Aurélie Solignac
email: aurelie.solignac@cea.fr

Georg Kurij

Ruben Guerrero

Guillaume Agnus

Thomas Maroutian

Claude Fermon

Myriam Pannetier-Lecoeur

Philippe Lecoeur

M Pannetier- Lecoeur

Spin electronic magnetic sensor based on functional oxides for medical imaging

Keywords: Spin electronic, magnetic sensors, magnetic tunnel junction, functional oxides, interface engineering

INTRODUCTION

Magnetic tunnel junctions (MTJ), consisting of two ferromagnetic electrodes separated by an insulating barrier, are one of the most promising devices for magnetic sensing purposes. However their high sensitivity is counterbalanced by a high level of noise, thus leaving the noise floor of the magnetic field sensors at the same level than other magnetoresistive devices [START_REF] Stutzke | Low-frequency noise measurements on commercial magnetoresistive magnetic field sensors[END_REF]. Several routes have been proposed to enhance the signal to noise ratio in such a case, most of them reducing the low frequency noise in MTJs [START_REF] Edelstein | Validation of the microelectromechanical system flux concentrator concept for minimizing the effect of 1/f noise[END_REF][START_REF] Guerrero | Low frequency noise in arrays of magnetic tunnel junctions connected in series and parallel[END_REF]. However another way to improve the performances is to increase the tunneling magnetoresistance (TMR) ratio by using other kinds of materials.

A good example of such an approach is the use of manganites as ferromagnetic electrodes instead of metallic ferromagnets. These materials exhibit a half metallic character, i.e. only one spin direction is present at the Fermi level, thus they exhibit a high TMR up to 2000% at low bias and at low temperatures [START_REF] Lu | Large magnetotunneling effect at low magnetic fields in micrometer-scale epitaxial La 0.67 Sr 0.33 MnO 3 tunnel junctions[END_REF][START_REF] Viret | Low-field colossal magnetoresistance in manganite tunnel spin valves[END_REF][START_REF] Bowen | Using halfmetallic manganite interfaces to reveal insights into spintronics[END_REF][START_REF] Werner | Improved tunneling magnetoresistance at low temperature in manganite junctions grown by molecular beam epitaxy[END_REF]. Moreover, manganites are part of the family of perovskite oxides possessing a large range of physical properties, which can be combined in epitaxial heterostructures to develop high performance devices. One example is mixed sensors which are ultrasensitive magnetic sensors targeted at the detection of biomagnetic signals such as neuronal signals [START_REF] Pannetier | Femtotesla magnetic field measurement with magnetoresistive sensors[END_REF]. The use of oxides is expected to improve the performances of these sensors at 77K. Indeed they will combine the gain induced by an YBaCuO superconductor flux-to-field transformer and the high magnetoresistance of a La 0.7 Sr 0.3 MnO 3 (LSMO) half metal-based MTJ's.

Nevertheless, we, as other authors in literature, do not reproduce the highest TMR ratios in LSMO-based tunnel junctions. The electrode/barrier interface has been shown to be primordial in MTJ and especially several effects could appear at the LSMO/barrier side, for example the presence of dead layer (magnetic and in transport), different terminations or oxygen vacancies [START_REF] Sun | Thickness-dependent magnetotransport in ultrathin manganite films[END_REF][START_REF] Ott | Interface magnetism of La 0.7 Sr 0.3 MnO 3 thin films studied by neutron reflectometry[END_REF][START_REF] Pailloux | Nanoscale analysis of a SrTiO 3 / La 2/3 Sr 1/3 MnO 3 interface[END_REF] which disturb the ideal interface. We present, in this paper, a systematic investigation of these effects and their impact on the TMR ratio by playing with the quality of the barrier, the interface and the electrode, thanks to a change of materials and growth conditions of the different MTJ layers.

SAMPLE PREPARATION

The MTJ stacks are deposited on a SrTiO 3 (100) substrate, with a nominal miscut lower than 0.1°, and a chemically and thermally treated surface at 900°C under pure oxygen in order to ensure a single TiO 2 termination [START_REF] Kawasaki | Atomic control of SrTiO 3 surface for perfect epitaxy of perovskite oxides[END_REF]. The stacking sequence and thicknesses are identical for all the junctions grown and studied in this paper. They are comprised of a SrTiO 3 (STO) insulating barrier and have the following structure: La 0.7 Sr 0.3 Mn 0.995 Ru 0.005 O 3 (LSMRO) (20nm) /LSMO (10nm) /STO (3nm) /LSMO (30nm) // STO. The Ru doping in the LSMRO layer increases its coercivity so that it plays the role of a hard layer in the MTJ [START_REF] Ishii | Perovskite manganite magnetic tunnel junctions with enhanced coercivity contrast[END_REF]. A 10 nm thick LSMO layer is maintained between LSMRO and STO to have a better interface. In order to determine the key parameters for the junction properties, we have grown junctions under different conditions and with different materials. The samples and their growth conditions are listed in Table 1. For each sample, one parameter is changed compared to the reference sample grown with "standard" conditions. This reference sample has been fully characterized and possesses properties similar to the ones presented in Ref. [START_REF] Guerrero | Low frequency noise in La 0.7 Sr 0.3 MnO 3 based magnetic tunnel junctions[END_REF]. With this serie of junctions, we wanted to test the effects of:

-The interface with the barrier, by replacing the STO with LaAlO 3 (LAO) and by using an STO(110) oriented substrate. The corresponding junctions should possess a reduced dead layer at the interface between the barrier and the electrode [START_REF] Valencia | Interfacial effects in manganite thin films with different capping layers of interest for spintronic applications[END_REF].

