental strain which needs to be reversible in electric field and will be described by odd powers terms of the electric field E. The lower power terms are usually sufficient to describe the behavior of a material. Producing strain by application of a field is described as the “inverse” piezoelectric effect, while the “direct” effect occurs when a strain produces a polarization, an effect that is only possible for noncentrosymmetrical materials. The difference derives from the fact that in the first case (“inverse” piezoelectric effect) one acts on the charge points, while in the second case (“direct” piezoelectric effect) one acts on the mass points. It results that for a ferroelectric material, the strain, x, experienced under an external electrical field, E, can be expressed as the sum of three terms, x1 + x2 + x3, as follows: x(E) = xS + dE + Qε 2E2

d = 2 Q eff (E) εPS

(3)

where the effective coefficient Qeff depends on the thin film and domain configurations that may differ somewhat from the hydrostatic bulk crystal values. The piezoelectric coefficient d is usually positive, but a number of systems with a negative coefficient were also reported. The thickness of the ferroelectric material has a tremendous incidence on the relative contribution to the total deformation for a given applied voltage since the electric polarization depends inversely on the thickness. To exemplify the effects, we consider several situations based on the most popular PbZrxTi1−xO3 (PZT)18 and BaTiO3 ferroelectrics.17,19−21 They have been widely studied over decades because of large piezoelectric moduli and all quantitative parameters were determined. Moreover, BaTiO3 is archetypical because it is a quite simple ferroelectric system due to the 4mm crystal symmetry reducing the electrostriction description to only four independent constants.17 In single crystals and single crystalline films, the polarization fully orients along the tetragonal direction. Let us consider typical parameters obtained from literature for each material. For PZT film the parameters may vary from εr = 1000, PS = 0.18 C/m2, Q = 0.018 18 to εr = 90, PS = 1 C/ m2 in ref 22. For BaTiO3 we found εr = 3000, PS = 0.26 C/m2, Q = 0.03 in refs 19 and 20, and PS = 0.35 C/m2 for a 15 nm BaTiO3 layer or PS = 0.4 C/m2 for 200 nm.23 We will focus on the 0−12 V electrical potential range because it is the typical range accessible in PFM setups. We calculated the strain under the electric field for four typical and illustrative situations, as summarized in Table 1 and Figure S3: (a) a bulk single BaTiO3 crystal with a macroscopic thickness of 1 mm, (b) a 0.2 μm, and (c) 15 nm thick BaTiO3 layer and (d) a typical ferroelectric device PZT film 0.2 μm thick. We assume that the sample is grounded and that the potential is applied on the surface without any loss. This is not completely true within PFM measurements because of a nonzero resistance of the overall setup, variable contact quality of the tip with the sample, and leakage through the FE thin film. The calculated strain corresponds thus to the upper limit of what could be observed. Relative deformations x1/x, x2/x, and x3/x were

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ACS Applied Nano Materials calculated for two sets of parameters (εr, PS, Q) A and B, which are indicated in Table 1. The curves are represented in Figure S3 in Supporting Information. Table 1 contains the minimum values of the potentials for which x2 (piezoelectric term) and x3 (electrostriction term) are larger than x1 (ferroelectric term) and for which x3 is larger than x2. The main observations that derive from the calculated expected strains in the 0−12 V electrical potential range are as follows. For the single BTO crystal case of 1 mm thickness (Figure S3a), the deformation is dominated by the spontaneous polarization (x1/x ∼ 100%) which is not surprising. It is not expected that a potential in the (0−12 V) range can substantially change the shape of a ferroelectric crystal. The actual deformation will depend on the eventual domain configuration. The electrostriction term (x3) contributes to (7 × 10−5)% to the total strain at 8 V, which is negligible. The piezoelectric strain (x2) becomes larger than the ferroelectric spontaneous one (x1) only above 5 kV, and the electrostriction term (x3) becomes larger than the piezoelectric term (x2) above 25 kV which is the typical range used to investigate single crystals in the past. For lower thicknesses of the BTO films, the situation is radically different. We observe that x2 > x1 for potentials greater than 1−1.5 V for a 0.2 μm thick BTO film and 0.075−0.1 V for a 15 nm thick BTO thin layer (Table 1 and Figure S3). Furthermore, x3 > x2 is reached for potentials higher than 4−6 V and 0.3−0.4 V, respectively. The electrostriction term, x3, for a 15 nm thick BTO layer at 8 V is huge; it represents 96% of the total strain, x. A typical 0.2 μm thick PZT film shows a similar behavior to a 0.2 μm thick BTO film. For the PZT film, potential thresholds for which x2 > x1 and x3 > x2 are 1.5−2 V and 6−8 V, respectively. At 8 V, x3 represents 44−54% of the total strain (x) that remains small (x = 5 × 10−3) compared to the 15 nm thick BTO case (x = 6). Such PZT ferroelectric devices with thicknesses of few hundreds of nanometers are typically used in the 2−5 V range for applications, as expected. We found that variations of the input parameters (parameter sets A and B) have only a little effect on the observations since the film thickness is the major parameter. In summary, for thick crystals large potentials are needed to influence ferroelectric crystals that otherwise mostly exhibit the spontaneous deformation due to the spontaneous polarization and the electrostriction becomes significant only for very large electrical potentials. For ferroelectric device layers of adequate thickness (hundreds of nanometers), the deformation is well approximated as being linear with the electrical field and is dominated by the piezoelectric moduli in the typical functioning range of a device. The deformation can be described as mainly piezoelectric with a modest modulation due to electrostriction. Finally, for thin films (tens of nanometers thick), the deformation can reach very high values at low potentials and is dominated by the electrostriction term. Moreover, the total deformation has to be understood as an electrostrictive deformation modulated by the piezoelectric coefficient that may enhance or reduce the electrostriction contribution. Such configurations appear as more suitable for microelectromechanical systems (MEMs). Importantly, from Hooke’s law it is obvious that the level of electrostrictive deformations that can be reached here is far above the elasticity limit. In Figure 7, we can observe step heights: ΔHPFM = h(V0), h(V0) values measured on BTO/Pt and BTO/Nb:STO samples after PFM poling. In all the cases, BTO layers experience a dilatation.

Interestingly, we observe that BTO layers grown on Pt or Nb:STO substrates are not equally affected by the electrostriction effect. Moreover, poling with positive or negative voltages do not induce the same dilatation values as expected (since the dominant deformation x3 ∼ E2). The largest residual dilatation is observed for BTO/Pt after negative (V < 0) poling. Smaller steps are obtained for the others three cases (on Pt for positive voltages, on Nb:STO for positive and negative voltages with ΔHPFM for Nb:STO always negative). This behavior strongly suggests that other parameters should be considered here that might contribute to balance the electrostriction and act as interfacial adhesion forces. One of these adhesion forces is the interface energy which is known to be fairly low for metal/oxide interfaces as compared to oxide/oxide ones. However, the strongly different results observed for positive and negative poling of BTO/Pt samples demonstrate that this adhesion force is insufficient to explain our observation. The systems herein require an interface force, which depends on the poling sign and on the substrate nature. As a matter of fact, ferroelectric polarization necessarily occurs concomitantly with the electrostriction since the field considered here is far above the field necessary to switch the polarization. Doing so, this polarization needs to be screened by available free charge carriers accumulated at the interface between the substrate and the ferroelectric layer. These concepts lead to the schematic representation given in Figure 8.

Figure 8. Schematic representation of the electrostriction vs interface adhesion effects for the Pt and Nb:STO substrate situations with respect to the ferroelectric layer polarization and required screening charge mobility. FE depicts the electrostatic forces that exist at the interface.