-The barrier quality by doing a growth of the STO barrier under lower O 2 pressure followed by an annealing under oxygen. These growth conditions are the optimal conditions for the STO in our PLD system. The second possibility is to use a crystalline STO target instead of the sintered STO target, so as to have a better STO layer [START_REF] Oh | Thermal conductivity as a metric for the crystalline quality of SrTiO 3 epitaxial layers[END_REF]. The monocrystalline STO target is a STO substrate of 30mm diameter (Crystec).

-The quality of the LSMO electrodes by growing them under higher O 2 pressure and at lower temperature.

All the different stacks are grown by pulsed laser deposition with a KrF laser (248 nm wavelength) operated at 2 Hz repetition rate with a fluence of 3J/cm 2 . Deposition rates of 0.13ML/s for the LSMO and the LSMRO, and 0.39ML/s for the STO layer are obtained by tuning an attenuator on the beam path. A homogenizer ensures a top hat profile of the laser beam on the target. The full stack is grown under 120mTorr of oxygen at 585°C. At the end of the deposition process, the chamber is filled with oxygen up to 300mTorr for 10 minutes before the cooling down step. In order to control the deposition rate a thickness calibration is performed with single layers measured by X-Ray reflectivity. During the deposition of the stack all the conditions are kept as in the mentioned calibration step in order to deposit the desired thicknesses.

After preliminary characterizations using Atomic Force Microscope (AFM), and Superconducting Quantum Interference Device (SQUID) magnetometry, the samples are patterned with a standard optical lithography process into 10x10 µm 2 square junctions for magneto-transport measurements. Following top electrode fabrication, the bottom electrode is etched by Ar ion milling. To control the etching, a secondary ion mass spectrometer (SIMS) is used and allows stopping the etching especially before the STO barrier during the first step. An additional STO barrier etching is necessary to access to the bottom electrode contact. The contact pads are defined by a 100nm thick palladium layer deposited by sputtering. After microfabrication, two junctions of 10 x 10 µm 2 are measured on each sample to verify the reproducibility and uniformity of the sample. The resistance area product (RA) was in the range of 0.1-10MΩµm 2 for all junctions.

PRELIMINARY CHARACTERIZATIONS

Preliminary characterizations are preformed to verify the magnetic and structural properties. The hysteresis cycles, measured at low temperature on the stacks described in Table 1 and presented in Fig. 2

TUNNELING MAGNETORESISTANCE

The transport measurements are performed after microfabrication in four point configuration (Fig. 1). The junctions are biased using a commercial current source and the voltage is sent to a set of homemade low noise amplifiers, the output being sent to a voltage acquisition card. The studied samples are cooled down in a He4 cryostat that contains also a superconducting coil in order to record the magnetoresistance curves and the current voltage curves at different magnetic fields.

- [START_REF] Brinkman | Tunneling Conductance of Asymmetrical Barriers[END_REF]. The fitting parameters are the height and the width of the junctions. (c) Dynamical conductance Gd=dI/dV measured at 5K in the parallel (black squares) and antiparallel (red circles) states; the parallel and antiparallel states are different because a TMR effect is present at 5K contrary than at 300K. First, the TMR ratio of the junction is measured, because it is a primordial parameter in MTJ. As example, Fig. 3 shows typical magnetoresistance measurements obtained at low temperature. We define the TMR ratio by

P P AP R R R TMR / ) ( 100 (%)  
, where AP R and P R are the antiparallel and parallel resistance respectively. In order to extract information on the quality of the barrier we performed other characterizations, especially I(V) characteristics at 300K and at 5K. At 300K (Fig. 3b), a simulation using a Brinkman [START_REF] Brinkman | Tunneling Conductance of Asymmetrical Barriers[END_REF] fit is realized, to extract the height and the width of the barrier. The estimated barrier height varies from 0.2 to 0.8eV between the different junctions (Table 1), which is much smaller than the value of the band gap in STO [START_REF] Van Benthem | Bulk electronic structure of SrTiO 3 : Experiment and theory[END_REF] but also slightly higher than values published for this type of junctions [START_REF] Viret | Low-field colossal magnetoresistance in manganite tunnel spin valves[END_REF]. The width is also smaller than the real STO layer thickness. These two parameters are strongly dependent on the quality of the barrier and a possible explanation of their observed reduction is the decrease of the STO quality inside the barrier compared to bulk material. Impurities could be present in the barrier, such as oxygen vacancies. The conductance at 5K (Fig. 3c) is no more parabolic as a function of the voltage and a strong drop of the conductance also called "zero bias anomaly" (ZBA) is observed at zero bias and is estimated at 5K by

mV mV mV G G G ZBA 0 0 100 / ) (  
. This ZBA has been attributed to the presence of metallic impurities in the barrier or at the interface with the electrodes which induce additional conduction channels [START_REF] O'donnell | Colossal magnetoresistance magnetic tunnel junctions grown by molecular-beam epitaxy[END_REF] or a Coulomb gap [START_REF] Sun | Temperature and bias dependence of magnetoresistance in doped manganite thin film trilayer junctions[END_REF]. A magnetic and transport dead layer in the LSMO at the interface with the STO creating an extra barrier may also be responsible for the ZBA. The effect of this double barrier disappears when the voltage increases. Thus a reduced ZBA will be characteristic of a better quality of the interface or of the barrier.

All the measurement curves are not shown in this paper but the key results are summarized in Fig. 4 for all the junctions (Table 1). First, the ZBA has been measured only on one junction for the PLD1350 and PLD1351 and explain the missing values in Fig. 4. The comparison of the samples PLD1353 and PLD1355 shows that replacing the crystalline STO target by a sintered one induces no visible change on the junction properties. The optimal growth conditions of the STO in the chamber are those used during the deposition of the PLD1348. The full stack is annealed during one hour under oxygen just after the growth in order to recover the oxygen stoechiometry. The quality of the barrier is improved with barrier heights of the PLD1348 junctions higher than for the sample PLD1350 grown with the standard conditions. The conductance anomaly around zero bias (ZBA) is also reduced, so that this anomaly seems indeed linked to the barrier quality. Nevertheless, the measured TMR is small, hinting at a limitation due to the LSMO electrodes and the interfaces with the STO more than to the barrier quality. The junctions with a LAO barrier present a very large resistance which makes them difficult to measure. Indeed, this large resistance induces an important capacitive effect and a very large response time of the junction. The I(V) curves could not be measured at 5K (no ZBA in Fig. 4). The LAO barrier is thus probably thicker than 3nm. Junctions with a smaller LAO barrier have to be grown in order to be able to compare the results with the reference junction.

The junctions grown on the STO(110) substrate (PLD1354) exhibit an important roughness as well as enhanced coercive fields as presented in the previous part. A zero TMR ratio has been measured and the RA product is weak. A leakage through the barrier induced by the roughness could explain the decrease of resistance and the loss of TMR. The growth conditions of the junction on top of STO(110) have to be optimized to avoid these pinholes. The growth temperature was found to have little impact on the junction characteristics. Indeed, the measurements performed on the junctions PLD1350 and PLD1353, for which only the growth temperature changes from 585° to 550°, are very similar with a slightly increased TMR and barrier height upon lowering the growth temperature. This is at variance with an increase of the O 2 pressure during the LSMO growth which degrades all the junction properties, with very small TMR and barrier height on the junctions PLD1351. The weak RA product could be the signature of pinholes in the barrier. Another possible explanation is a columnar rather than layer-by-layer growth mode of the LSMO grown at high O 2 pressure. The STO of the barrier would then be comprised of small grains following the LSMO columnar structure and leakage due to oxygen vacancies could appear at the grain boundaries. Nevertheless we expect in this case a change in the LSMO coercive field, which is not what is observed (Fig. 2 (a)). Thus this study seems to show that the interface between the barrier and the electrode, as the roughness or the LSMO properties is the key parameter for the junction properties and especially the TMR ratios. The quality of the barrier as soon as the barrier is insulating and without pinholes, has less impact on the TMR ratios.

CONCLUSION

We have studied LSMO based tunnel junctions with various growth conditions, barrier structure and substrate. By playing with materials and growth conditions, the importance of the quality of the electrode on the junction properties has been highlight. The quality of the barrier as soon as the barrier is insulating and without pinholes, has less impact on the TMR ratios. Moreover, we observed big impacts on the properties due to little modifications of the growth conditions and materials. Besides, this fact could explain the variety of TMR results observed in literature due to a variety of growth conditions and techniques. A more complete study of the best junctions could be performed through transport noise study, to go a step further toward uses and medical applications.

Figure 1 .

 1 Figure 1. Picture of a processed 10x10µm² junction. The light green part is the bottom electrode, the red crosshatched square is the top electrode and the light blue shapes are the Pd contacts.

Figure 2 .

 2 Figure 2. (a) Hysteresis cycles measured by SQUID magnetometry at 10K on the as grown stack described in Table 1. Two magnetization switches are observed and correspond to the coercive fields of the two electrodes of the junctions. (b) Typical AFM image 2x2µm² observed on top of the junctions after growth. The steps coming from the vicinal substrate are reproduced on top of the junction, which is characteristic of a layer-by-layer smooth growth mode. The rms (root mean square) roughness is lower than 0.2nm.

Figure 3 .

 3 Figure 3. Measurements performed on a 10x10 µm 2 junction of the PLD 1355 sample. (a) TMR versus magnetic field measured at 5K with 0.1nA bias current. (b) Dynamical conductance Gd=dI/dV measured at 300K. The red line is the simulation based on the Brinkman model[START_REF] Brinkman | Tunneling Conductance of Asymmetrical Barriers[END_REF]. The fitting parameters are the height and the width of the junctions. (c) Dynamical conductance Gd=dI/dV measured at 5K in the parallel (black squares) and antiparallel (red circles) states; the parallel and antiparallel states are different because a TMR effect is present at 5K contrary than at 300K.

(

Figure 4 ..

 4 Figure 4. Summary of the measurements performed on the junctions: product resistance x area (RA) at 300K, TMR at 5K and at zero bias, zero bias conductance anomaly (ZBA) at 5K in the parallel state mV mV mV G G G ZBA 0 0 100 / ) (  . The heights of the tunnel barrier are the ones obtained thanks to the Brinkman fit realized at 300K on the I(V) characteristics.

Table 1 :

 1 Summary of the growth conditions for all the samples: name, stack, growth temperature, O 2 pressure during and after the growth. The parameters in red and italic are modified compared to the standard growth conditions used for the PLD1350 highlighted

	in bold. The thicknesses of the junctions are identical: LSMRO (20nm) /LSMO (10nm) /Barrier (3nm) /LSMO (30nm) // substrate.
	The type of target used for STO is indicated into brackets: monocrystalline "STO(mono)" or sintered "STO(sintered)".
	N°	Junction composition	Growth temperature	O 2 pressure during the growth	O 2 pressure after the growth
	PLD1350	LSMRO/