Consequently, an electrostatic force between opposite sign charges accumulated on both sides of the interface occurs. For the Pt substrate under negative poling, the orientation of the electric polarization requires screening by positive charges coming from the substrate, which are not available for Pt because it is a metal with only negative charges available at its Fermi level. Thus, in this case, no interface force is able to reduce the effect of electrostriction deforming the layer well above the elastic limit, leading to the huge residual step that we observe. In the reverse case, when the BTO/Pt is exposed to a positive potential, the required negative screening charges are easily available in Pt, and thus an important interface adhesion will result from electrostatic forces and the deformation will be 3563


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The external potential was applied between a contact on sample surfaces and the substrate as shown on the images from Figure S4. Leakage currents as a function of the applied potential are presented in Figure S5. The equivalent circuit of the sample is a ferroelectric capacity (CFE) in parallel with a resistance and in series with a diode associated with the interface. Low (high) resistance is related to high (reduced) leakage current through the oxide layers. For these measurements, the acquisition is realized when the current is stabilized, i.e., a long time after the potential was set. We recorded the leakage current induced by a step potential (step height = 0.5 V, step duration = 2 s, data acquisition between 70% and 90% of the step duration). In this case, the equivalent impedance of the sample is completely resistive, the leakage current (Ileak (V)) is mainly determined by conduction through pinholes, Schottky-like interfaces, oxygen vacancies migration, defects,30,31 etc. The Ileak(V) characteristic obtained on Nb:STO (Figure S5a and Figure S5b) confirms a diode-like electric conduction. As for PFM poling, the potential is applied on the surface and the substrate is grounded. The forward (respectively reversed) bias polarization of the junction corresponds to V < 0 (respectively V > 0). This result confirms that at the interface between Nb:STO and BTO, a potential barrier was formed with positive (respectively negative) charge accumulation on the BTO (respectively Nb:STO) side of the interface. Thus, the majority carriers in the BTO film are negative charges. Such a configuration demonstrates that the BTO film is more n-doped than the Nb:STO substrate and implies also that the BTO electronic affinity is smaller than the Nb:STO one. The diode behavior is less obvious from the measurements on samples of BTO grown on Pt. The Ileak(V) characteristic obtained in this case is mostly resistive (linear curve). This is likely a consequence of a greater number of pinholes, miscut, and/or defects participating in the electrical measurement when the metallic contact on the sample surface is larger (millimeter size) than for micrometric contacts (like on Nb:STO samples). Therefore, as observed, BTO/Pt samples appear to be less resistive than the Nb:STO ones as seen in Figure S5. However, from a theoretical point of view, the n-doped BTO/Pt interface is an archetypical case n-type semiconductor/metal Schottky junction with the Pt work function fairly larger than the BTO electronic affinity. Thus, the favorable electric flow configuration is expected to be identical to the BTO/Nb:STO case; i.e. when connecting the positive (or ground) terminal of a generator to the Pt and negative terminal to the n-type BTO, one attends to create a forward-biased state and, thus, large current flows. In summary, all interfaces studied here consist of junctions with Schottky barriers that are forward-biased for negative voltages. The fact that the BTO layer is more n-doped than the 1% Nb doped STO substrate stresses the existence of a very large number of available carriers, most likely due to a noticeable amount of oxygen vacancies. These concepts are important in order to understand the properties of our interfaces as described in the other sections. Electroresistance and Electromigration Effects. For each I(V) curve represented in Figure S5, we observe the existence of a clear hysteretic behavior. Very interestingly, the hysteresis directions (indicated by black arrows on each graph) for samples grown on Nb:STO and Pt substrates are opposite when the external voltage is swept between −VMAX and +VMAX.

minimal. This is consistent with the PFM image from Figure 6c, where the negative poled parts show assemblies of balls due to the excessive plastic electrostriction deformation, and the positive poled regions remain very flat with modest overall deformation. In the Nb:STO substrate case, we have to take into account a gradient in charge carriers in BTO and Nb:STO and the presence of a depleted junction region as will be detailed in the next section. Positive charges accumulate on the BTO side of the interface. This results in an additional interface electric field and thus in an additional polarization, Pnn, that can be parallel or antiparallel to the ferroelectric polarization of the layer, P. This results in a modulation of the interface adhesion, which overall remains always possible because both positive and negative charge carriers are available in the Nb:STO substrate. However, the interface polarization is parallel to the ferroelectric one when the poling potential is negative, providing a larger adhesion balancing the electrostrictive effects. Thus, from this analysis, we expect huge positive step height differences for Pt substrates and modest negative ones for Nb:STO substrates, which is exactly what we have observed (Figures 6 and 7). Electrical Behavior: Interface Junctions. When dealing with semiconductors and/or metal/semiconductor interfaces, the electric current flow is typically determined by the interface configurations and relative doping of the layers, i.e., the type of majority charge carriers. These concepts need to be addressed because they apply equally for continuous macroscopic interfaces and for PFM experiments since the tip is nothing more than a conducting top contact. Both Nb:STO and Pt substrates are conductive. In the former case, the conductivity originates from doping the SrTiO3 oxide semiconductor with 1 atom % Nb. Nb:STO is used very often in order to avoid charge effects that may harm in running some experiments, like in our case the PFM or KPFM. The second one, Pt, is an excellent reference metal. When BTO is grown on different substrates, the energy band structure at the interface depends on the values of the work functions (WF) and electronic affinities (χ) for semiconducting materials of both materials in contact; a depleted space charge region is formed at the interface following Fermi levels alignment. A Schottky-like diode behavior is thus expected when applying an external voltage.24 When the junction is forward biased, the respective majority charge carriers are repelled from the electrodes and the depletion region collapses for a small turn-on voltage (typically 0.2 V) leading to a conducting state. To the contrary, a reverse biasing expands the Schottky barrier and prevents the flow of electric current until, at very high voltages, the barrier breaks down leading to a high reverse current flow, damaging the component. A band diagram usually allows evaluating the type of junction. In the literature, work functions and electronic affinities for the materials used in this work are well documented. For instance, work function values for Pt may vary between 5.3 eV 25 and 5.6 eV.26 For Nb:STO the electronic affinity values between 3.64 and 4.1 eV (3.64,27 4.0 eV,28 4.1 eV 26,29) were reported. Finally, for BTO the electronic affinity varies between 3.76 and 3.9 eV (3.76 eV 27 to 3.9 eV 28). Because of the partial overlapping of the values, describing potential barriers at the interfaces considering only these considerations is thus difficult and a direct evaluation of the electric behavior is desirable. Therefore, in order to characterize the leakage current through these samples, we realized conduction measurements on four samples: BTO/ Nb:STO, CFO/BTO/Nb:STO, BTO/Pt, and TiO2/BTO/Pt. 3564


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Figure 9. Electric resistances calculated from Figure S5 as R = V/I for (a) 13 nm BTO/Nb:STO, (b) 15 nm CFO/27 nm BTO/Nb:STO, (c) 10 nm BTO/Pt, and (d) 10 nm TiO2/10 nm BTO/Pt. (e, f) sSchematics of O2− migration during positive and negative poling in Nb-STO and respective Pt substrate cases.