LSMO/STO(mono)/LSMO STO(100) substrate 585°C 120 mTorr 300 Torr

	PLD1348	LSMRO/LSMO/STO(mono)/LSMO STO(100) substrate	585°C	1.5.10 -2 mTorr (STO) 120 mTorr (other)	1h at 500° under 300 Torr
	PLD1349	LSMRO/LSMO/LAO/LSMO STO(100) substrate	585°C	2.2.10 -4 mTorr (STO) 120 mTorr (other)	1h at 500° under 300 Torr
	PLD1351	LSMRO/LSMO/STO(mono)/LSMO STO(100) substrate	585°C	250mTorr (LSMO) 120 mTorr (other)	300 Torr
	PLD1353	LSMRO/LSMO/STO(mono)/LSMO STO(100) substrate	550°C	120 mTorr	300 Torr
	PLD1354	LSMRO/LSMO/STO(mono)/LSMO STO(110) substrate	585°C	120 mTorr	300 Torr
	PLD1355	LSMRO/LSMO/STO(sintered)/LSMO STO(100) substrate	550°C	120 mTorr	300 Torr


S Datta

M Rioult

D Stanescu

H Magnan

A Barbier

Manipulating the ferroelectric polarization state of BaTiO 3 thin films

Keywords:

Controlling the ferroelectric polarization at macroscopic or microscopic levels is crucial in the framework of the development of ferroelectric materials used in yet challenging photo-electrochemical (PEC) cells and spintronic applications. We report here on polarization methods allowing to electrically polarize prototypical samples of BaTiO 3 (001) films. Epitaxial single crystalline layers were grown up to a thickness of 25 nm by atomic oxygen assisted molecular beam epitaxy on 1 at.% Nb doped SrTiO 3 (001) single crystals. The samples were both microscopically and macroscopically polarized using Piezoresponse Force Microscopy and electrochemical poling in an electrolyte respectively. In addition we demonstrate the possibility to retrieve a quasi-native mixed ferroelectric polarization state after annealing. These polarization methods may be applied to many other ferroelectric thin films.

Introduction

There is currently a strong interest towards developing ferroelectric materials for various applications such as their use in multiferroics systems [START_REF] Tokunaga | Composite domain walls in a multiferroic perovskite ferrite[END_REF][START_REF] Evgeny | Multifunctional Oxide Heterostructures[END_REF][START_REF] Jedrecy | Strong magnetoelectric coupling in multiferroic Co/BaTiO3 thin films[END_REF][START_REF] Barbier | Antiferromagnetic long-range spin ordering in Fe-and NiFe2-doped BaTiO3 multiferroic layers[END_REF], photoelectrochemical cells (PEC) [START_REF] Rioult | Tailoring the photocurrent in BaTiO3/Nb:SrTiO3 photoanodes by controlled ferroelectric polarization[END_REF][START_REF] Ji | Epitaxial ferroelectric BiFeO3 thin films for unassisted photocatalytic water splitting[END_REF][START_REF] Meng | BaTiO3 photoelectrodes for CdS quantum dot sensitized solar cells[END_REF][START_REF] Park | A ferroelectric photocatalyst for enhancing hydrogen evolution: polarized particulate suspension[END_REF], and heat and gas sensors [START_REF] Dong | Gas-sensing and electrical properties of perovskite structure p-type barium-substituted bismuth ferrite[END_REF][START_REF] Navrotsky | Thermochemistry of perovskite-related oxides with high oxidation states: superconductors, sensors, fuel cell materials[END_REF][START_REF] Tasaki | Fabrication of Sm-based perovskite-type oxide thinfilms and gas sensing properties to acetylene[END_REF]. Ferroelectricity is a particular property of non-centrosymmetric materials exhibiting a spontaneous electrical polarization that can be reversed by the application of an external electric field. In the recent years, an impressive amount of contributions have focused on the ferroelectric properties of perovskite-based oxides such as BiFeO 3 , BaTiO 3 , PZT (Lead Zirconium Titanate) [START_REF] Yun | Giant ferroelectric polarization beyond 150 μC/cm 2 in BiFeO3 thin film[END_REF][START_REF] Yang | Metalorganic chemical vapor deposition of lead-free ferroelectric BiFeO3 films for memory applications[END_REF][START_REF] Iijima | Preparation of ferroelectric BaTiO3 thin films by activated reactive evaporation[END_REF][START_REF] Young | Critical thickness of ultrathin ferroelectric BaTiO3 films[END_REF][START_REF] Fukushima | Preparation of ferroelectric PZT films by thermal decomposition of organometallic compounds[END_REF][START_REF] Moazzami | Electrical characteristics of ferroelectric PZT thin films for DRAM applications[END_REF]. For such crystals, in the ferroelectric state, the barycenter of cations and anions in the unit cell are slightly spatially separated, inducing an electric dipole, i.e. an electric polarization. Dipoles rearrange themselves to minimize the sample energy forming different ferroelectric domains wherein the electric polarization points towards a given direction. Usually in an as-grown thin ferroelectric film the polarization is not uniform and the sample exhibits several ferroelectric domains where the polarization vector is pointing towards different directions. However, it is possible to control the orientation of the domains through the application of a saturating external electric field that allows the majority of the electrical dipoles within a domain to be oriented in a particular direction. Because of the wealth of technologically important applications deriving from the inclusion of ferroelectric layers in smart devices, it is of major importance to understand and master the manipulation of the domains electric polarization orientation in thin films, at the microscopic and macroscopic scales. Recently multiferroic BiFeO 3 samples were considered within this framework. M. Singh et al. have reported on the local polarization of BiFeO 3 samples at the microscopic scale by applying a 10 V potential on an AFM tip [START_REF] Singh | Metallorganic chemical-vapor-deposited BiFeO3 films for tunable high-frequency devices[END_REF]. Cao et al. have independently demonstrated the poling of BiFeO 3 samples at the macroscopic scale in a solution of Lithium Chlorate (LiClO 4 ) in propylene carbonate [START_REF] Cao | Switchable charge-transfer in the photoelectrochemical energy-conversion process of ferroelectric BiFeO3 photoelectrodes[END_REF]. In the present work we consider the prototypical ferroelectric material BaTiO 3 that has the major advantages as exhibiting a simple perovskite crystalline structure and featuring a para-to ferro-electric transition at relatively high temperature (130 °C for bulk BaTiO 3 and even much higher for strained layers [START_REF] Choi | Enhancement of ferroelectricity in strained BaTiO3 thin films[END_REF]). Here we demonstrate routes to master the manipulation of the polarization state at microscopic and macroscopic scales of BaTiO 3 /Nb:SrTiO 3 samples, as well as the possibility to erase an applied polarization state by an air annealing of the very same samples.