substrates or the surrounding environment (air), corresponding to the “oxygen reservoir”.36 When a positive potential is applied to the sample surface, O2− ions migration from the substrate toward the surface (Figure 9e and Figure 9f) is favored, and if they accumulate in a favorable oxide layer, they may eventually break reversible conducting filaments. Within these considerations and assuming our BTO layers are favorable resistive switching oxide layers, the substrate nature will have a tremendous consequence. Indeed, Nb:STO is a doped oxide that possesses an infinite reservoir of oxygen atoms and oxygen vacancies with respect to the thin BTO layer whereas Pt is unable to provide any. Consequently, the resistive switching mechanism will necessarily be of a different nature. (i) The Nb:STO substrate case is the most usual. If we assume that the substrate is the “oxygen reservoir”, we may consider that oxygen ions can flow between the layer and the substrate. If the surface is negatively poled, V < 0, oxygen vacancies will move toward the sample surface and accumulate inside the BTO layer, and the resistance will drop as soon as enough vacancies are available to form a conductive filament. Eventually vacancies may combine with oxygen from the air at the surface. Under a positive polarization, V > 0, the vacancies will move toward the substrate and the conduction filaments break up, resulting in a high resistance state. The vacancies migration determines leakage current increasing/decreasing or resistance increasing/decreasing as observed for samples on Nb:STO in Figures 9a. Importantly, in this scenario one expects the resistive drop to occur for positive poling potentials. (ii) For the Pt substrate case, no oxidation reaction is possible between Pt and O, similar to the oxidation reaction with Ag substrates in ref 33. Since we observe a clear resistive

In order to better visualize this aspect, we represent in Figure 9 the equivalent resistance values calculated as the ratio V/I. For each sample, we identify clear resistance hysteresis loops corresponding to memristive resistive switching effects.32,33 Interestingly, hysteresis direction was found to depend on the substrate: it appears clockwise on Pt and counterclockwise on Nb:STO as shown on the Figure 9. Importantly, the deformations in topography observed at remanence after poling with the PFM (Figures 6 and 7) are fully correlated to the resistive changes. The maximum step height in topography occurs when the sample resistance switches to the high resistance state, under potential values for which the conduction mode is suddenly modified (red arrows in Figure 9 indicate the resistance switch that also corresponds to the appearance of the large deformations). In summary, on Nb:STO substrates a maximum step height is observed at remanence after V > 0 poling, on Pt after V < 0 poling, and in both cases the step occurrence is associated with a sudden and important increase of the sample resistance. Surface deformations after PFM poling were reported recently by Vaghefi et al., and these authors attributed it to Joule heating induced by the current through the sample during the poling.34 Our results fully exclude this interpretation because the deformation occurs when the low current state sets in and consequently when the Joule effect (if any) is less important. From the literature, we know that memristor or resistive switching effects are typically associated with oxygen vacancies migration induced by poling voltages exceeding a threshold value.35 A conductive state is reached when oxygen vacancies form continuous conducting filaments and a more resistive state occurs when the filaments break up, most usually due to oxygen atoms filling the vacancies. Oxygen vacancies migration can be realized by oxygen ions migration from or toward 3565


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blurred patterns while TOL/BTO structures are almost completely screened with free charges from TOL leading to better defined patterns. Although one could intuitively expect that a metal substrate would be favorable for the electric poling of a ferroelectric layer, our findings show that this is not the case. Under the application of a negative voltage, the lack of oxygen vacancies diffusion in the metal leads to a detrimental weak interface adhesion and thus to a maximization of the electrostrictive effects. This results in huge plastic deformations in the oxide layer, which can be strong enough to reorganize the material in submicrometer balls. However, if positive poling is sought, a metal substrate will be very efficient. Doped oxide substrates appear as more convenient since electrostrictive effects remain in a smaller range while the resistive switching is of higher magnitude. The possibility to tune the local resistance state in nanometric ferroelectric films is a remarkable concept that deserves further investigations. These observations are of prime importance in the understanding of what is actually going on during a PFM poling and pattern reading.

switching phenomenon with Pt substrates and since Pt is unable to provide the necessary oxygen vacancies, the “oxygen reservoir” must come from another source which is most likely the air. In this case, the positive V > 0 potential will repel the surface oxygen vacancies toward the Pt substrate depleting the surface region allowing for the incorporation of new vacancies upon O2 leaving the material until the concentration of vacancies is high enough to create conduction filaments and thus a drop of resistance. To the contrary, for V < 0 the vacancies migrate toward the surface where they may be annihilated through O2 absorption from the air and the high resistance state will occur when the filaments break up because the density of vacancies can only decrease with a substrate unable to provide any compensating vacancies. Importantly, with a scenario where air is the “oxygen reservoir”, we expect the resistive switching toward the small resistance state at negative voltages as observed in Figure 9c. For the present resistive switching phenomenon, the presence of a top oxide layer (TOL) modifies the amplitude and shape of the hysteresis loop but not the fundamental mechanisms that are here driven by the nature of the substrate (Figure 9b and Figure 9d). Interestingly, the clockwise or counterclockwise resistance hysteresis loop behavior determines the actual “oxygen reservoir”: the air or the TOL in the first case and the substrate in the second case. We expect intuitively that using air as the “oxygen reservoir” hampers the phenomenon, which is consistent with a smaller relative electroresistance variation for the Pt substrate while the topographic deformation is huge, at least the largest observed here.

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CONCLUSION We have studied the electric poling in crystalline BaTiO3 (BTO) layers and the role of interfaces on the ferroelectricity. Our correlated analysis between PFM, KPFM, and FS-AFM experiments demonstrates that both interfaces between BTO and substrate (Nb:SrTiO3(001) or Pt(001)) or between BTO and a top oxide layer (α-Fe2O3 or CoFe2O4) play a very important role on ferroelectric properties of BTO-based heterostructures. We found that the BTO layer remains ferroelectric both as a single layer deposited on two different substrates (Nb:STO and Pt) or when it is covered by different top oxides like CoFe2O3 or α-Fe2O3. All the heterostructures considered herein can be poled with an external perpendicular electric field between a conductive tip and the substrate. Poled patterns on unscreened BTO surfaces are blurred and less stable due to larger in-plane electric fields on the surface. Heterostructures containing a ferroelectric BTO layer covered by an oxide have screened surfaces, and the poled patterns are sharper and more stable. Our ferroelectric BTO-based heterostructures consist of Schottky junctions forward biased for negative voltages and can easily be polarized electrically. Significant electrostriction deformations of the topography were observed when poling is performed with voltages higher than a threshold potential. This effect is more significant for negative poling voltages when the BTO is grown on Pt and for positive poling voltages when grown on Nb:STO. In both cases, the step formation is associated with a sudden increase of the sample resistance. Clockwise and counterclockwise resistance hysteresis loops were obtained for BTO/Pt and for BTO/Nb:STO, respectively. Resistance variation is well explained by oxygen vacancies migrations due to high electric fields and an oxygen reservoir depending on the nature of the substrate. We demonstrated that PFM poling in the 0−12 V potential range leads to the concomitant occurrence of a number of phenomena that have to be taken into account during a PFM investigation. Complementarily to the ferroelectric nature of our samples, a number of important physical phenomena of practical use in devices have been evidenced here like memristor behavior and Shottky barriers. Moreover, this PFM study of high quality

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DISCUSSION In the present study, we have investigated the PFM responses of a large number of samples containing a nanometric BTO layer and having conductive metallic or oxide substrates and eventually top oxide layers. While the assumption of a linear piezoresponse for PFM poling of layers in the micrometer thickness range, typical for device applications, is generally correct, it needs to be revised in the case of sub-100-nm-thick layers. Using a PFM tip to poll a nanometric ferroelectric layer leads not only to ferroelectric polarization switching but also to a concomitant large number of electro-driven phenomena that we have disentangled because of a combination of samples and additional techniques. At large electric fields that can be reached with modest electric potentials for thin enough layers, electrostriction becomes the dominant phenomenon that leads to irreversible plastic deformation of the layers that may be moderated in the case of favorable interface adhesion. These plastic deformations are easily imaged qualitatively in PFM and let appear a remanent (residual) topography state in amplitude images, while the phase image corresponds of course to the ferroelectric orientation as usually expected. The actual morphology deformation can however only be accurately determined by a clean AFM measurement with an insulating tip. At the same time and for the same voltage range, the huge electrical field that is applied leads to electromigration of oxygen vacancies in amounts large enough to produce memristor effects. The mechanisms involved depend on the nature of the substrate and are different if the substrate can provide oxygen vacancies or not, leading to electroresistance switching when topographic steps are produced at positive or negative potentials. Finally, the top layer is also of importance; single BTO layers are predominantly unscreened yielding 3566