Experimental details

The BaTiO 3 epitaxial layers were deposited on single crystalline 1 at.% Nb:SrTiO 3 (001) (i.e. ~0.5 wt% Nb) substrates using atomic oxygen assisted molecular beam epitaxy (AO-MBE), a technique that makes possible the deposition of single crystalline layers of controlled morphology, stoichiometry and thickness [START_REF] Barbier | Surface composition of BaTiO3/SrTiO3(001) films grown by atomic oxygen plasma assisted molecular beam epitaxy[END_REF]. The doped substrate is conductive enough to use it as electrode during the BaTiO 3 overlayer poling procedures. The oxide layers were formed by thermal evaporating high purity metals from Knudsen cells, while exposing the sample to an atomic oxygen plasma (RF power = 350 W) in an ultrahigh vacuum vessel (10 -7 mbar working conditions, 10 -10 mbar base pressure). In order to have a homogeneous deposition and to reach the expected stoichiometry, the samples are rotated continuously around their surface normal during the deposition while kept at a temperature of ~900 K. The oxide deposition rate was about 0.12 nm/min. To investigate the quality and the structure of the grown epitaxial films, in situ reflection high-energy electron diffraction (RHEED) patterns were acquired during the film growth. The thickness of the BaTiO 3 layers investigated in the present work lies in-between 15 and 30 nm. In this range no influence of thickness on any measurements could be observed within the techniques used in this paper which is noticeably consistent with the previous report by Kim et al. [START_REF] Kim | Critical thickness of ultrathin ferroelectric BaTiO3 films[END_REF] who reported a critical thickness for ferroelectricity of BaTiO 3 of 4 nm and very little changes of the remnant polarization with thickness above 15 nm.

In situ XPS spectra of the Ba 3d, Ti 2p and O 1s core levels, using Al-K α radiation, were systematically recorded just after deposition to check the consistency of the BaTiO 3 layer stoichiometry with previous reports [START_REF] Barbier | Surface composition of BaTiO3/SrTiO3(001) films grown by atomic oxygen plasma assisted molecular beam epitaxy[END_REF]. XPS spectra were also recorded after the different polarization treatments to check eventual compositional changes in the layers.

We used Atomic Force Microscopy (AFM) and Piezoresponse Force Microscopy (PFM) to investigate the surface morphology and ferroelectric properties of our samples respectively. The images were realized with a Brüker™ scanning probe microscope and a Nanoscope V controller.

The samples were electrically poled using two major strategies as depicted schematically in Fig. 1. (i) Local poling (microscopic scale) was achieved using PFM, which was used in writing mode to create well defined ferroelectric domains in our samples, by poling with various positive or negative potentials applied on the tip (in the [0 V, 10 V] range), Fig. 1a andb. Since, the back of the sample was always connected to the ground, the sample surface was thus exposed to potentials in the [-10 V, +10 V] range. This poling procedure will be hereafter referred to as PFM-poling. The tip velocity was 0.4 μm/s and the writing resolution was 20 nm. (ii) Macroscopic poling was realized using a non-aqueous LiClO 4 electrolyte (0.1 M LiClO 4 in propylene carbonate), with a large electrochemical window (i.e. a large voltage range in between which the substance is neither oxidized nor reduced), in an electrochemical cell with 2 electrodes. The sample was mounted as the working electrode and a Pt wire as the counter electrode. The geometry of the experimental setup is shown on Fig. 1c andd and a scheme of the experiment is given on Fig. 1e. This poling procedure will be hereafter referred to as EC-poling. During the EC-poling procedure, the sample was kept at 0 V. Potentials of + 8 V and -8 V were applied several times alternatively (while monitoring the current) on the platinum counter electrode (during 10 s each) in order to reinforce the ferroelectric state as proposed by Roelofs et al. [23]. To obtain a remnant electric polarization pointing towards the substrate (by convention called downward polarization), the last applied potential was +8 V. Indeed, in a previous work [START_REF] Rioult | Tailoring the photocurrent in BaTiO3/Nb:SrTiO3 photoanodes by controlled ferroelectric polarization[END_REF], we demonstrated that +8 V EC-poling leads to downward polarization of the BaTiO 3 layer. Lastly the electric current monitoring is of crucial importance during EC-poling since overcurrent appeared to be strongly detrimental to the samples as we will demonstrate later in this paper.