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ACS Applied Nano Materials ultrathin nanometric oxide and ferroelectric films supported by conducting substrates, as they can be produced nowadays, also paves the way to studies of materials exposed to very high electric fields within light setups, without requiring the need for very high electric potentials. The electrostriction regime has been demonstrated to be accessible within the 0−12 V range for films of thicknesses below 15 nm.

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For that, we calibrated deflection sensitivity parameter (in nm/V) before each writing/reading experiment. We consider that when a contrast in the amplitude image occurs between two regions, the polarizations in these regions have different magnitudes. Nonetheless, this does not necessarily pertain to opposite orientations. When a contrast in the phase image appears, it means that the polarization vectors are oriented in opposed directions but their magnitudes are not necessarily equal. Phase images presented herein have phase scales spanning 180°. KPFM measures the surface potentials of ferroelectric multilayers poled with the PFM. On the IMAFM platform, we can easily switch between PFM and KPFM experimental modes, using the same tip (SCM-PIT) and head (ICON). For KPFM measurements, the surface is scanned in tapping mode and potentials are applied on the tip in order to minimize the electrostatic forces between the tip and the surface; KPFM images map these potential values. Both PFM and KPFM require the use of metallic cantilevers, typically metal-coated silicon or silicon nitride. Artifacts on heights images originate from cross-talk effects with the electrostatic forces, electrostriction, flexoelectricity, and electrochemical strain.46,47 One should note however that charge effects on phase and amplitude images may appear in PFM, which can induce frequently thought erroneous responses in the height sensor images; they normally disappear after a few scans (several minutes). In order to overcome these issues, topography measurements were also realized using a nonconductive AFM setup, on the same microscopy platform with the FastScan head and a silicon nitride insulating cantilever (FastScan-A, Bruker AFM probe). Hence, the artifacts mentioned above are eliminated, and the FastScan head authenticates the step height measurements. It is important to note here that using a gold lithography grid or some impurities (on samples that do not allow patterning) enables a coordination system ensuring the same scan zone for all microscopy techniques used in our study that need swapping the AFM head, scanning mode, and tip. A table that presents technical details (head, tip, scan mode) is presented in the Supporting Information, Table S1. Leakage Current Measurement with TF Analyzer 1000. In order to characterize electric conduction through BTO-based heterostructures, we use the TF analyzer 1000 from aixACCT Systems GmBH (Aachen, Germany). Leakage measurements were realized by applying an external step potential between a metallic contact on the sample surface and the substrate. A soft contact was realized on the sample surface (Figure S4) using a spring tip applied (i) on a micrometer-sized gold pad patterned by laser lithography for samples grown on Nb:STO or (ii) on a millimeter-sized silver paste contact on samples deposited on Pt substrates. For both situations, we checked the reproducibility of all the measurements presented here.

EXPERIMENTAL SECTION

Epitaxial BaTiO3-Based Heterostructures Elaboration. The elaboration of BTO-based heterostructures is realized in a dedicated atomic oxygen assisted molecular beam epitaxy (AO-MBE) setup.37 Substrates herein concern single crystalline (001) oriented 1% Nb doped SrTiO3 (Nb:STO) and Pt surfaces. Crystals were cut and polished within state-of-the-art surface science procedures providing surfaces with less than 0.1° nominal miscuts. In situ reflection high energy electron diffraction (RHEED) characterizes the crystallographic structures of the films during the growth process. As shown in Figure S8, all the samples are characterized by very well contrasted RHEED patterns with intense 2D streaks corresponding to a perovskite crystal structure. In situ Auger electron spectroscopy (AES) and X-ray photoelectron spectroscopy (XPS) permit the characterization of the chemical composition and stoichiometry of the films after the growth. More details concerning film deposition and in situ characterizations can be found in our previous works for BTO,4,38,39 CFO,40 or HEM.41,42 A more detailed analysis concerning BTO growth on Pt single crystals will be reported in a dedicated paper. Importantly, our epitaxial layers grown on Nb:STO experience high strains in surface plane growth due to substrate clamping. For total thicknesses up to 30 nm, the strain is shared between the BTO and the overlayer as we have reported in ref 43. Epitaxial (001) BTO layers are thus more tetragonal than the bulk counterpart and are wellknown to favor a pure out-of-plane electric polarization.10 Multifunctional Atomic Force Microscopy To Correlate Ferroelectricity with Potential Surface and Topography. The interdisciplinary multiscale atomic force microscopy platform (IMAFMP, based on the Bruker FastScan/Icon Nanoscope V atomic force microscope) at the SPEC laboratory enables near field microscopy measurements on BTO-based heterostructures. This platform enables various microscopy experiments like PFM, KPFM, and FS-AFM, in order to correlate ferroelectricity, surface potential, and topography, respectively, of the ferroelectric heterostructures. PFM is used both to polarize specific patterns at positive or negative potentials and to measure the remnant ferroelectric state (see Supporting Information, Figure S6). Details concerning the PFM technique and measurements at resonance44 can be found in Supporting Information. A single PFM scan provides a wealth of information via the acquisition of simultaneous images, including (i) the surface topography, (ii) the amplitude signal, proportional to d33DA, and (iii) the phase, φ, of the piezoelectric signal. From these images, we can deduce the mean polarization vector (P ⃗ ) inside the film, at remnant state. The polarization vector magnitude is proportional to the amplitude signal, and its orientation is given by the phase: in-phase (0°) or antiphase (180°) values corresponding to the two orientations of polarization (up or down). Nevertheless, a phase difference less than 180° frequently occurs between regions polarized in opposite directions. This effect is linked to parasite complex background signals45 due to the experimental setup (cantilever stiffness, tip pressure on the surface, contact quality between the substrate, and AFM’s sample plate, etc.). Small changes of the ac electric field frequency may decrease this background. Moreover, if the DA is too large during the PFM reading, it may disturb and even reverse the polarization state of the sample via the reading procedure. The images presented herein are optimized as follows: the DA was chosen to be low enough not to change the ferroelectric remnant state of poled regions but large enough to measure a significant ferroelectric signal. We mention here that the force exerted by the tip on the FE surface was kept constant during the scans and it has approximately the same value for all the samples.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsanm.9b00517.

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PFM poling and reading experimental details, scanning modes, tips, heads, and height sensor images and data; relative deformation calculated data; electrical measurements setup and leakage currents; complementary RHEED patterns at the end of the growth (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: +33 1 69 08 75 48. ORCID

Dana Stanescu: 0000-0003-1881-8495 Helene Magnan: 0000-0002-2239-5408 Maxime Rioult: 0000-0002-8361-7238 Antoine Barbier: 0000-0002-1032-299X 3567


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ACS Applied Nano Materials Notes

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The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge the “Agence Nationale de la Recherche (ANR)” for their funding through the PHOTOPOT (Grant ANR-15-CE05-0014) and IOBTO (Grant ANR15-CE09-0005-01) projects. This work was also supported in part by “Triangle de la Physique” and “Ile-de-France” (C’Nano IdF, DIM-OXYMORE, and ISC-PIF) under the IMAFMP and MAEBA grants. We acknowledge support from DIM OxyMore under grant BANCMULTI.

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