The PFM-poled written ferroelectric domains were observed by PFM and Low Energy Electron Microscopy (LEEM). Beyond the recognized full field morphology characterization capabilities of LEEM [START_REF] Bauer | Low energy electron microscopy[END_REF][START_REF] Bauer | Low-Energy Electron Microscopy[END_REF], it has been previously reported that surface charge changes, expected when the ferroelectric polarization orientation changes, provide a visible LEEM contrast [START_REF] Rault | Thickness-dependent polarization of strained BiFeO3 films with constant tetragonality[END_REF]. This phenomenon was used to observe ferroelectric contrasts between polarized and unpolarized regions of the sample regions. The samples were optically patterned by laser lithography and e-beam evaporation to include metallic Au landmarks on the sample surface that allow finding the PFM polarized regions easily. The 200 nm thick Au landmarks were evaporated on top of a 15 nm thick Ti seed layer. For a broader view, one may also use Kelvin probe force microscopy [START_REF] Kim | Local surface potential distribution in oriented ferroelectric thin films[END_REF][START_REF] Wang | A new non-destructive readout by using photo-recovered surface potential contrast[END_REF] or scanning surface potential microscopy [START_REF] Chen | Surface potential of ferroelectric thin films investigated by scanning probe microscopy[END_REF] to map the surface charge changes due to different ferroelectric domains polarization orientations, and DARTSSPFM for P-E loops measurements [START_REF] Rodriguez | Dual-frequency resonancetracking atomic force microscopy[END_REF].

The high chemical stability in air of BaTiO 3 , as well as of the Nb:SrTiO 3 substrates, allows for air annealing at high temperature without composition changes. Since annealing above the ferroelectric orderdisorder temperature is easily possible in air, it appears as a method of choice to "reset" the ferroelectric domains orientation after PFM-or EC-poling. This was realized in a high temperature furnace above the Curie temperature of BaTiO 3 with typical rates of 14 °C/min up to 280 °C.

Results and discussion

Growth

In situ RHEED patterns were observed and acquired during film growth to monitor the quality and the structure of the grown epitaxial samples. Upon growth, the RHEED patterns switched from features typical of the perovskite Nb:SrTiO 3 substrate to the ones characteristic of the BaTiO 3 thin epitaxial film (Fig. 2), similarly to observations reported using undoped SrTiO 3 (001) substrates [START_REF] Barbier | Surface composition of BaTiO3/SrTiO3(001) films grown by atomic oxygen plasma assisted molecular beam epitaxy[END_REF]. By convention we indexed the reciprocal space using the reciprocal space unit vectors (10) ⁎ and (01) ⁎ to describe the primitive reciprocal surface lattice as depicted in Fig. 2e. The RHEED reciprocal space streaks remain mostly 2D rod-like with some hints of island formation for high thicknesses (intensity modulations along the streaks at high thickness). The surface lattice parameter can easily be extracted from the streaks position while the error bar is derived from the parameter dispersion obtained using several azimuths. The surface lattice parameter evolution (Fig. 2f) reveals an almost complete structural relaxation for thicknesses above 10 nm. From our RHEED observations we can conclude that the layers are mainly epitaxial, single crystalline and flat. The chemical composition of the as-grown layers was investigated by in situ XPS. A wide scan recorded over the entire accessible binding energy region (Fig. 3) shows the presence of the expected species (Ti, O and Ba) and no additional eventual surface contamination. As previously reported in reference [START_REF] Barbier | Surface composition of BaTiO3/SrTiO3(001) films grown by atomic oxygen plasma assisted molecular beam epitaxy[END_REF] a Ba rich surface oxide is necessary to promote a good epitaxial growth leading to Ba/Ti ratio above unity (typically ~1.3). The binding energy of the Ba 3d 5/2 line was found at 780.4 eV from detailed measurements of the Ba 3d XPS lines as reported in Fig. 4.

Microscopic polarization

The ferroelectric nature of the layers was investigated by PFM. Scanning the tip-surface potential in the [-10 V, + 10 V] range evidences two distinct polarization states. In order to determine the ferroelectric domain repartition we locally polarized the sample by applying + 8 V and -8 V on the tip (substrate at ground) accordingly to the pattern shown in Fig. 5a. After the PFM-poling the pattern was read again by PFM using a sinusoidal potential applied on the sample (tip at ground) with a drive amplitude of 4000 mV. The resulting PFM phase image is shown in Fig. 5b. The integrated phase profile through the image of the written pattern (Fig. 5c) shows 3 different phase levels (58°, -60°a nd 54°) corresponding to the different remnant states: (i) + 8 V PFM-poled, (ii) -8 V PFM-poled and (iii) outside the writing region (i.e. 0 V or 'as-grown'), respectively. One may note that in this case the PFM phase of the as-grown region (54°) is close to the one after +8 V polarization (58°). This result shows that the as-grown state behaves more like the +8 V polarization states rather than like the -8 V ones and that the ferroelectric domain mix between up and down orientations is different in each region. Moreover we observe that the PFM phase difference between the +8 V and -8 V remnant states is about 118°, which is smaller than the expected 180°value. This phenomenon has been examined in more details in previous reports and can have different origins: (i) the domains are not fully polarized (i.e. a mix of upward and downward domains); (ii) the PFM signal is disturbed by a background signal [START_REF] Jungk | Quantitative analysis of ferroelectric domain imaging with piezoresponse force microscopy[END_REF] inherent to the PFM setup; (iii) an asymmetry in polarization orientation may arise from an internal built-in electric field at the bottom interface which may result in incomplete 180°s witching [START_REF] Gruverman | Asymmetric nanoscale switching in ferroelectric thin films by scanning force microscopy[END_REF]; and/or the drive amplitude was large enough to induce a charge build-up. Although the later effect is unlikely since successive scans did not evidenced any evolution of the patterns. However, one may note that an improved contrast using lower drive amplitude may be obtained thanks to the DARTSSPFM method [START_REF] Rodriguez | Dual-frequency resonancetracking atomic force microscopy[END_REF], unfortunately this mode was not available in our setup. From region to region or sample to sample (within a set of about 10 samples) the observed domain mix, in the as-grown region, spans a wide range from nearly equiprobable to almost single domain. We can thus conclude that for the asgrown samples, the ferroelectric domains are arranged in a random fashion, i.e. a mix of upward and downward states.

PFM written patterns were subsequently investigated using full field LEEM imaging. Contrarily to PFM that can be used for writing and reading patterns, the LEEM technique allows a rapid investigation of large areas of the sample without eventual unwanted modifications as could be induced by a scanning tip, like during PFM measurements. Prior to LEEM imaging the sample was outgassed for several hours at 80 °C. Working close to the Mirror Electron Microscopy (MEM) to LEEM transition (starting voltage close to 1 V) allows the best imaging conditions and contrast. Thanks to the Au landmarks, the written pattern can easily be located as it can be seen in Fig. 6b (50 μm field of view (FOV). A closer inspection with a FOV of 10 μm (Fig. 6c) confirms the presence of 3 different contrasts and thus the mixed initial polarization states of the as-grown samples. The LEEM observations are fully consistent with the previous PFM analysis.

Macroscopic polarization

After a PFM-poling (using the pattern shown on Fig. 6a), the sample was macroscopically EC-poled by keeping the current below 1 mA and observed again by LEEM. As it can be seen in Fig. 6d the pattern completely vanished upon this +8 V EC-poling procedure, indicating that the macroscopic polarization overcomes the PFM writing. To confirm that the ferroelectric properties of the sample were unaltered by the EC-poling we performed again a PFM-poling using the initial PFM pattern in identical conditions and observed it subsequently by LEEM. As it can be seen in Fig. 6e the pattern is clearly visible defining again 3 contrasts states (+ 8 V remnant state, -8 V remnant state and ECpoled remnant states). In addition one may note that the surface morphology does not evolve upon EC-poling, remaining smooth and flat. Moreover, as revealed by XPS the composition remained unchanged after EC-poling since the peak intensities only marginally varied (Fig. 4). The peak position of the Ba 3d 5/2 line (Fig. 4b) is shifted to 780.2 eV. This shift is likely linked to a global surface charge effect although it is only slightly above the experimental error bar. Therefore we can conclude that EC-poling a BaTiO 3 epitaxial layer in LiClO 4 while keeping the current below 1 mA during the poling procedure changes neither the surface composition nor the morphology of the sample.

Overcurrent

By studying EC-poling with different successive increasing currents flowing through the sample, we found out that a current flow above 1 mA leads to detrimental results. We investigated in more details the effect of not limiting the current flowing through the sample during the EC-poling procedure. AFM images of a sample having experienced EC-poling without limiting the current (overcurrent) is reported in Fig. 7. During the EC-poling procedure, the sample was mounted in a special sample holder and fixed with a sealing joint in order to avoid the electrolyte to be in contact to the back of the substrate. As a matter of fact, the macroscopic polarization of the samples required some particular attention. Since BaTiO 3 is ferroelectric and in consequence necessarily also piezoelectric, a mechanical strain may compete or even overcome the effect of an external electric field when one desires to tune the internal ferroelectric polarization. It is thus mandatory to get rid of any mechanical strain while performing an external poling procedure. In the case of our EC-poling procedure this was achieved by etching out a ring down to the substrate on the sample that corresponds to the position of the electrochemical cell sealing joint (Fig. 7). The etched ring for sealing purpose separates the sample in two regions: outside the ring the 30 nm thick BaTiO 3 layer was not exposed to the electrolyte (Fig. 7a) while the inside region which was EC-poled in LiClO 4 was the siege of eventual overcurrent situations (Fig. 7b). This ring patterning allows comparing, on the same BaTiO 3 sample, surfaces which underwent overcurrent upon EC-poling (Fig. 7b, inside) or not (Fig. 7a, outside). Integrated topography profiles at locations corresponding to different contrasts (Fig. 7d) show that the difference of height between the B and C regions is around 17 nm. Since the nominal thickness of this BaTiO 3 film was 30 nm, we can conclude that the darker regions correspond to sample zones where the BaTiO 3 film partially disappeared. The film morphology and thickness seem unperturbed in the lighter (unexposed) regions. The corresponding global XPS scan (Fig. 3) recorded on this sample confirms the massive loss of BaTiO 3 after the overcurrent exposure of the surface, which is most likely due to the dissolution in the electrolyte.

Annealing

Macroscopic single polarization states as likely produced by ECpoling may be interesting for specific applications. For example we could show a 2.5 times increased photocurrent density during solar water splitting using a polarized BaTiO 3 layer with respect to a nonpolarized configuration [START_REF] Rioult | Tailoring the photocurrent in BaTiO3/Nb:SrTiO3 photoanodes by controlled ferroelectric polarization[END_REF]. However, for many aspects, it is very desirable to be able to reset the sample domain configuration. To ensure the complete control of the ferroelectric polarization and the repeatability of our different poling procedures, we performed an air annealing of an EC-poled samples at 280 °C (above the Curie temperature of BaTiO 3 ). After annealing we could write again patterns using PFM indicating that after crossing the Curie temperature the sample came back to a more mixed polarization state (effective ferroelectric polarization "reset"). XPS showed unchanged composition and the Ba 3d 5/2 lines where found at a position intermediate between the as-grown situation and the EC-poled ones (Fig. 4b). The ability to reset the polarization state was investigated in more details by PFM starting from a virgin sample. The ferroelectric domain configuration of a 15 nm thick BaTiO 3 /Nb:SrTiO 3 layer was investigated before and after annealing in air at 280 °C for 20 min (Fig. 8). The sample possessed Au landmarks, realized by lithography, to allow a precise positioning of the tip on a given region that was observed prior (Fig. 8a) and after writing (Fig. 8b). The very same region was imaged again after a sample annealing realized at 280 °C for 20 min. As can be seen on Fig. 8c the pattern was "erased" after the annealing procedure. Finally, we checked that the sample showed similar ferroelectric behavior after annealing by writing a pattern on another region of the same sample, as observed in Fig. 8d. From these observations we can conclude that after an annealing cycle the ferroelectric configuration of the layer is "reset" to a new mixed polarization state configuration.

Conclusion

In conclusion, we have demonstrated efficient methods to manipulate the polarization state of a prototypical BaTiO 3 ferroelectric layer and also to retrieve a new mixed polarization state after macroscopic polarization. We showed that our layers can be locally polarized using PFM within the [-10 V; +10 V] potential range. We have also shown that the BaTiO 3 layer can be downward polarized macroscopically in a solution of LiClO 4 in propylene carbonate chloride, provided that the current is maintained below 1 mA. Upon overcurrent the BaTiO 3 layer is likely to dissolve. Moreover, we found that an air annealing at a temperature above the Curie temperature performs a sort of "reset" of the ferroelectric polarization. Our study presented a set of tools in order to master the polarization states of ferroelectric layers, which is of top interest in the framework of the use of ferroelectric materials in new devices having tremendous potential in many technologically important fields like multiferroics for spintronics or photoelectrochemical water splitting devices in particular.

Fig. 1 .

 1 Fig. 1. (a) to (d) Geometry and expected polarization states using PFM-poling with -8 V (a) and +8 V (b) poling potentials and using EC-poling with -8 V (c) and +8 V (d) poling potentials. (e) Scheme of the EC-poling experiment.

Fig. 2 .

 2 Fig. 2. (Left panel) RHEED patterns of a 28 nm BaTiO 3 film grown by AO-MBE on Nb:SrTiO 3 at different deposition times/thicknesses (steps (a) to (d)) over the two lowest Miller index surface diffraction directions. (e) The corresponding surface reciprocal lattice is represented; making explicit the diffraction directions, the elementary cell (in the reciprocal space) is also shown. (f) In-plane lattice parameter derived from the RHEED data with the same method as in reference [21]. Dashed lines stand for the in-plane lattice parameter of bulk BaTiO 3 (bluetop) and Nb:SrTiO 3 (red-bottom) at the deposition temperature of 900 K (taking into account the thermal expansion).

Fig. 3 .

 3 Fig. 3. Wide XPS scan of a 15 nm thick BaTiO 3 /Nb:SrTiO 3 (001) sample elaborated by AO-MBE. ((blue, top curve) As grown sample and (red, bottom curve) for the same sample having experienced overcurrent during the macroscopic EC-poling procedure. The O 1s line (located at ca. 530.1 eV) of the as-grown sample was used as a reference position and the curves were shifted vertically for the sake of clarity.

Fig. 4 .

 4 Fig. 4. Ba 3d XPS photoemission lines (a) recorded for a 20 nm BaTiO 3 /Nb:SrTiO 3 (001) layer grown by AO-MBE and (b) enlarged region of the main Ba 3d 5/2 line: (-, red solid line) as-grown layer, (-, black dashed line) after +8 V EC-poling in a LiClO 4 bath and (+, blue cross-plus) after +8 V EC-poling and a subsequent annealing in air at 280 °C.

Fig. 5 .

 5 Fig. 5. PFM measurements performed on a 15 nm thick BaTiO 3 /Nb:SrTiO 3 (001) sample realized by AO-MBE. (a) Sketch of the written pattern used to locally polarize the sample; (b) observed PFM phase pattern in reading mode after writing at ±8 V the pattern as described in (a) and using a drive amplitude of 4000 mV; (c) integrated intensity along the blue dashed box depicted in (b).

Fig. 6 .

 6 Fig. 6. LEEM images of a 15 nm thick BaTiO 3 /Nb:SrTiO 3 (001) sample elaborated by AO-MBE. (a) PFM pattern used to write ferroelectric domains using ±10 V electric potentials; (b) and (c) as grown situation with written PFM pattern for two field of views (FOV) (50 μm and 10 μm respectively) -The rectangular features are Au landmarks, as indicatedthe arrow in (b) highlights the ease of observation of a the written pattern with respect to the landmarks. (d) Observation of the same region (50 μm FOV) than (a) after macroscopic EC-poling of the sample with a potential of +8 V in LiClO 4 , the arrow indicates the original position of the PFM pattern; (e) observation of the same region than (d) (10 μm FOV) after writing again the PFM pattern depicted in (a) after EC-poling (10 μm FOV). All images were taken around the MEM-LEEM transition potential with starting voltages ranging from 0.8 to 1.2 V.

Fig. 7 .

 7 Fig. 7. Fast Scan AFM images recorded on a 30 nm thick BaTiO 3 /Nb:SrTiO 3 (001) sample elaborated by AO-MBE on both sides of an etched ring after the sample having been exposed to an overcurrent during the macroscopic EC-polling polarization in LiClO 4. (a) Region non-exposed to the electrolyte (no overcurrent, unpolarized as-grown, left region) and a fraction of the ring (right region); (b) fraction of the ring (left region) and region exposed to the electrolyte (EC-poled, overcurrent, right region); (c) schematic drawing of the sample with the etched ring and (d) integrated step heights as measured on zones labeled A, B and C in (a) and (b).

S. Datta et al. / Thin Solid Films 607 (2016) 7-13

Acknowledgements

This work was funded by the CEA (DSM-Energie program under the PAPI grant) and supported in part by Triangle de la Physique and Ile-de-France (C'Nano and ISC-PIF) under the IMAMFP grants.

Conflict of interest

The authors declare no conflict of interest.